Nanotechnology: Nanofabrication, Patterning, and Self Assembly

NANOTECHNOLOGY SCIENCE AND TECHNOLOGY SERIES

NANOTECHNOLOGY: NANOFABRICATION, PATTERNING AND SELF ASSEMBLY

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NANOTECHNOLOGY SCIENCE AND TECHNOLOGY SERIES

NANOTECHNOLOGY: NANOFABRICATION, PATTERNING AND SELF ASSEMBLY

CHARLES J. DIXON AND

OLLIN W. CURTINES EDITORS

Nova Science Publishers, Inc. New York

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. Any parts of this book based on government reports are so indicated and copyright is claimed for those parts to the extent applicable to compilations of such works. 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 Nanotechnology : nanofabrication, patterning, and self assembly / editors, Charles J. Dixon and Ollin W. Curtines. p. cm. Includes index. ISBN 978-1-61761-771-3 (Ebook) 1. Nanostructured materials. 2. Nanotechnology. I. Dixon, Charles J. II. Curtines, Ollin W. TA418.9.N35N3574 2009 620'.5--dc22 2009004990

Published by Nova Science Publishers, Inc. Ô  New York

CONTENTS Preface

xi

Research and Review Studies

1

Chapter 1

Electrochemical Nanofabrication Di Wei

3

Chapter 2

Fabrication and Application of Novel Two-Dimensional Nanowebs via Electrospinning Bin Ding, Chunrong Li, Dong Wang and Seimei Shiratori

Chapter 3

Nano-scale Characterization and Spectroscopy of Strained Silicon Norihiko Hayazawa and Alvarado Tarun

Chapter 4

Nanotechnologies for Cancer Diagnostics and Treatment Phong Tran and Thomas J. Webster

Chapter 5

Mechanical Characterization at Nanometric Scale of Ceramic Superconductor Composites J.J. Roa, X.G. Capdevila and M. Segarra

151

Chapter 6

ZnO Nanowire Arrays: Template-free Assembly Growth and Their Physical Properties Bingqiang Cao and Weiping Cai

237

Chapter 7

Spatially Resolved Control of Electrical Resistivity in Organic Materials —Development of a New Fabrication Method of Junction Structures Toshio Naito

Chapter 8

Fabrication of Electrical Contacts on Individual Metal Oxide Nanowires and Novel Device Architectures Francisco Hernandez-Ramirez, Juan Daniel Prades, Roman Jimenez-Diaz, Olga Casals, Albert Cirera, Albert Romano-Rodriguez, Joan Ramon Morante, Sven Barth and Sanjay Mathur

51 71 107

275

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Contents

Chapter 9

Functionalization of Nanoparticles, Nanotubes and Nanowires by Surface-Initiated Atom Transfer Radical Polymerization Jinying Yuan, Mi Zhou and Yingwu Yin

309

Chapter 10

Synthesis and Applications of Nano-sized Ferroelectrics via Mechanochemical Activation L.B. Kong, Z. Xu and T.S. Zhang

331

Chapter 11

Preparation and Characterization of Monoatomic Carbon Chains: Unraveling, Field Ion Microscopy, and Field Emission Igor M. Mikhailovskij

371

Chapter 12

Sequential Nucleation and Growth of Complex Nanostructures by a Two-Step Strategy Li Yang, Paul W. May and Lei Yin

409

Chapter 13

Progress of Self-standing Diamond Film Fabricated by DC Arc Jet Plasma CVD G.C. Chen, F.X. Lu, B. Li, C.M. Li, W.Z. Tang, J.H. Song, L.F. Hei and Y.M. Tong

Chapter 14

Nanoshell Arrays: Fabrication and Enhanced Photoluminescence Zhipeng Huang and Jing Zhu

Chapter 15

A Strategy for the Incorporation of Trivalent Lanthanide Ions into Anatase Tio2 Nanocrystals Wenqin Luo, Chengyu Fu, Renfu Li and Xueyuan Chen

Chapter 16

Nanocrystallite Superhard Titanium Nitride Film in Multi-arc Ion Plating Xiang Yu, Chengbiao Wang, Meng Hua, Yang Liu and Shengli Ma

Chapter 17

Embedded Optical-electrical Nanomateriales Fabricated by Ion Implantation X.T. Zu, X. Xiang, S. Zhu and L.M. Wang

525

Chapter 18

Structural, Dynamical and Optical Properties of Self-assembled Porphyrins at the Mesoscopic Scale Valentina Villari, Norberto Micali and Luigi Monsú Scolaro

559

Short Communications Short Communication A The Influence of Thiophene Addition on Catalytic Pyrolysis of Poly (Dimethyl Siloxane) K.F. Cai, C.W. Zhou, A.X. Zhang and J.L. Yin

435

459

479

509

603

605

Contents

ix

Short Communication B Nanofinishing of Cotton Textiles N. Vigneshwaran and Virendra Prasad

615

Index

621

PREFACE This new book is dedicated to outstanding research in nanotechnology which is a “catchall” description of activities at the level of atoms and molecules that have applications in the real world. A nanometer is a billionth of a meter, about 1/80,000 of the diameter of a human hair, or 10 times the diameter of a hydrogen atom. Nanotechnology is now used in precision engineering, new materials development as well as in electronics; electromechanical systems as well as mainstream biomedical applications in areas such as gene therapy, drug delivery and novel drug discovery techniques. Nano- and micro-fabrications have been largely used in the applications such as integrated circuits, micro/nano electro-mechanical systems (M/NEMS), micro-optics and countless others. The methodology of nanofabrication can be divided into two types, topdown and bottom-up processes, which themselves can be further divided. Top-down process refers to approaching the nanoscale from the top (or larger dimensions), such as lithography, nanoimprinting, scanning probe and E-beam technique etc.. In bottom-up fabrication processes, the nanotechnology process builds nanoscale artifacts from the molecular level up, through single molecules or collections of molecules that agglomerate or self-assemble. Using a bottom-up approach, such as self-assembly enables scientists to create larger and more complex systems from elementary subcomponents (e.g. atoms and molecules). In general, top-down processes that transfer minute patterns onto material are more matured than bottomup processes. An exception is epitaxial processes that create layers through layer-by-layer growth with registry at the atomic level. Electrodeposition has actually been used for decades to form high quality, mostly metallic, thin films. It has recently been shown that high quality copper interconnects for ultra large scale integration chips can be formed electrochemically on Si wafer [1;2]. Electrodeposition has thus been shown compatible with state of the art semiconductor manufacturing technology. The largest semiconductor companies, for example, IBM, Intel, AMD, Motorola etc. are installing wafer-electroplating machines on their fabrication lines [1]. The electrodeposition of Cu with the line width 250 nm was used in the mass-production of micro-processor Pentium III in 1998. In 2003, the line width of the CPU was reduced to 130 nm in Pentium IV. Electrochemistry was largely used in chip fabrication [3] and the packaging of micro-electronics [4]. However, comparing with other nanofabrication techniques, electrochemical nanofabrication is still a maiden area which needs further development and fulfilment.

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Chapter 1 summarized the most recent developments in electrochemical nanofabrications. It includes not only the conventional technique, under potential deposition (UPD), which deals with the deposition of a single metal-ion on a definite substrate but also some new developments using ultrashort voltage pulsing and template methods for 3D construction of nano-materials. Electrochemical nanofabrication is a versatile method, which includes both top-down nanofabrication (e.g. electrochemical lithography) and bottom-up process such as electrochemical atomic layer epitaxy (EC-ALE). Nano-templates including anodized aluminum oxide (AAO) membranes, colloidal polystyrene (PS) latex spheres, single/aligned carbon nanotubes, selfassembled monolayers (SAMs), blocked copolymers and cyclodextrin molecules can be used for the preparation of various types of nanowires, nanotubes, ordered arrays of nanoparticles and nanodots electrochemically. Combining electrochemistry with other nanofabrication techniques such as focused ion beam (FIB) and self-assembly provides many novel strategies in the fabrication of nanomaterials with specific design. Selective areas in the nanoscale can be modified by electrochemical nanostructuring with metals, metal oxides and conducting polymers using a bipolar electrochemical technique. The traditional lithography and pattern technique is costly. In the construction of soft matters such as conducting polymers, traditional spin casting cannot guarantee nanostructures due to the fast speed of solvent evaporation. Electrochemical technique provides an innovative, versatile and economic way of nanofabrication. It especially offers better alternative to construct the soft matter nano-structures in a controllable manner. In general, electrochemical nanofabrication offers simplicity, efficiency, low-temperature processing, cost-effectiveness, the possibility in preparing large area deposits and precise control of the deposit thickness, which are the essential advantages than other nanofabrication techniques till date. Additionally, it can be used to prepare a wide range of materials comprising the inorganic and the organic. The former includes quantum dots, metallic and semiconducting (e.g. ZnO, TiO2) nanotubes and nanorods. The latter includes conducting polymer nanotubes and nanowires. Chapter 2 reviews our recent progress on the novel two-dimensional nanowebs by the optimization of processing parameters during electrospinning. Using high applied voltage and low relative humidity in chamber, the by-product of micro-sized defect films can be splitted into nanowebs due to the fast phase separation of the charged droplets which flight with high moving speed in electric field from capillary tip to collector. The electrospun fibers act as a support for the “fishnet-like” nanowebs comprising interlinked one-dimensional nanowires. The average diameter of the nanowires contained in typical nanowebs is about one order of magnitude smaller than that of conventional electrospun fibers. Nanowebs together with common electrospun nanofibers can be assembled into a three-dimensional fibrous mat. So far, nylon-6, polyacrylic acid (PAA), poly(vinyl alcohol) (PVA)/SiO2 nanoparticles, and PVA/zinc acetate have been found to have the possibility forming nanowebs. The formation, morphology, and area density of the nanowebs in electrospun fibrous mats are strongly affected by the applied voltage, ambient relative humidity, kinds of solvents, solution concentration and conductivity, and distance between capillary tip to collector. The expanded applications of electrospun fibers are expected due to the formation of nanowebs, such as the nano-sized controllable filters, high efficient catalysts, catalyst supporter, and sensors. The preliminary data showing that the sensitivity of PAA nanowebs to ammonia is 2.5 times higher than that of electrospun PAA nanofibers. Additionally, PAA nanowebs show much

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quicker absorption speed and larger capacities than that of PAA nanofibers during the ammonia absorption test. Strained silicon (ε-Si), the fundamental material of integrated circuit, is finding tremendous attention because it boosts the speed and reduces the power consumptions of electronic devices. However, poor homogeneity distribution of strain in ε-Si layers can degrade performance of electronic devices. Raman spectroscopy is used to study strain fluctuations in silicon because the optical phonons in Raman spectra are strongly influenced by strain. Though silicon are Raman active devices, the Raman efficiency of a nanometer layer of ε-Si is extremely weak and is often eclipsed under the Raman scattering of underlying buffer substrates. Micro Raman measurements show only uniform features in the nano-scale because of averaging effect from diffraction-limited spatial resolution. In Chapter 3, the authors utilized surface enhancement in Raman scattering to overcome weak emission problems and to suppress averaging effect. Thin ε-Si layers were covered with thin Ag layer to invoke surface enhanced Raman spectroscopy (SERS). Results show that SERS effectively enhanced the Raman signal from ε-Si layer and it stands distinctly apart from the Raman signal originating from the buffer layer. This technique is promising but it lacks the spatial resolution in the nano-scale due to diffraction limit from the probing light. In order to achieve nano-scale spectroscopy, point-surface-enhancement was used, rather than a large surface enhancement. The authors used a silver-coated sharp tip, just like SERS, but only the sample region very close to the tip apex is characterized. This technique, known as the tipenhanced Raman spectroscopy (TERS), provides nanometric resolution in our measurement. The authors observed localized strains by employing TERS. The TERS spectra revealed clear nano-scale variation in Raman frequency. Now that the authors can distinctively separate ε-Si from underlying buffer layer, signal-to-noise ration (SNR) needs further improvement. They improve TERS SNR in two ways: optical field enhancement using different metallic tip and background signals reduction arising from bulk materials. The tip-enhancement is more important for homogenous nano-materials or for samples with very weak signals whereas the background signal reduction is indispensable for nano-materials that consist of different thin layers with strong signals such as ε-Si or samples with strong signal level. Accordingly, the authors introduce several approaches mainly for the suppression of background signals arising from other bulk materials. The authors will discuss the utilization of UV light source, specialized tip, sample orientation relative to probing polarization, and depolarization configuration to obtain high contrast Raman signal. The characterization techniques describe above is applicable to other nano-materials. Cancer treatment usually uses drugs (chemotherapy) to reduce tumor size, followed by surgery to remove the tumor (if possible). Then, more chemotherapy and radiation therapy is used to kill as many tumor cells as possible. The goal of this collective treatment is to target and kill cancerous tissue while minimizing side effects on healthy cells. Due to their non specificity, current cancer therapies have poor therapeutic efficacy and can also have severe side effects on normal tissues and cells. In addition, cancer is often diagnosed and treated too late, i.e., when the cancer cells have already invaded and metastasized (i.e. spread) to other parts of the body. At this stage, treatment methods are highly limited in their effectiveness. Thus, scientists have been focusing efforts into finding alternative methods to detect cancer at earlier stages and kill such cancerous tissues more effectively. Nanoparticles (that is, particles with at least one dimension less than 100 nm) have become very attractive for improving

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cancer diagnosis and treatment due to their novel optical, magnetic and structural properties not available in conventional (or micron) particles or bulk solids. Nanoparticles have been extensively studied for various applications including delivering anti-cancer drugs to tumorous tissues and/or enhancing imaging capabilities to better diagnose and treat cancer. In Chapter 4, recent work related to the improved targeted therapy for specific cancers (whether by developing more specific anti-cancer agents or by altering delivery methods) are summarized. Discussions on the advantages and disadvantages of the most widely studied nanoparticles (i.e., liposome nanoparticles, polymer-based nanoparticles, quantum dots, nanoshells, and superparamagnetic particles) in cancer imaging followed by anti-cancer drug delivery are highlighted. Lastly, bone cancer and current research in using nanoparticles for treating bone cancer, with an emphasis on the novel use of selenium (a natural anti-cancer element found in our bodies), are addressed. The nanoindentation or indenter testing technique (ITT) is a functional and fast technique that can give us a lot of information about the mechanical properties of different materials at nanometric scale, from soft materials, such as copper, to brittle materials, such as ceramics. The principle of the technique is the evaluation of the response of a material to an applied load. In a composite material, if the size of the residual imprint resulting from a certain load is lower than the size of the studied phase, then is possible to determine its mechanical properties, and therefore its contribution to the global mechanical properties of the composite. Depending on the tipped indenter used, different equations should be applied to study the response of the material and calculate stress-strain curves and parameters such as hardness, Young’s modulus, toughness, yield strength and shear stress. These equations are related to the different deformation mechanisms (elastic, plastic or elastoplastic) that the material undergoes. In the case of most of the ceramic composites, when a spherical tipped nanoindenter is used, elastic deformation takes place, and Hertz equations can be used to calculate the yield stress, shear stress and the strain-stress curves. On the other hand, when a Berckovich indenter is used, plastic deformation takes place, then Oliver and Pahrr equations must be applied to evaluate the hardness, Young’s modulus and toughness. Nevertheless, in the hardness study, Indentation Size Effect (ISE) must be considered. In Chapter 5, the mechanical properties of a ceramic superconductor material have been studied. YBa2Cu3O7-δ (YBCO or Y-123) textured by Bridgman and Top Seeding Melt Growth (TSMG) techniques have been obtained and their mechanical properties studied by ITT. This material presents a phase transition from tetragonal to orthorhombic that promotes a change in its electrical properties, from insulating to superconductor, and that can be achieved by partially oxygenating the material. On the other hand, the structure of the textured material is heterogeneous, and two different phases are present: a Y-123 as a matrix and Y2BaCuO5 (Y-211) spherical inclusions. Moreover, the texture process induces an anisotropic structure, thus being the ab planes the ones that transport the superconductor properties while the c axis remains insulating. The purpose of this study is the characterization of the mechanical properties, in elastic and plastic range, of orthorhombic phases of YBCO samples textured by Bridgman and TSMG technique. With the ITT technique, the oxygenation process can be followed and its kinetics established. In Chapter 6, growth and physical properties of ZnO nanowire arrays were reviewed. It begins with some general remarks on semiconductor nanowires and basic properties of ZnO.

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In the second part, different kinds of growth methods that have been applied to grow ZnO nanowires are summarized. Vapor phase methods usually based on VSL or VS mechanism, depending on the presence or absence of a metal catalyst, were discussed in general. Typical solution methods for growth of ZnO nanowires were discussed separately as there is no common growth mechanism that can be applied to describe them. A new template-free strategy based on self-assembly process to grow ZnO nanowire into arrays were emphasized and discussed in detail. The obtained sample qualities were characterized with scanning electron microscopy, transmission electron microscopy, X-ray diffraction and energydispersive X-ray spectrum. The third part deals with the physical properties of ZnO nanowire arrays. Raman spectrum, including resonant Raman spectrum, was applied to test the crystal quality and phonon interaction of ZnO nanowires. Temperature-dependent photoluminescence spectra were measured to probe the intrinsic exciton and defect-related emission process of ZnO nanowires. Field emission properties of such ZnO nanowire arrays were also studied in view of the possible application for flat-panel displays. Some brief conclusions were summarized at the end. Organic materials (OMs) are diverse and are interesting in terms of application in electronic devices. In particular, organic charge transfer salts (OCTSs), which are typically composed of positive and negative ionic (and often radical) organic molecules, attract continued attention. Without carrier doping, they generally have high conductivity, magnetism, and well-defined unique nanostructures in their crystalline form. In order to apply the OCTSs to electronic devices, they should be made junction structures. Although there are established ways and advanced methods for doping and fabrication of junction structures in the current industrial techniques for the silicon devices, few of them are applicable to the OMs due to totally different chemical and physical properties between inorganic and organic materials. In Chapter 7, the authors would like to discuss a new method for simultaneous realization of doping and junction structures beginning with the single crystals of the OCTSs. The method utilizes a photo-induced chemical reaction, and produces a stable solid state composed of well-defined different parts of different conducting/magnetic properties. With reference to our recent and previous work as well as related studies of other groups, discussion will briefly cover experimental methods, preparation of materials, examination of irradiation conditions and resultant solids’ characterization, outline of mechanism of this photochemical modification, and remaining problems to be explained or overcome. Metal oxide nanowires exhibit novel properties due to their high surface-to-volume ratio and high surface stability. For this reason, they are considered excellent candidates to be incorporated into a new generation of devices with improved performance. Nevertheless, reaching complete control of their physical, chemical and electrical properties is needed before they can be widely used in our everyday life. This objective can be only fulfilled if reproducible electrical measurements on individual nanomaterials are performed. However, the fabrication of electrical nanocontacts in a fast and well-controlled process is still an unsolved issue. In Chapter 8, the main nanofabrication techniques that are commonly used to electrically access individual metal oxide nanowires, and to study their intrinsic properties are presented. Advantages and limitations of these methodologies are discussed in detail. By integrating bottom-up and top-down techniques, the first functional prototypes based on individual nanowires have already been implemented, paving the way to the future developments of nanoscale electronics, optoelectronics and chemical sensing devices.

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The inorganic-polymer hybrid nanomaterials have many excellent properties. So they are becoming increasingly important for various applications ranging from biomaterials to semiconductors in many fields and arouse much interest of scientists all over the world. Chapter 9 highlights the development of surface-initiated living radical polymerizations from the inorganic materials, including nanoparticles and one-dimensional (1D) nanostructures, by surface-initiated atom transfer radical polymerization (SI-ATRP). The emphasis is put upon the new developments of SI-ATRP taken to prepare hybrid nanomaterials in the recent years. Ferroelectric materials have been found to be promising candidates in applications of a wide range of electronic devices, such as high-dielectric constant capacitors, piezoelectric sonar or ultrasonic transducers, pyroelectric security sensors, medical diagnostic transducers, electro-optical light valves, and ultrasonic motors, and so on. Ferroelectric materials were conventionally fabricated via solid-state reactions at relatively and sometimes extremely high temperatures for calcining and sintering. Due to the presence of volatile components, such as lead (Pb), bismuth (Bi) or lithium (Li), in most ferroelectric compounds, high temperatures processing would brought out the problems of losing of the elements, which often resulted in the deteriorations in microstructures and thus electrical performances of the ferroelectric materials. To reduce the fabrication temperatures of ferroelectric ceramics, it is necessary to use ultrafine powders. High-energy mechanochemical technique, as an alternative method, has been used to synthesize nanosized ferroelectric powders directly from their oxide and other precursors. Chapter 10 serves as an overview of progress in the synthesis of various ferroelectric materials by using various mechanochemical milling facilities. In addition, applications of nanosized ferroelectric powders in materials preparation and device fabrication will be also be included. Linear forms of carbon are important in a wide variety of application, ranging from highly conducting interconnects to field emission materials. By methods of field ion microscopy (FIM) and mass-spectrometry, it was revealed the presence of linear carbon chains at the surface of carbon fibers after high-voltage treatment. The authors present in Chapter 11 a brief review of these research emphasizing recent developments. The carbon chains attached to the specimen tips can be produced in situ in a field ion microscope using low-temperature pulsed evaporation by electric fields of the order of 1011 V/m. Atomic Cchains are produced during the high-field unraveling of nanofibers. The experimental procedures used in FIM carbon chains studies are reviewed and the results in relation to the atomistics of unraveling processes are discussed. Molecular dynamics simulations and high resolution FIM experiments are performed to assess the evaporation of atomic chains under high-field conditions. Carbon exhibits a very rich dynamics of bond-breaking that allows transformation from graphenes to atomic chains. High-field experiments, theories leading to carbon chain formation, and methods to extract quantitative information on a variety of chainsurface interactions are described in detail. Isolated atomic carbon chains can be obtained at different temperatures, pulling speeds and forces. Current versus voltage field electron characteristics of monoatomic carbon wires were investigated. These results lend strong support to the conjecture of Smalley that linear carbon chains may provide the ultimate atomic-scale field emitters. Self-assembled nanostructures are new forms of materials which are interesting from a fundamental scientific perspective, as well as having many potential technological applications. As explained in Chapter 12, it is believed that the ability of nanostructures to self-assemble with controlled crystalline orientation, size, complexity and crystal

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morphology, provide potential applications in data storage, functional devices, communications and technology. Recently, a two-step strategy was successfully developed in our lab to produce two-dimensional or three-dimensional carbon nitride well-defined hierarchical complex structures. This strategy is a combination of a novel laser-induced deposition technique followed by self-assembly. In the first step, a suspension of carbon nitride nanoparticles was prepared by liquid-phase pulsed laser ablation (LP-PLA). In the second stage, this suspension was deposited onto a silicon substrate to act as a ‘seed’ layer. Via controlling the rate of evaporation of the liquid phase part of the seed suspension, and the size and the quantity of nanocrystals within the droplet, it was possible to create a range of nanoscale structures, including dense nanospheres, highly-symmetric flowers, hollow coreshell and uniform grass-like structures. The growth of such complex structures is governed by an evaporation-driven self-assembly process. As the droplet dries, small building blocks, such as nanoparticles (NPs) or nanorods (NRs) nucleate upon the existing crystals and template, sharing the same edges, to form a close-packed arrangement. By varying the design of the building blocks, materials combination, interfacial chemistry, and confining dimensions, it is expected to extend this synthetic approach to a range of new structured materials with useful functional properties. Self-standing diamond films were fabricated by a 30 kW DC Arcjet CVD system. The novel progresses, including layer-structured film (nano-/micro-crystalline layer) fabrication, high orientated film deposition with high growth rate at very high ratio of CH4/H2, crack-free thick and large area films growth, and single crystal fabrication, were reported. Layer-structured self-standing films, 2- and 4-layered ones, were fabricated by fluctuating the ratio of methane to hydrogen with deposition time. Results of scan electronic microscopy (SEM) and Raman spectra showed that the layered films were constructed by the micro-crystalline grains layer / nano-crystalline grains layer. The residual stress within the films were balanced, and even diminished in the certain layer. The layer containing nanocrystalline grains due to a plenty of secondary nucleation could weakly inherit the columnar growth feature of the overlaid layer containing micro-crystalline grains. The grain size and growth orientation of the layer containing micro-crystalline grains could be adjusted by introduction a mid-layer containing nano-crystalline grains. Growth rate was over 10μm/hr in layered film fabrication. The effect of very high concentration of CH4 in H2, 10%≤CH4/H2≤ 25%, was studied on the film morphology and orientation. Diamond films with morphology containing nice faceted micro-sized grains were obtained with CH4/H2 up to 17%. The film composition change was found by Raman spectra. High (111)-oriented films were deposited under the condition of CH4/H2=15% at the maximum growth rate about 50μm/h. Deposition temperature could influence both the morphology and orientation of the diamond films. The higher deposition temperature, the higher CH4/H2 could be allowed to deposit micro-sized grain-containing films. However, high deposition temperature would spoil (111)-orientation. As a consequence, (220) and (311) would be enhanced. Crack patterns occurring in self-standing films were classified as network shape, river shape and circle shape. The distribution and style of dominating crystalline surface was found to influence the strength of self-standing film. The films with 60-120mm of diameter and 2mm of thickness were successfully deposited by controlling of dominating crystalline surface in the films.

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A new approach to single diamond crystal fabrication by arc jet was proposed and discussed in Chapter 13. This method was named as “stable-tip” method which was applied to overcome the morphology instability. Single crystal, 1×1×0.6mm3 in size, was successfully fabricated by this method. The synchrotron radiation topography was adopted to characterize this single crystal diamond. Low emission efficiency of silicon (Si) based light emission devices (LED) still blocks application of Si-LED. Therefore, studies focusing on improving light emission of Si-LED still attract researches’ passion. Recently, the authors have developed a new method combining nanosphere lithography and pulsed laser deposition to fabricate Si-based arrays nanostructures, and have obtained remarkably enhanced photoluminescence (PL) from these structures. The Si based nanostructures are hemisphere shell arrays (HSSAs) or nanoflower arrays assembled by silicon-germanium (SiGe) alloy. These structures include non-closepacked and close-packed ones, single layer and multilayer ones, as well as arrays on different substrates. In Chapter 14, the authors investigated the photoluminescence of these arrays structures, and found that all these structures could enhance the photoluminescence intensities. Among them, the enhancement of light emission from SiGe double layer HSSAs (DL-HSSAs), which is as high as 700 folds, is the highest among those of all structures. Employing transmission electron microscopy (TEM), scanning electron microscopy (SEM), time-resolved PL, and electromagnetic simulation etc, the authors found the enhancement of light emission in Si based nanostructures originated mainly from the increase of extraction efficiency of photons from the nanostructures. The electromagnetic simulation of enhancement matched well the experiment data. The authors also found that these enhancements are related to degree of order of arrays. In highly order arrays, the enhancement is higher than that in other arrays. Trivalent lanthanide (Ln3+) ion-doped semiconductor nanocrystals have attracted extensive attention due to the ability to tailor their optical properties via size control and to achieve highly efficient luminescence through sensitization by the host. To date, finding a way to dope the “undopable” Ln3+ ions into semiconductor nanomaterials via chemical methods remains a challenge. In Chapter 15, recent progress in the doping of Ln3+ ions in TiO2 nanomaterials has been reviewed. A novel sol-gel-solvothermal method has been developed to effectively incorporate Ln3+ ions (Eu3+, Er3+, Nd3+ and Sm3+) into anatase TiO2 nanoparticles via the self-assembly and crystallization process of previous amorphous nanoparticles, in spite of a large mismatch in ionic radius and charge imbalance between Ln3+ and Ti4+. The crystallization process of Ln3+ doped TiO2 nanoparticles were systematically studied by means of thermogravimetric-differential thermal analyses (TG-DTA), powder Xray diffraction (XRD), and transmission electron microscope (TEM). Photoluminescence (PL) spectra of Ln3+:TiO2 samples exhibit resolved and sharp emission and excitation lines from the intra f-f transitions of Ln3+ ions (Ln=Nd, Sm, Eu, Er), indicating regular crystalline surroundings of Ln3+ ions. Multiple sites of Eu3+, Sm3+ and Nd3+ ions in anatase TiO2 were detected by means of high- resolution site-selective spectroscopy at 10K, whereas only single site emission of Er3+ in TiO2 were observed. Very intense near-infrared luminescence around 1.53 µm was also observed, which originated from the single lattice site of Er3+ ions incorporated in TiO2 nanocrystals. The luminescence dynamics and CF levels of Ln3+ at different sites have been analyzed. Highly efficient emissions of Nd3+ and Sm3+ sensitized by the TiO2 host were observed upon the excitation above the TiO2 band gap energy at room

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temperature (RT), which is of particular interest for material applications. A growth mechanism for the incorporation of Ln3+ in the anatase lattice is also suggested. Titanium nitride (TiN) films synthesized by multi-arc ion plating (AIP) normally have a columnar microstructure, and are likely to induce surface defects due to the formation of macroparticles and neutral particles in the vicinity of cathode arc sources. Hence, the achievable microhardness of the normal AIP TiN films only ranges between 20~30 GPa. A systematic study for fabricating an adherent nano-superhard titanium nitride (TiN) film on M2 high speed steel substrate by a vacuum cathode multi-arc ion-plating (AIP) system was initiated. To understand the relationship of the film processing-structure-property, their microhardness, film-to-substrate adhesion, frictional property, and microstructure of the film were investigated using Vickers hardometer, scratch tester, ball-on-disc tester, X-ray diffractometer, and transmission electron microscope. Results in Chapter 16 show that: (i) the achievable film microhardness ranges between 35 GPa and 45 GPa; (ii) the critical load (Lc) of the superhard TiN film is at 64 N approximately; (iii) the friction coefficient, under a highload and a high rotating-speed, of the film is ranging from 0.5 to 0.8; and (iv) the nm scale mean main grain-sizes of the film are approximately 12.7 nm for TiN111, 19.7 nm for TiN200 and 9.6 nm for TiN220. The maximum achievable microhardness 45 GPa is more than twice of the 22 GPa for standard TiN film. Such hardness enhancement is anticipated as mainly due to: (a) the formation of nanoscaled crystalline grains; (b) the preferential orientation and growth of grains in the close-packed plane (111); and (c) the induced residual stress within the film by ion bombardment. In Chapter 17, nanoparticles embedded in insulators, e.g., Al2O3, MgO, YSZ and TiO2 single crystals, were fabricated by ion implantation and subsequent thermal annealing, including metallic Ni, Zn and their oxides, and intermetallic nanoparticles. Optical, magnetic and mircostructural properties of nanoparticles have been studied. The metallic nanoparticles have surface plasmon resonance absorption, and oxide nanoparticles show good photoluminenscence. The magnetic nanoparticles, e.g., metallic Ni and intermetallic CoxNi1-x nanoparticles, show strong ferromagnetism behaviors. The ion fluence can affect the concentrations and the intensities of the surface plasmon absorption of metallic nanoparticles. Ion flux is another important parameter to fabricate nanoparticles. An example of effects of ion flux on the nanoparticles has been presented in this data review. The relationship between annealing temperature and optical, magnetic and microstructural properties of nanoparticles has also been systematically studied. Ion implantation provides a versatile and powerful technique for synthesizing nanometerscale clusters embedded in the near-surface region of a variety of host materials. The embedded nanoparticles have attracted considerable attention because of their unique opticalelectrical properties that are different from those of the bulk matrix. Metallic nanoparticles embedded in insulators have pronounced optical effects, including surface plasma resonance (SPR) absorption, and strong third-order nonlinear optical (NLO) susceptibility. The former suggests applications as optical filters, including eye-glass coatings. The latter has potential application in all-optical-memory or switching devices. Oxide nanoparticles have good photoluminescence. They have promising application in light-emitting devices. Magnetic metallic nanoparticles often show a ferromagnetic behavior with a larger coercivity than that of the corresponding bulk materials, which may provide potential application of the nanocomposite as magneto-optical materials for a high density magnetic data storage device.

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Organized self-assembly of molecules, driven by noncovalent intermolecular interactions, is the most versatile tool for accessing new materials with desired optical and electronic properties. Porphyrins are particularly attractive species to incorporate into supramolecular assemblies because their rich photochemistry may impart functionality, provide insight into the mechanisms of biological processes such as photosynthesis, serve as probes into the features of self-assembled structures and as models for molecular organization and energy/electron transfer processes. The close molecular packing in a self-assembled porphyrin aggregate leads to different electronic coupling and delocalization of the excitation energy, which can be exploited for applications in non-linear optical devices, photoelectric cells, recording devices. The possibility to control and tune either shape and size of the porphyrin clusters opens the way for their use as potential nanodevices. Chapter 18 aims to collect some recent developments in the field of porphyrin self-assembly and to frame all the reported topics into the current theories. The influence of thiophene addition on the pyrolysis of poly(dimethyl siloxane) catalyzed by ferrocene at ~1050 oC in Ar was studied in the first Short Communication. The assynthesized product was characterized by X-ray diffraction, scanning electron microscopy, transmission electron microscopy and high-resolution transmission electron microscopy. The thiophene addition caused several changes. Firstly, the yield of the product was increased by several times and the diameters of the product were somewhat increased. Secondly, the product was changed from only SiC/SiO2 nanocables to a mixture of SiC/SiO2 nanocables and SiC-SiO2 side-by-side nanowires. Thirdly, more “Y” type nanostructures were found. Finally, the growth process of the product was altered as the nanostructures each had a polyhedral FeS nanoparticle rather than spherical Fe nanoparticle. However, lengths of the product were still on the millimeter scale. The promotion mechanism of thiophene addition was also analyzed. As discussed the second short communication, nanotechnology revolutionized every field in science and technology. Recently, its usefulness in nanofinishing of cotton fabrics by imparting functional properties like antimicrobial, UV-resistance, self-cleaning and drugdelivery is well documented. In addition, enhancement in comfort properties of cotton textiles is also being evaluated with the help of nanofinishing. With judicial use of nanomaterials, keeping in view their bio-safety and environmental impact issues, nanofinishing will be a great boon to the users of cotton textiles.

RESEARCH AND REVIEW STUDIES

In: Nanotechnology… Editors: C.J. Dixon and O.W. Curtines, pp. 3-50

ISBN: 978-1-60692-162-3 © 2010 Nova Science Publishers, Inc.

Chapter 1

ELECTROCHEMICAL NANOFABRICATION Di Wei Nokia Research Centre, c/o Nanoscience Centre at University of Cambridge, 11 JJ Thomson Avenue, CB3 0FF, Cambridge, UK

Abstract Nano- and micro-fabrications have been largely used in the applications such as integrated circuits, micro/nano electro-mechanical systems (M/NEMS), micro-optics and countless others. The methodology of nanofabrication can be divided into two types, top-down and bottom-up processes, which themselves can be further divided. Top-down process refers to approaching the nanoscale from the top (or larger dimensions), such as lithography, nanoimprinting, scanning probe and E-beam technique etc.. In bottom-up fabrication processes, the nanotechnology process builds nanoscale artifacts from the molecular level up, through single molecules or collections of molecules that agglomerate or self-assemble. Using a bottom-up approach, such as self-assembly enables scientists to create larger and more complex systems from elementary subcomponents (e.g. atoms and molecules). In general, topdown processes that transfer minute patterns onto material are more matured than bottom-up processes. An exception is epitaxial processes that create layers through layer-by-layer growth with registry at the atomic level. Electrodeposition has actually been used for decades to form high quality, mostly metallic, thin films. It has recently been shown that high quality copper interconnects for ultra large scale integration chips can be formed electrochemically on Si wafer [1;2]. Electrodeposition has thus been shown compatible with state of the art semiconductor manufacturing technology. The largest semiconductor companies, for example, IBM, Intel, AMD, Motorola etc. are installing wafer-electroplating machines on their fabrication lines [1]. The electrodeposition of Cu with the line width 250 nm was used in the mass-production of micro-processor Pentium III in 1998. In 2003, the line width of the CPU was reduced to 130 nm in Pentium IV. Electrochemistry was largely used in chip fabrication [3] and the packaging of micro-electronics [4]. However, comparing with other nanofabrication techniques, electrochemical nanofabrication is still a maiden area which needs further development and fulfilment.

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Di Wei This chapter summarized the most recent developments in electrochemical nanofabrications. It includes not only the conventional technique, under potential deposition (UPD), which deals with the deposition of a single metal-ion on a definite substrate but also some new developments using ultrashort voltage pulsing and template methods for 3D construction of nano-materials. Electrochemical nanofabrication is a versatile method, which includes both top-down nanofabrication (e.g. electrochemical lithography) and bottom-up process such as electrochemical atomic layer epitaxy (EC-ALE). Nano-templates including anodized aluminum oxide (AAO) membranes, colloidal polystyrene (PS) latex spheres, single/aligned carbon nanotubes, selfassembled monolayers (SAMs), blocked copolymers and cyclodextrin molecules can be used for the preparation of various types of nanowires, nanotubes, ordered arrays of nanoparticles and nanodots electrochemically. Combining electrochemistry with other nanofabrication techniques such as focused ion beam (FIB) and self-assembly provides many novel strategies in the fabrication of nanomaterials with specific design. Selective areas in the nanoscale can be modified by electrochemical nanostructuring with metals, metal oxides and conducting polymers using a bipolar electrochemical technique. The traditional lithography and pattern technique is costly. In the construction of soft matters such as conducting polymers, traditional spin casting cannot guarantee nanostructures due to the fast speed of solvent evaporation. Electrochemical technique provides an innovative, versatile and economic way of nanofabrication. It especially offers better alternative to construct the soft matter nano-structures in a controllable manner. In general, electrochemical nanofabrication offers simplicity, efficiency, low-temperature processing, cost-effectiveness, the possibility in preparing large area deposits and precise control of the deposit thickness, which are the essential advantages than other nanofabrication techniques till date. Additionally, it can be used to prepare a wide range of materials comprising the inorganic and the organic. The former includes quantum dots, metallic and semiconducting (e.g. ZnO, TiO2) nanotubes and nanorods. The latter includes conducting polymer nanotubes and nanowires.

Electrochemical Deposition Under Potential Deposition (UPD) Surface limited reactions are well known in electrochemistry and are generally referred to as under potential deposits [5;6]. Underpotential deposition (UPD) is the formation of an atomic layer of one element on a second element at a potential under, or prior to, that needed to deposit the element on itself. The shift in potential results from the free energy of the surface compound formation. ;Early UPD studies were carried out mostly on polycrystalline electrode surfaces [7]. This was due, at least in part, to the difficulty of preparing and maintaining single-crystal electrodes under well-defined (and controlled) conditions of surface structure and cleanliness [8]. For example, cadmium (Cd) can be underpotentially deposited on Cu(111) and Cu(100) [9]. There are a number of excellent reviews on this topic [5;6].

Electrochemical Atomic Layer Epitaxy (EC-ALE) Electrochemical atomic layer epitaxy (EC-ALE) is the combination of UPD and ALE, which uses UPD for the surface limited reactions in an ALE cycle [10]. Fundamental to forming high quality structures and devices with thin films of compound semiconductors is the concept of epitaxy. The definition of epitaxy is variable, but focuses on the formation of

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single crystal films on single crystal substrates. This is different from other thin film deposition methods where polycrystalline or amorphous film deposits are formed even on single crystal substrates. Homoepitaxy is the formation of a compound on itself. Heteroepitaxy is the formation of a compound on a different compound or element, and is much more prevalent. The principle of ALE is to grow the deposit one atomic layer at a time [11]. Surface limited reactions are developed for the deposition of each component element in a cycle to directly form a compound via layer-by-layer growth, avoiding 3D nucleations. With cycle, a compound monolayer is formed, and the deposit thickness is controlled by the number of cycles. In an EC-ALE process, each reactant has its own solution and deposition potential, and there are generally rinse solutions as well. Control of growth at the nanoscale is a major frontier of materials science. The manipulation of a compound’s dimensions, or unit cell, at the nanoscale, can result in materials with unique properties. By constructing superlattices, nanowires and nanoclusters or forming nanocrystalline materials, the electronic structure (bandgap) of a semiconductor can be engineered. EC-ALE has been developed as an electrochemical methodology to grow compound semiconductors with nanoscale or atomic layer control. In an EC-ALE synthesis of CdTe, for example, atomic layers of tellurium and cadmium are alternatively electrodeposited to build up a thin layer of CdTe [12;13]. The necessary atomic level control over the electrodeposition of these two elements is obtained by depositing both elements using UPD. The thickness of the CdTe layer prepared by EC-ALE can be specified by controlling the number of Cd and Te layers that are deposited.

Figure 1. Transmission electron micrograph (TEM) of a CdTe deposit formed using 200 cycle of CdTe via EC-ALE [14]. Reproduced by the kind permission from the publisher.

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Fig. 1 shows the TEM figure of the CdTe usng EC-ALE with 200 cycles [14]. The regular layered structure, parallel with the substrate Au lattice planes, suggests the epitaxial nature of the deposit. There are a number of ways to introduce dopants into an EC-ALE deposit and they can be introduced homogeneously throughout the deposit. Initial doping studies of ZnS were run with the idea of forming phosphor screens for flat panel display applications [15].

Electrochemical Deposition Methods for Semiconducting Nanocompounds Electrodeposition normally leads to small particle size, largely because it is a low temperature technique, thereby minimizing grain growth. It, however, possesses the additional feature of a very high degree of control over the amount of deposited material through Faraday’s law, which relates the amount of material deposited to the deposition charge. This feature is particularly desirable when isolated nanocrystals are to be deposited on a substrate. Several methods and variations have been developed to electrodeposit compounds. Oxides are probably the largest group of electrodeposited compounds (for example, aluminium anodization). The electrodeposition of II-VI compounds has been extensively studied and is well reviewed in a number of articles [16-18]. A number of reviews of semiconductor electrodeposition also exist which describe the various methods used [19;20]. The most prominent electrodeposition methods for semiconducting compounds include: codeposition, precipitation and various two-stage techniques. Semiconductor film can be electrodeposited either by EC-ALE or by co-deposition [21]. The most successful methodology to form II-VI compounds has been codeposition [22-26], where both elements are deposited at the same time from the same solution. Stoichiometry is maintained by having the more inactive element as the limiting reagent, and poising the potential where the less noble element will underpotentially deposit only on the more noble element. The classic example is CdTe formation [22], where the solution contains Cd2+ and HTeO2+, usually at pH 2. The potential is set to reduce HTeO2+ to Te on the surface at a limiting rate, while Cd2+ is reduced on the Te at an under potential, a potential where no bulk Cd is formed. Cd2+ ions are present in a large excess, to deposit quantitatively on Te as it is formed, resulting in stoichiometric CdTe. Although the structure and morphology of codeposited compounds are variable, some having been described as ‘cauliflower’ like, high quality deposits have been formed [27]. There are a number of papers in the literature concerning the formation of compound semiconductor diodes by electrodeposition, the most popular structure being a CdS-CdTe based photovoltaic. CdS was generally deposited first on an ITO/glass substrate, followed by a layer of CdTe, usually by codeposition [28-34]. Sailor and Martin et al. grew an array of CdSe-CdTe nano-diodes in 200 nm pore alumite [35-37], using a compound electrodeposition methodology called sequential monolayer electrodeposition [38]. A commercial process is being developed by BP Solar to form CdTe based photovoltaics using codeposition. Relatively rapid deposition rate has been achieved by codeposition and it is presently the most practical compound semiconductor electrodepostion methodology. Codeposition holds great promise if greater control can be achieved. At present the main parameters of control are solution composition and the deposition potential. There have

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been a number of attempts to improve the process by using variations in reactant concentration, pH [39;40], and the potential program [38;41-45]. In most cases, the deposits are improved by annealing. In the application of photovoltaic applications, annealing is used to convert CdTe from the as-deposited n-type material to the desired p-type [34;46]. Semiconductors such as polycrystalline ITO on glass have been used to form deposits of ZnS with no obvious problems [15]. Ideally lattice matched semiconductor substrates could be used to form deposits. For instance, InSb is lattice matched with CdTe and could be used as a substrate. Good quality deposits of CdSe have been formed on InP and GaAs substrates using codeposition by Maurin and Froment et al. [47;48]. Their work clearly show the applicability of high quality commercial compound semiconductors wafers as substrates for compounds electrodeposition. The precipitation method involves electrochemical generation of a precursor to one of the constituent elements, in a solution containing precursors to the other elements [49-52]. The reaction is essentially homogeneous, but as one reactant is formed at the electrode surface, most of the product precipitates on the surface. This method resembles passive oxide film formation on reactive metals, where metal ions react with the solvent, oxygen or hydroxide. The film thickness is controlled by the amount of electrogenerated precursor. However, as the method resembles precipitation, the quality of the resulting deposit is questionable, and the process is difficult to control. Film thickness is necessarily limited by the need for precursor transport through the deposit. A classic example is the formation of CdS by oxidizing a Cd electrode in a sulphide solution [49;53-57]. Two stage methods are where thin films of the compound element, or an alloy, are first deposited, at least one by electrodeposition [58]. A second stage, annealing, then results in inter-diffusion and reaction of the elements to form the compound. The deposits are annealed in air, inert gas, or a gaseous precursor to one of the compound’s component elements. For instance, electrodeposited CuIn alloys have been annealed in H2S to form CuInS2 [59]. Given the need for annealing, this methodology has limitations for the formation of more involved device structures. In general, annealing has been used to either form or improve the structures of compound films formed by the electrodeposition methods described above. The primary tool for understanding compound electrodeposition and for improving control over the process has been the methodology of EC-ALE [60;61].

Electrochemical Synthesis of Quantum Dots Quantum dots are semiconductor particles having diameters that are smaller than about 10 nm. Such semiconductor nanoparticles exhibit a bandgap that depends on the particle diameter: the smaller nanoparticle, the larger the bandgap. Because quantum dots possess a ‘size-tunable’ bandgap, these diminutive particles have potential applications in detectors, light emitting diodes, electroluminescent devices, and lasers. Electrochemical methods can synthesize size-monodisperse quantum dots on graphite surfaces, which provide an electrical connection to the graphite in situ. The essential features of these methods can be depicted as in Fig. 2.

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Figure 2. The electrochemical and chemical method to synthesize the quantum dots and other semiconductor nanocrystals on graphite.

The first step involves the electrodeposition of metal nanoparticles onto a graphite surface from a solution containing the corresponding metal ions. The metal nanoparticles are electrochemically oxidized to yield a metal oxide (MO), in which the oxidation state of the metal matches the oxidation state in the final product. Finally metal oxide nanoparticles are converted into nanoparticles of a semiconducting salt (MX) via a displacement reaction in which oxide or hydroxide is replaced by the desired anions (X). Examples of using these methods can be shown in the synthesis of CuI [62], CdS [63] and ZnO [64] quantum dots. Ultrathin films of quantum dots with deposits of non-connected nanocrystals and thick films of more than 10 nm in average thickness can be made by electrochemical methods [65-67].

Electrochemical Deposition Methods for Metallic Nanostructures Inorganic nanoparticles can be fabricated by many different techniques. Electrochemical and wet-chemical methods are demonstrated to be effective approaches to make metal nanostructures under control without addition of a reducing agent or protecting agent. An in situ electrochemical reduction method for fabricating metal nanoparticles on carbon substrates simultaneously assembling into ordered functional nanostructures was developed [68]. Ag+ was adsorbed on a highly oriented pyrolytic graphite surface modified by 4aminophenyl monolayer with coordination interaction, and then homogeneously dispersed Ag nanoparticles could be obtained through pulsed potentiostatic reduction. Multilayered metal nanostructures on glassy carbon electrodes have been obtained by extending this method [69;70]. The larger electrochemical window of ionic liquids in comparison of aqueous electrolytes enables the investigation of electrodeposition of metal and semi-conductor elements and compounds in nanoscale. Nanoscale electrocrystallization of metals such as Ni, Co and the electrodeposition of semiconductors (Ge) on Au (111) and Si (111):H have been studied in the underpotential and overpotential range from ionic liquids [71]. For example, 3D growth in Co electrodeposition on Au (111) from ionic liquids based on imidazolium cations starts at potentials below -0.17 V vs. Co/Co(II). 3D and 2D structures of Co and Ni deposition in the nanoscale were illustrated. In addition, nanocrystalline aluminium can be obtained by electrodeposition from ionic liquids containing imidazolium cations without additives[72].

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The crystal refinement is due to a cathodic decomposition of the imidazolium ions to a certain extent giving rise to nanocrystalline aluminium. Metal nanowires can be obtained using solution phase reduction [73], template synthesis [74-77], and physical vapour deposition (PVD) [78] onto carbon nanotubes. Metal wires with widths down to 20 nm and lengths of millimetres can be prepared on silicon surface using electron beam lithography [79] or by PVD [80]. However, none of these methods are useful to prepare free-standing metal nanowires that are longer than 20 μm. Penner et al. [81] have used the step edge defects on single crystal surfaces as templates to form metal nanowires by electrochemical step edge decoration. Metallic molybdenum (Mo) wires with diameters ranging from 15 nm to 1 μm and lengths of up to 500 μm were prepared in a two-step procedure on freshly cleaved graphite surfaces [82]. Molybdenum oxide wires were electrodeposited selectively at step edges at -0.75 VSCE and then reduced in hydrogen gas at 500 °C to yield Mo metal. Such nanowires can be obtained size selectively because the mean wire diameter was directly proportional to the square root of the electrolysis time. Parrallel arrays of long (> 500 μm), dimensionally uniform nanowires composed of molybdenum, copper, nickel, gold, and palladium can also be electrodeposited by the same strategy [81]. They were firstly prepared by electrodepositing nanowires of a conductive metal oxide such as NiO, Cu2O or MoO2. Nanowires of the parent metal were then obtained by reducing the metal oxide nanowires in hydrogen at elevated temperature. Nanowires with diameters in the range from 15 nm to 750 nm were obtained by electrodeposition onto the step edges present on the surface of highly oriented pyrolytic graphite electrode. After embedding the nanowires in a polymer film, arrays of nanowires could be lifted off the graphite surface thereby facilitating the incorporation of these arrays in devices such as sensors. Vertical arrays of metal nanowire hold promise for making chemical and biological sensors in addition to electron emitters in field-emission displays. But the difficulty of growing well-defined arrays has kept these technologies at bay. Electrochemical nanofabrication using crystalline protein masks solved this problem [83]. A simple and robust method was developed to fabricate nanoarrays of metals and metal oxides over macroscopic substrates using the crystalline surface layer (S-layer) protein of deinococcus radiodurans as an electrodeposition mask. Substrates are coated by adsorption of the S-layer from a detergent-stabilized aqueous protein extract, producing insulating masks with 2-3 nm diameter solvent-accessible openings to the deposition substrate. The coating process can be controlled to achieve complete or fractional surface coverage. The general applicability of the technique was demonstrated by forming arrays of Cu2O, Ni, Pt, Pd, and Co exhibiting longrange order with the 18 nm hexagonal periodicity of the protein openings. This protein-based approach to electrochemical nanofabrication should permit the creation of a wide variety of two-dimensional inorganic structures.

Electrochemical Nanolithography In addition to its well-known capabilities in imaging and spectroscopy, scanning probe microscopy (SPM) has shown great potentials for patterning of material structures in nanoscales with precise control of the structure and location. Electrochemical nanolithography using SPM, which includes scanning tunnelling microscopy (STM) and atomic force microscope (AFM), has been used to fabricate of patterned metal structures

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[84-87], semiconductors [87;88] and soft matters such as conducting polymers [89;90]. The electrochemical processing of material surfaces at nanoscale both laterally and vertically can be conducted by scanning probe anodization/cathodization, which used the tip-sample junction of a scanning probe microscope connected with an adsorbed water column as a minute electrochemical cell. A review on nanofabrication by scanning probe microscope lithography examines various applications of SPM in modification, deposition, removal, and manipulation of materials for nanoscale fabrication [91]. Comprehensive reviews of SPMrelated lithography can be found in the literature [92]. STM has a tremendous potential in metal deposition studies. The initial stages of metal deposition and the Ag adlayer on Au (111) have been studied by Kolb et al. [93].The inherent nature of the deposition process which is strongly influenced by the defect structure of the substrate, providing nucleation centres, requires imaging in real space for a detailed picture of the initial stage. This is possible with an STM, the atomic resolution helps to understand these processes on a truly atomistic level. The following figure demonstrates wealth of structural detail supplied by STM. In situ STM investigations for Ag UPD on Au (111) in the potential range from 600 to 200 mV vs. Ag/Ag+ were carried out [93]. Underpotentially deposited Ag revealed a series of ordered adlayer structures with an increasing density of adatoms when decreasing applied potentials. Nanostructuring can be also achieved by involving electrochemical processes into the overall procedures. For example, Li et al. [94;95] deposited Ag and Pt clusters on a graphite surface by applying positive voltage pulses to the STM tip in a solution containing the respective metal cations. This effect was attributed to nucleation within holes which were created on the graphite substrate surface by the voltage pulses. Kolb et al. [96], on the other hand, were able to detach Cu clusters from a tip, where they had been previously deposited electrochemically, onto a Au substrate by mechanical contact. However, all these techniques suffer from restrictions, which could be largely avoided if controlled nanostructuring could be achieved by a direct local electrochemical reaction on the substrate, with the geometry being determined by the location of the tip which acts as a local counter electrode. By applying ultrashort voltage pulses (≤ 100ns), holes of about 5 nm in diameter and 0.3 to 1 nm depth on Au substrate can be created by local anodic dissolution, while cathodic polarization led to the deposition of small Cu clusters [97]. The development of allowing the generation of small metal clusters, with the help of an STM tip, and placing them at will onto single crystal electrode surfaces was reported [96;98]. This so-called tip-induced metal deposition involves conventional electrolytic deposition of a metal onto the tip of an STM, followed by a controlled tip approach during which metal is transferred from the tip to the surface (so-called jump-to-contact [99]). Small copper clusters, typically two to four atomic layers in height, were precisely positioned on a Au(111) electrode by a process in which copper was first deposited onto the tip of the STM, which then acted as a reservoir from which copper could be transferred to the surface during an appropriate approach of the tip to the surface [100]. Tip approach and position were controlled externally by a microprocessor unit, allowing the fabrication of various patterns, cluster arrays, and ''conducting wires'' in a very flexible and convenient manner [96]. The formation of such clusters with the tip of a STM is simulated by atom dynamics and subsequently the stability of these clusters is investigated by Monte-Carlo simulations in a grand-canonical ensemble. It leads to the conclusion that optimal systems for nanostructuring are those where the metals participating have similar cohesive energies and

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negative heats of alloy formation. In this respect, the system Cu-Pd(111) is predicted as a good candidate for the formation of stable clusters [101]. In addition to producing the metal nano-clusters, STMs can also be applied to form nanoscale pits in thin conducting films of thallium (III) oxide [102] as well as to write stable features on an atomically flat Au (111) surface [103]. Pit formation was only observed when the process was performed in humid ambient conditions. The mechanism involved in pit formation was attributed to localized electrochemical etching reactions beneath the STM tip. By applying voltage pulses (close to 3 V) across the tunnelling junction in controlled atmosphere with the presence of water or ethanol vapour, nano-hole can be produced. The smallest hole formed is 3 nm in diameter and 0.24 nm in depth. This nano-hole represents the loss of about 100 Au atoms in the top atomic layer of gold surface, there is no atomic perturbation seen inside and outside the nano-hole. Different nanostructures (lattice of dots, legends, map, etc.) can also be fabricated. The threshold voltage for the formation of a nanohole depends on the relative humidity, however, the relationship between the threshold voltage and the relative humidity is basically independent of the tip material. The application of conducting AFM probe anodization to nanolithography was also used in the fabrication and patterning of materials in a similar manner as STM probes. AFM-tipinduced and current-induced local oxidation of silicon and metals were reported and this novel local oxidation process can be used to generate thin oxide tunneling barriers of 10-50 nm.[88]. The direct modification of silicon and other semiconductor and metal surfaces by the process of anodization using the electric field from a SPM is one promising method of accomplishing direct-writing lithography for the electron device fabrication. This technique involves the application of an electrical bias to both the conducting probe and the sample substrate to locally oxidize selected regions of a sample surface. Since most of lithographic works with organic resists have been exclusively carried out on silicon surface for practical application, so several organic resists with different function groups have been studied in order to investigate the surface group effect on the anodic anodization using AFM. Silicon (Si) samples whose surfaces were terminated with hydrogen (Si-H samples) or covered with an organosilane monolayer were generally used to deposit metallic and/or semiconductor nanomaterials [86]. Silicon oxide (SiO2) patterns can be prepared on the Si-H sample surfaces by the use of anodization of Si, while in the second case the organosilane layer was selectively degraded by anodic corrosion. Furthermore, pattern transfer processes that fabricated metal nanostructures using these patterned SiO2 or organosilane layers as templates were developed. These processes are based on area-selective electroless plating where selectivity depends on the difference in the chemical reactivity between the surface modified by scanning probe anodization and the unmodified surface. Nanostructures down to a few ten nanometers in size have been fabricated with Langmuir–Blodgett (LB) films and self-assembled monolayers (SAMs) using SPM lithography [104]. The SAMs can be prepared with organosilane etc. as ultrathin resists on Si substrate. The effect of such functional groups of molecules on the AFM anodization, which was performed under contact mode has been studied in the optimized process conditions. Applied voltage between the AFM tip and sample, the scanning speed and the relative humidity in air are also important factors for nanometer-scale lithography of the ultrathin films. The high structural orderness and perfect thickness of ultrathin organic molecular assemblies are the major advantages as required for nanoscale lithography.

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Cleaned p-type Si (100) wafer was firstly oxidized by piranha solution (3:1 mixture of sulfuric acid and 30% hydrogen peroxide) to make the surface hydrophilic. The octadecyldimethylmethoxysilane (ODMS) molecules react with OH groups on a silicon oxide surface resulting in the formation of a SAM in a thickness of 1.7 nm. The local degradation of the monolayer on selected area where the probe tip of the AFM was scanned with a bias voltage occurred due to the anodic reaction. The degraded regions became hydrophilic indicating that the ODMS molecules were decomposed and replaced with OH groups as a result of the probe scan. AFM images showed that the tip scanned regions were protruded compared to the surrounding regions as Fig.3 shows. This is due to volume expansion resulting from anodization of the Si substrate, which immediately followed the tip-induced degradation of the SAM.

Figure 3. Anodized pattern formation on Si-wafer. Si wafer was firstly oxidized to form hydrophilic oxide layer for the selfassembly of ODMS. ODMS SAMs act as the resist. Reproduced by the kind permission from the publisher.

Figure 4. The AFM image of ODMS monolayer after AFM anodization at the different applied currents [104]. Reproduced by the kind permission from the publisher.

Fig. 4 shows the AFM image of ODMS monolayer after AFM anodization at the different current conditions. Typically with ODMS monolayers, protruded patterns are fabricated with line-width of 70 nm at a scan rate of 120–200 μm/s and an applied voltage of 20–25 V. The

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height of protruded lines can be controlled by changing the current between the tip and substrate. When the applied current was 5–7 nA, the height of protruded line was 1.0 nm. With increasing the current up to 20–23 nA, the anodized height increased to 3.0 nm. Thus, the height of the protruded lines increases due to the effective electric field strength for a given applied voltage and a scan rate. The electric field plays an important role in the formation of protrusion. AFM anodization was successfully carried out with SAMs on the silicon substrate. Applied voltage between the AFM tip and sample, the scanning speed, surface group, and the relative humidity in the laboratory are very important factors for nanometer-scale lithography of the ultrathin films. The high structural orderness and perfect thickness of ultrathin organic molecular assemblies are the major advantages as required for nanoscale SPM lithography. Soft matters such as conducting polymers can also be patterned in the nanoscale by such lithographic method. Cai et al. described the first observation of localized electropolymerization of pyrrole and aniline on highly oriented pyrolytic graphite (HOPG) substrates under AFM tip-sample interactions [89]. A scanning or oscillating AFM tip, providing the horizontal scratching force and the vertical tapping force, is essential as the driving force for the surface modification with the conducting polymer. It was shown that under the AFM tip interaction, the electropolymerization can be blocked on the bare HOPG substrate or enhanced on the as-polymerized film. The localized electropolymerization in selected surface areas enables the nanomodification of lines, square platforms, or hollows of polypyrrole and polyaniline on the substrates. The result indicates that AFM can be used as a unique tool for nanofabrication of conducting polymers. Nano-writing of intrinsically conducting polymer was also achieved via a novel electrochemical nanolithographic technique using tapping mode electrochemical AFM. Conducting polymer (polythiophene derivatives) nanolines as small as 58 nm in width were obtained and the line width is controlled as a function of the writing speed and writing potential [90].

Figure 5. Electrochemical nanolithography [90]. Reproduced by the kind permission from the publisher.

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The electrochemical nanolithography process is shown in Fig. 5. Conductive AFM probes, gold coated silicon nitride (SiN4) are used as working electrode. Silver wire and platinum wire were used as reference electrode and counter electrode, respectively. Higher writing potential and slower writing speed produce wider conducting polymer nanolines due to enhanced propagation. The great benefit of this method lies in no specific restriction in the choice of substrates and the ease of controlling feature size, which is expected to facilitate to fabrication of all plastic nanoelectronic devices. As stated previously, many SPM lithography techniques based on anodization of Si surfaces [88], electrochemical reactions in solution using electrochemical STM tips [96;98] and electrochemical decomposition of self-assembled monolayers [104] have been developed in the past decade. More recently, a “dippen” nanolithography (DPN) method was invented that uses an AFM tip as a “nib” to directly deliver organic molecules onto suitable substrate surfaces, such as Au [105-107]. When AFM is used in air to image a surface, the narrow gap between the tip and surface behaves as a tiny capillary that condenses water from the air. This tiny water meniscus is actually an important factor that has limited the resolution of AFM in air. “Dip-pen” AFM lithography uses the water meniscus to transport organic molecules from tip to surface. By using this technique, organic monolayers can be directly written on the surface with no additional steps, and multiple inks can be used to write different molecules on the same surface. By coupling electrochemical techniques, the DPN are not limited to deliver organic molecules to the surface. Electrochemical “dip-pen” nanolithography (EC-DPN) technique can be used to directly fabricate metal and semiconductor nanostructures on surfaces.

Figure 6. Schematic sketch of the EC-DPN experimental setup [108]. Reproduced by the kind permission from the publisher.

The tiny water meniscus on the AFM tip was used as a nano-sized electrochemical cell, in which metal salts can be dissolved, reduced into metals electrochemically and deposit on the surface as shown in Fig. 6. In a typical experiment, an ultrasharp silicon cantilever coated with H2PtCl6 is scanned on a cleaned p-type Si (100) surface with a positive DC bias applied on the tip. During this lithographic process, H2PtCl6 dissolved in the water meniscus is electrochemically reduced from Pt(IV) to Pt(0) metal at the cathodic silicon surface and deposits as Pt nano-features as shown in the result below [108].

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Figure 7. AFM image and height profile of two Pt lines drawn at different scan speed. a) line at 10 nm/s and b) line at 20 nm/s. The voltage applied at the tip is 3V for both lines and the relative humidity is 43% [108]. Reproduced by the kind permission from the publisher.

Electrochemical AFM "dip-pen'' nanolithography has significantly expanded the scope where DPN nanofabrication can be applied. It combines the versatility of electrochemistry with the simplicity and power of the DPN method to produce nanostructures with high resolution. Electrochemical STM-based methods require that the substrates be metallic, but substrates used in EC-DPN do not have to be metallic since the control feedback of the AFM does not rely on the current between the tip and surface. Si wafers coated with native oxide provides enough conductivity for the reduction of the precursor ions. This development significantly expands the scope of DPN lithography, making it a more general nanofabrication technique that not only can be used to deliver organic molecules to surfaces but is also capable of fabricating metallic and semiconducting structures with precise control over location and geometry. Local electrochemical deposition of freestanding vertically grown platinum nanowires was demonstrated with a similar approach, electrochemical fountain pen nanofabrication (ECFPN) [109]. The EC-FPN exploits the meniscus formed between an electrolyte-filled nanopipette ('the fountain pen') and a conductive substrate to serve as a confined electrochemical cell for reducing and depositing metal ions. Freestanding Pt nanowires were continuously grown off the substrate by moving the nanopipette away from the substrate while maintaining a stable meniscus between the nanopipette and the nanowire growth front. High quality and high aspect-ratio polycrystalline Pt nanowires with diameter of similar to 150 nm and length over 30 μm were locally grown with EC-FPN. The EC-FPN technique is shown to be an efficient and clean technique for localized fabrication of a variety of vertically grown metal nanowires and can potentially be used for fabricating freeform 3D nanostructures.

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3D Electrochemical Nanoconstruction In combination with lithographic patterns, electrochemistry has taken a key position in products and manufacturing processes of microtechnology, which has established a multibillion dollar market with applications in information, entertainment, medical, automotive, telecom and many other technologies such as lab on a chip etc. [110]. Different strategies and techniques such as electrochemical etching, LIGA and ultrashort voltage pulsing, which have been used in microtechnology, were also applied to construct 3D structures in the nanoscale.

Electrochemical Etching and LIGA Technique Electrochemical etching with ultra-short voltage pulses allows to dissolve electrochemically active materials within an extremely narrow volume and to manufacture three dimentional (3D) microstructures. Micro- and nanoporous silicon can be generated by anodization of silicon wafers in hydrofluoric acid. Since etching proceeds preferably in the (100) direction of the single-crystalline silicon wafer, the pore shape is nearly straight and the depth is equal for all pores. The pores start to grow on a polished wafer in a random pattern, and their arrangement is usually defined by transferring a suitable lithographic pattern and generating a corresponding pattern of pits by alkaline etching as shows in Fig. 8. The pattern of the deep pores generated subsequently by the electrochemical etching process corresponds to the pattern of the shallow starter pits. The cross section of the pores is usually square with rounded corners and their size can be varied by changing the etching current. More complex cross sections can be generated by overlapping of pores using additional etching steps and corresponding pattern definition (e.g. by means of lithography).

Figure 8. Macroporous silicon structure generated by means of electrochemical etching of singlecrystalline silicon (Source: V. Lehmann, Siemens AG) [110]. Reproduced by the kind permission from the publisher.

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The macroporous silicon structure generated by means of electrochemical etching of single-crystalline silicon was demonstrated above. In the electrochemical set-up, hydrofluoric acid (HF), with or without ethanol and/or water is used as electrolyte and platinum is the standard cathode. The etch rate of ca. 1 μm/min is observed. Both electro-polishing and pore formation take place in the anodic regime. Depending on current density, silicon can be etched in such two different modes: pore formation and electro-polishing. In pore formation, etching proceeds vertically downwards, leaving a silicon ‘skeleton’ with up to 80% empty space, whereas in electropolishing, the whole surface is being etched. Pore formation starts at the wafer surface from a defect or an intentional initial pit. Electronic holes from the bulk silicon are transported to the surface, and they react at the defect or pit. Further etching occurs at the newly formed pore tips, because they attract more holes due to higher electric field strength, and the process leads to a uniform porous layer depth as the holes are consumed by the growing tips and other surfaces are depleted of holes. This etching mode takes place under low hole concentration and it is limited by hole diffusion, and not by mass transfer in the electrolyte cell. If hole density increases, some holes reach the surface and react there, leading to surface smoothing. This is the electro-polishing regime, in which ionic transfer from the electrolyte plays a role. Illumination contributes to hole concentration in n-type silicon (but not in p-type silicon) and the anodic etching of n-type silicon happens under illumination whilst p-type silicon etches in dark. A very wide range of pore size from 0 to 20 μm can be etched by varying electrolyte concentration, current density and illumination [111]. As a rule of thumb, pore diameter in micrometre is half the resistivity in ohm/cm: for 1 um pore, 2 ohm/cm n-type silicon is suitable. For small pores, low resistivity is needed; for large pores, high resistivity material has to be used. If pore formation starts from an unobstructed surface, a random pore array results. If initial pits are prepared by lithography and etching, pores can be arranged at will [112]. When an n-type silicon wafer is an anode in an alkaline etching solution (e.g. KOH) biased positively above passivation potential, the surface will be oxidized, which stops silicon dissolution whilst p-type silicon was etched. The n-type layer of a p/n-structure can similarly be protected. Etching of p-type silicon continues until the diode is destroyed, and n-type silicon is then passivated. The confined etchant layer technique has been applied to achieve effective three-dimensional (3D) micromachining on n-GaAs and pSi. This technique operates via an indirect electrochemical process and is a maskless, lowcost technique for microfabrication of arbitrary 3D structures in a single step [113]. It has also been presented that free-standing Si quantum wire arrays can be fabricated without the use of epitaxial deposition or lithography by electrochemical and chemical dissolution of wafers [114]. This novel approach uses electrochemical and chemical dissolution steps to define networks of isolated wires out of bulk wafers. Electrochemical methods, either alone or in combination with other techniques, have been developed for shaping materials. 3D microstructures with extremely high precision and aspects ratio can be manufactured by means of LIGA technology, which combines deep lithography, electrodeposition and moulding process steps. The acronym LIGA is derived from the German expressions for these manufacturing steps and offer high potential regarding miniaturization, freedom of design and mass production. Micro-gear system produced from a nickel iron alloy by means of LIGA was shown in Fig. 9.

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Figure 9. Micro-gear system produced from a nickel iron alloy by means of LIGA technology (Source: Micromotion, IMM) [110]. Reproduced by the kind permission from the publisher.

However, the ultra precise microstructures with extreme aspect ratio could only be generated by deep X-ray lithography. Difficulties, for example, the access to synchrotron radiation facilities have limited the commercialization of LIGA technique in mass fabrication.

Micro- and Nano-machining by Ultrashort Voltage Pulsing Technique The application of electrochemistry in micro-machining can be found in book [115]. In contrast to the conventional processes of electrochemical micromachining, where the gap between the electrodes is usually 0.1 mm and direct current is applied, a novel electrochemical micofabrication method using a gap in ‘μm’ range and very short voltage pulses of some tens of nanoseconds was developed. The short pulse confines electrochemical processes correspondingly, removal of material to a very narrow volume to enable a precise nanomachining. Ultrashort pulses can be employed to machine conducting materials with lithographic precision [116]. Resolution can be improved significantly through the use of ultrashort voltage pulses comparing to the use of conventional direct current anodization. Three dimensional complex nanostructures, lines, curved features, and arrays can be machined in substrates in single-step processing. The method is based on the application of ultrashort voltage pulses of nanosecond duration, which leads to the spatial confinement of electrochemical reactions, e.g. dissolution of material. The electrochemical dissolution rate of the material has to be intentionally varied over the workpiece surface by applying inhomogeneous current density distribution in the electrolyte and at the workpiece surface. This can be achieved by the geometric shape of the tool, locally very small tool-workpiece distances, partial insulation of the tool or workpiece, and high overall current densities etc. This situation is illustrated in Fig 10. The workpiece is preferentially etched within the gap region between the front face of the tool and the workpiece surface. This approach for local confinement of electrochemical reactions is based on the local charging of the electrochemical double layer (DL) and the resulting direct influence on the electrochemical reaction rates.

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Figure 10. Sketch of the experimental setup and principle of electrochemical micromachining with ultrashort pulses. RE and CE are abbreviates of reference and counter electrode.

The potentials of the workpiece and tool are controlled by the low-frequency bipotentiostat. The voltage pulses are supplied by the high frequency pulse generator. An ultrashort voltage pulse limits the charging of double layer capacitance to the vicinity of the tool. The current distribution between the DL is also illustrated in Fig. 10. The pulsing time constant is given by the DL capacitance multiplied by the resistance of the electrolyte along the current path. The latter factor is locally varying, depending on the local separation of the electrode surfaces [116]. Therefore, upon proper choice of the pulse duration, DL areas where the tool and workpiece electrodes are in close proximity are strongly charged by the voltage pulses, whereas at further distances the charging becomes progressively weaker. The pulse duration provides a direct control for the setting the machining accuracy. Machining precisions below 100 nm were achieved by the application of 500 ps voltage pulses [117].

Figure 11. Spiral trough with a depth of 5 mm, machined into a Ni sheet with a W tool in 0.2 M HCl (3 ns, 2 V pulses, 33 MHz repetition rate, Φworkpiece≈ -0.1 VAg/AgCl, Φtool≈ -0.3 VAg/AgCl [117]. Reproduced by the kind permission from the publisher.

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The application of ultrashort voltage pulses between a tool electrode and a workpiece in an electrochemical environment allows the three-dimensional machining of conducting materials with nanoscale precision. The principle is based on the finite time constant for DL charging, which varies linearly with the local separation between the electrodes. During nanosecond pulses, the electrochemical reactions are confined to electrode regions in close proximity. The performance of electrochemical micromachining with ultra-short voltage pulses was demonstrated in a number of experiments where microstructures were manufactured from various materials like copper, silicon and stainless steels [116]. Three dimentional structures with high aspect ratios can be achieved by using suitable microelectrodes and piezo-driven micropositioning stages. The spiral shown in Fig. 11 was manufactured by machining the Ni sheet with 3 ns pulses. Walls of similar thickness with surface roughness and radii of curvature less than 100 nm were readily machined [117].

Figure 12. Scanning electron micrographs. (a) the tool; (b) structure in Ni substrate. Experimental conditions: Usub=-0.35 V, Utool=-0.3 V, 2 ns pulse duration, 2.2 V amplitude, 1:10 pulse to pause ratio, and 0.05 M HCl electrolyte. The structure was machined 400 nm into the surface in less than 2 min [118]. Reproduced by the kind permission from the publisher.

Small tools can be used to make very small features. High aspect ratio nanometre accurate features were machined in nickel using ultrashort voltage pulse electrochemical machining [118]. Two tools (one is the rotunda tool, presented in Fig 12a and other is 2×2 array of cubes, presented in Fig. 13a) were firstly fabricated by focused ion beam (FIB) milling and then used in the machining. The potentials of the shaped tungsten tool and nickel substrate electrodes were controlled with a bipotentiosts which kept the potentials of the tool and substrate constant versus an Ag/AgCl reference electrode by applying a potential to Pt counter electrode. All experiments were conducted in an aqueous HCl solution with various concentrations. Supplementary circuitry was present to allow the additional of ultrashort (order of 1 ns) pulses to the potential of the tool electrode. The separation of the tool and the substrate workpiece electrodes was controlled with piezoactuators and the tool was fed into the workpiece with a constant feeding speed, avoiding mechanical contact between the electrodes by monitoring the dc current between tool and workpiece. Structures with 90 nm widths were made by applying 2ns voltage pulses for the parallel lines in the centre of the structure in Fig 12b. To reach a depth of 400 nm, total electrochemical machining time of 1

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min 45s was applied. Examples of patterns made with the 2×2 array of cubes tool are shown in Fig. 13b. It indicates that the feature resolution improves with decrease in pulse duration.

Figure 13. Scanning electron micrographs. (a) Tungsten tool; (b) machined Ni substrate. Experimental conditions: Usub=-1.7 V, Utool=-1.0 V, pulse duration indicated, 2 V amplitude, 1:10 pulse to pause ratio, and 0.2 M HCl electrolyte. Feature resolution and edge sharpness increased as pulse duration decreased [118]. Reproduced by the kind permission from the publisher.

Three-dimensional machining of electrochemically active materials including construction of unconventional island patterns on a surface with nanoscale resolution was realized by this method [97;119-121]. Thus, electrochemical machining can be applied to microelectro-mechanical systems (MEMS) [122], and even in the nanoelectro-mechanical systems (NEMS). Electrochemical methods can realize the nanofabrication in a selective place and make the complicated 3D nanostructures. Conducting polymers can also be made in this way. Similar to the electrochemical machining, by application of short voltage pulses to the tool electrode in the vicinity of the workpiece electrode, the electropolymerization of pyrrole can be locally confined with micrometre precision [123]. As the produced nuclei of conducting polymers will grow preferentially vertically to the surface, fibre-like morphologies were found in the local polypyrrole electrosynthesis with short voltage pulses. A polypyrrole ring with needle-like feature can be selectively nanofabricated, which was grown with 1 μs pulses in 0.2 M H2SO4/0.2M Pyrrole. To obtain such structure, a 50 μm diameter flattened cylindrical Pt wire tool was brought to a distance of 1 μm from the workpiece substrate before applying the pulses. The small distance between tool and the substrate strongly hinders the supply of monomer at the reaction site. The consequent relatively low supersaturation leads to a low density of nuclei, which grow and form isolated polypyrrole fibres. With the increase of the distance, a more compact ring-like structure can be obtained.

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Template Free Methods for Conducting Polymer Nano-architecture Although oriented carbon nanotubes, oriented metal and semiconductor nanowires have attracted wide attention, there have been few reports on oriented polymer nanostructures such as nanowires. The assembly of large arrays of oriented nanowires containing molecularly aligned conducting polymers without using a porous membrane template to support the polymer was reported recently [124]. The uniform oriented nanowires were prepared through controlled nucleation and growth during a stepwise electrochemical galvanostastic deposition process, in which a large number of nuclei were first deposited on the substrate using a large current density. After the initial nucleation, the current density was reduced stepwise in order to grow the oriented nanowires from the nucleation sites created in the first step. The usefulness of these new polymer structures is demonstrated with a chemical sensor device for H2O2, the detection of which is widely investigated for biosensors. It offers a general approach to control nucleation and and has potential for growing oriented nanostructures of other materials. Fig. 14 demonstrates the steps to electrochemically fabricate arrays of oriented conducting polyaniline (PANI) nanowires by well controlled nucleation and growth without templates. After initiating the nucleation of the conducting polymer at high current densities the current density is reduced to avoid formation of further nuclei. The existing nuclei grew preferentially vertically to the surface. A typical procedure involves electrochemical deposition in an aniline-containing electrolyte solution, by using the substrate as the working electrode. This process involves: 0.08 mAcm-2 for 0.5 h, followed by 0.04 mAcm-2 for 3 h, which was then followed by another 3 h at 0.02 mAcm-2. The stepwise growth produced uniform, oriented nanowires on a variety of flat and rough surfaces as Fig. 15 shows.

Figure 14. Schematic drawing of the steps for growing oriented polymer nanowires. a) Schematics of the reactions in the electrochemical cell. b) Schematics of the nucleation and growth.[124]. Reproduced by the kind permission from the publisher.

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Figure 15. SEM micrographs of oriented polyaniline on Pt. a) Low magnification face-on. b) High magnification face-on. c) Tilted view, low magnification. d) Tilted view, high magnification. The insert in Figure 3 a is the image of the oriented nanowires on Si substrate[124]. Reproduced by the kind permission from the publisher.

The PANI films prepared by the above method appear to be fairly uniform with the diameters of the tips ranging from 50 nm to 70 nm. The direct electrochemical synthesis of large arrays of uniform and oriented nanowires of conducting polymers with a diameter much smaller than 100 nm, on a variety of substrates (Pt, Si, Au, carbon, silica), without using a supporting template was thoroughly studied [125]. Ordered PANI nanowires tailored by such stepwise electrochemical depositions showed remarkably enhanced capacitance [126]. The superior capacitive behaviors of PANI nanowires show great potential in the application of supercapacitors and rechargeable batteries. Conducting polymer nanowires are also promising one-dimensional nanostructured materials for application in nanoelectronic devices and sensors [127;128] due to their light weights, large surface areas, chemical specificies, easy processing with low costs and adjustable transport properties. The electrochemical growth of nanowire devices using e-beam-patterned electrolyte channels potentially enables the controlled fabrication of individually addressable sensor arrays [128]. This approach should be highly efficient and scalable, while meeting the current requirements for nanoelectronics technologies, i.e., an integration of bottom up production methods (electropolymerization of the nanoframeworks) and top-down fabrication (lithographic fabrication of Pt electrodes in array). The same galvanostastic template free, site specific electrochemical method was developed to precisely fabricate individual and addressable conducting polymer nanowires on electrode junctions in a parallel-oriented array [129]. Electrochemical polymerization at low and constant current levels was used to fabricate 10 nano-framework-electrode junctions simultaneously with uniform diameter (ca. 40 to 80 nm) PANI nanowires interwined to nanoframeworks. Electropolymerization was carried in an aqueous solution containing 0.5 M

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aniline and 1.0 M HCl. Firstly, a constant current (50 nA) was applied for ca. 30 min to introduce the PANI nuclei onto the Pt working junction electrodes. The current was then reduced to 25 nA for 180 min. The PANI nanoframeworks begin to propagate from the working junction electrodes to the other set of the junction electrodes. In the last step, the current was decreased to 12 nA for 180 min, and the 10 polyaniline nano-frameworkelectrode junctions were obtained simultaneously with each PANI nanoframe work positioned precisely within the 2 μm gap between its electrodes. All these nanoframe work electrode junctions can be covered by such conducting polymer wires simultaneously and sitespecifically in a parallel fashion. This device can be used as miniaturated resistive sensors for real-time detection of NH3 and HCl gases. Such two-terminal devices can be easily converted to three-terminal transistors by simply immersing it into an electrolyte solution along with a gate electrode. Electrolyte-gated transistors based on conducting polymer nanowires junction arrays was developed and the field induced modulation can be applied for signal amplification to enhance the device performance [130]. Conducting polymer nanowires including PANI (ca. 50-80 nm), PPy (ca. 60-120 nm) and PEDOT(ca. 80-150 nm) were introduced to the 10 paralleled 2μm-wide gaps by the template free electropolymerization method. The preparation of electrolyte gated transistors can be completed simply immersing the conducting polymer nanowire-based two-terminal resistive device along with a gate electrode (a Pt wire or Ag/AgCl electrode) into a buffered electrolyte solutions containing NaCl. P-channel transistor characteristics at pH7 and n-type behaviour in basic media were observed for both PANI and PPy nanowires. Whereas PEDOT nanowire based device only exhibit depletion mode behaviors in neutral solutions. These open new opportunities to fabricate sensor arrays with conducting polymer nanowires to realize the ultrasensitive, realtime and parallel detection of analytes in solutions.

Template Methods The growth of thin films displaying special features like aligned pores perpendicularly to the substrate surface and nano-porous structures have attracted the attention of many research groups in the last decade and, with that aim several techniques such as ion beam lithography have been used [131]. On the other hand, a lot of bottom up techniques, particularly those in which, self-assembly processes play a relevant role in the growth mechanisms of that nanostructures have been reported. Among them, electrochemical techniques constitute one of the most used to fabricate highly ordered nanostructures to be used as templates for replicating other nanostructured materials and for growing functionalized material arrays. An overview on the nanofabrication techniques is done mainly of those related with the nanostructures fabrication based on ordered and nanoporous anodic aluminium oxide membranes (AAO), anodic titania membranes, colloidal polystyrene (PS) latex spheres. Templates from carbon nanotubes, self-assembled monolayers (SAMs) and selfassembled block copolymers etc. are also summarized. Template methods offer very important strategies for complicated nano-structures.

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Anodized Aluminum Oxide (AAO) Membranes A simple and completely nonlithographic preparation technique for free-standing nanostructured films with a close-packed hexagonal array of nanoembossments has been developed by porous anodic aluminium oxide (AAO) membrane templates with different pore diameters. They have been largely used to construct different free-standing inorganic and organic nanowires [132-134], nanotubes [135] and ordered arrays of nanoparticles [136] since invented. Aluminium anodization provides a simple and inexpensive way to obtain nanoporous templates with uniform and controllable pore diameters and periods over a wide range. The usual electrochemical method for producing the AAO film is the anodization of high purity Al plates at constant voltage (e.g. anodized at 22 V from an Al foil and detached by the reverse-bias method) [137]. Membranes with several different pore sizes can be made, for example, in the following electrolytes: Aqueous solutions of H2SO4 at 10-20 V for pores ~1025 nm, H2C2O4 at 40-80 V for pores ~40-100 nm, and H3PO4 at 100-140 V for pores ~100170 nm. The pore diameter is linearly related to the anodizing voltage (1.2 nm/V). A voltage reduction was done to thin the barrier layer that inhibits anodic current during electrodeposition [138]. Other attempts have been made to create nano-porous symmetries other than hexagonal packing [139]. Recently, a novel AAO membrane with a six-membered ring symmetry co-existing with the usual hexagonal structure has been fabricated by constant current anodization [140].The pore sizes of this structure can be tailored by changing the processing conditions. Ordered arrays of nanodots with novel structure have been fabricated by this AAO template. In the final stage, the porous alumina substrate can be removed by etching in KOH. Moreover, one of the interesting possibilities afforded by the anodization process is that the anodization can take place on arbitrary surfaces, such as curved surfaces. Unique features including cessation, bending, and branching of pore channels are observed when fabricating AAO templates on curved surfaces [141]. The new structures may open new opportunities in optical, electronic and electrochemical applications. Many strategies have been ingeniously implemented to fabricate complicated nanostructures based on the AAO templates. For example, hexagonally ordered Ni nanocones have been fabricated using an a porous AAO template where the pores are of a cone shape [142].The conical AAO film was found to exhibit hexagonal order with a period of 100 nm. The Ni nanocones and the surface morphology of the nano-conical film exhibit the same periodic structure of the template as shown in Fig. 16. The hexagonally ordered Ni nanocones and nano-conical film were produced using anodization and metal plating techniques. The conical AAO template was produced using a process of repeated applications of the anodization and pore-widening steps, applying the two steps alternately. The Ni nanocones were produced by electroless Ni deposition onto the conical AAO template, with the pores filled with Ni particles. The resulting Ni nanocones exhibit the same ordered structure. The Ni nano-conical film was produced by detaching the deposited Ni layer, with the surface morphology also a hexagonally ordered array with a period of 100 nm. These nanostructures were produced using the wet-processes of anodization, a pore widening treatment, the pulsed deposition of Pd particles and electrochemical Ni deposition. Clearly, this complete process can produce good quality results. A double-templating approach using simple electrochemical methods to create aligned arrays of nanotubes on substrates was also introduced [143].

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Figure 16. (A) Surface normal, (B) surface angled and (C) cross-section view of the hexagonally ordered conical Ni film [142]. Reproduced by the kind permission from the publisher.

Figure 17. Schematic of method to fabricate nanotube arrays on substrates. Nanorods are electrodeposited into nanoporous anodic alumina template films, the alumina is removed, and then the nanorods are used as secondary templates for nanotube electrodeposition [143]. Reproduced by the kind permission from the publisher.

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The method used to fabricate nanotube arrays is shown schematically in Fig. 17. Initially, nanorod arrays are fabricated by electrodeposition into AAO templates. By varying the anodization conditions, the pore diameter, spacing, and height can be tuned, and the pore ordering can also be controlled. Nickel nanorods are then electrodeposited into the pores of the alumina. After deposition, the exposed ends of the nanorods are modified by anodization in a dilute KOH solution. Because the anodization is performed when the alumina is still in place, only the top ends of the nanorods are anodized. Then the alumina template is removed by a selective chemical etch, leaving an array of nickel nanorods with anodized tips. This nanorod array is then used for electrodeposition of gold nanotube arrays. The nanotube material deposits uniformly across the entire surface of the nanorod arrays, except at the anodized tips of the nanorods. Finally, the nickel nanorod array template is selectively removed, resulting in an array of open-ended nanotubes on the substrate. This approach allows both fabricate and organize nanostructures over large areas on substrates in the same process. This method is demonstrated to prepare arrays of electrodeposited, open-ended nanotubes aligned vertically on substrates. The nanotube inner diameter, spacing, and height are determined by the nanorod dimensions, which in turn correspond to the alumina film characteristics. The nanotube morphology and wall thickness can be controlled by the electrodeposition parameters. Utilizing electrochemically prepared textured aluminum sheets as a replication master in conjunction with electrochemical deposition of metals revealed a highly facile and economical way for the production of periodic metallic nanostructures in a large area with high fidelity in pattern transfer as well as with a good degree of flexibility in materials [144]. The growth of metal nanowires using AAO membranes as hard templates has been reviewed [145]. These kinds of metal nanowires can be applied to superconductivity, optical spectroscopy, sensing, and catalytic conversion, and energy harvesting. The AAO template method provides access to arrays of single-crystal metal nanowires and to quasi-onedimensional metal nanostructures with controlled compositional variation along their length. Important semiconductors such as ZnO and TiO2 with nanostructures can also be manufactured in such template or related anodization techniques.

Zinc Oxide (ZnO) ZnO exhibits many unusual properties including uniaxial piezoelectric response and ntype semiconductor characteristics. Such properties can be used in applications such as fieldemission materials [146], light emitting diodes (LEDs) [147], solar cells [148] and gas sensors [149]. Electrodeposition of ZnO films has been reported by several groups and used in fabrication of oriented nanowire and nanorod arrays [150-155]. ZnO films can be electrodeposited cathodically in aqueous chloride solutions using dissolved oxygen as a precursor. The deposition reaction in the electrolyte is: Zn 2+ + 1 O 2 + 2 e − → ZnO . The 2

deposition mechanism was analyzed in terms of electrochemical induced surface precipitation due to an increase of local pH resulting from the oxygen reduction reaction. A dramatic effect of temperature on the formation of ZnO was observed. When the temperature increases (Ttransition≈ 50˚C), a transition between amorphous insulating zinc oxyhydroxide (ZnOx(OH)y) to well-crystallized and conducting ZnO happens [156]. The dimensions and the deposition

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rate of electrodeposited ZnO nanowire arrays can be controlled by changing the chemical nature of the anions in solution. A significant variation of the diameter (65 to 110 nm) and length (1.0 to 3.4 μm) of the nanowires can be obtained by changing only the nature of the anions in solution [157]. Recently, large-scale, single-crystalline ZnO nanotube arrays were directly fabricated onto F-doped SnO2 (TCO) glass substrate via an electrochemical deposition method from an aqueous solution [158]. The tubes had a preferential orientation along the [0001] direction and hexagon-shaped cross sections. The novel nanostructure could be easily fabricated without a prepared layer of seeds on the substrate. The surface condition of substrate material and the experimental conditions played a key role in the nanotube formation. The growth of ZnO arrays and the deposition reaction take place by applying a cathodic potential (-0.7 V vs. SCE) to the TCO substrate at 80˚C. The dissolved oxygen serves as the oxygen source for the growth of ZnO arrays in the electrolyte. SEM image of the result ZnO nanotubes are shown in Fig. 18. The electrochemically deposited single crystalline ZnO nanowires can be applied in LEDs [154;155].The ZnO nanowire films can be embedded in an insulating spin-coated polystyrene layer. The spin-coating parameters are carefully finetuned to completely fill out the space between the ZnO nanowires and produce only a very thin coverage of the nanowire tips. The polystyrene layer thickness at the tips can be further reduced by the plasma etching treatment to make the n-type ZnO tip junctions outside. A top contact consisting of a thin ptype poly(3,4-ethylene dioxythiophene): poly(styrenesulfonate) (PEDOT:PSS) layer and an evaporated Au film are provided to serve as the hole injection anode in the LED. A flexible LED can be realized when electrochemical depositing the ZnO nanowire on flexible transparent polymer substrates (e.g. polyethylene terephtalate, PET) coated with indium-tinoxide (ITO) [147]. The infrastructure of such flexible LED is illustrated in Fig. 19.

Figure 18. SEM image of the ZnO nanotube film [158]. Reproduced by the kind permission from the publisher.

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Figure 19. Design scheme for a flexible LED structure consisting of vertically oriented single crystalline nanowires grown electrochemically on a polymeric ITO-coated substrate. The top contact consists of p-type polymer (PEDOT:PSS) and an evaporated Au layer. Light is emitted through the transparent polymer [147]. Reproduced by the kind permission from the publisher.

This device exhibits electroluminescence over most of the visible spectrum at moderate forward bias.

Titanium Dioxide (TiO2) Another semiconductor which has similar bandgap with ZnO, titanium dioxide (TiO2) thin films have been widely exploited in many applications such as microelectronics [159], highly efficient catalysts [160], microorganism photolysis [161], antifogging and selfcleaning coatings [162], biosensors [163], gate oxides in metal-oxide-semiconductor field effect transistors (MOSFET) [164] and more recently in dye-sensitized solar cells (DSSCs) [165]. The geometry of anodic TiO2 nanotubes can be controlled over a wide range by the applied potential in H3PO4/HF aqueous electrolytes in contrast to other electrolytes. It was found that for potentials between 1 and 25 V, tubes could be grown with any desired diameter ranging from 15 to 120 nm combined with tube length from 20 nm to 1 μm. The diameter and the length depend linearly on the voltage [166]. Aqueous HCl electrolyte can be used as an alternative to fluoride containing electrolytes to obtain the vertically oriented TiO2 nanotube arrays by anodization of titanium thin films [167]. Nanotube arrays upto 300 nm in length, 15 nm inner pore diameter and 10 nm wall thickness can be made using 3 M HCl aqueous electrolyte for anodization potentials between 10 and 13 V. A highly ordered TiO2 nanotube array with a unique surface morphology can be fabricated by electrochemical anodization of titanium in an organic electrolyte (e.g. 1:1 mixture of DMSO and ethanol) containing 4% HF. The TiO2 nanotube arrays with improved

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photochemical response can be obtained using electrochemical anodization of titanium in fluorinated organic electrolytes by optimizing etching time, applied potential, solvents and the HF concentration [168]. Using an electrochemical approach in organic electrolytes, the growth of more than 250 μm thick self-organized TiO2 nanotube layers is possible [169]. Combining the electrochemical parameters in an optimal way, ordered TiO2 nanotube layers with a length of over 250 μm have been obtained at 120 V with 0.2 M HF. The tubes can grow as a hexagonal close packed pore array. Although the TiO2 nanotubes fabricated in organic solution have the longest length and the largest surface area, its conductivity may be lower than the one synthesized in aqueous solutions [163]. Crucial parameters that decide on the dimensions are the fluoride ion concentration, the voltage and the anodization time. The different length of TiO2 nanotubes synthesized by electrochemical anodization can be controlled, and is shown in Fig. 20.

Figure 20. SEM images of TiO2 nanotubes prepared in (A) 0.1M HF acid solution at 20 V (B) 1.0M NaHSO4 containing 0.1M KF at 20 V and (C) ethylene glycol containing 0.25% NH4F at 60 V for 1 h [163]. Reproduced by the kind permission from the publisher.

Highly ordered transparent TiO2 nanotube arrays produced by electrochemical anodization have been used in DSSCs. It suggests superior electron transport in nanotubular TiO2 based DSSCs [170]. Remarkable photoconversion efficiencies were expected to be obtained with increase in the length of the nanotube arrays. Carboxylated polythiophene derivatives can be selfassembled onto the TiO2 arrays produced by anodizing titanium foils in

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ethylene glycol based electrolytes [171]. Such self-assembled hybrid polymer-TiO2 nanotube array heterojunction solar cells can yield power efficiency of 2.1% under AM 1.5 without dyes. It was found that the formation mechanism of TiO2 nanotubes is similar to the porous alumina case under high electrical field. TiO2 nanotube arrays can be fabricated by anodic oxidation of titanium foil in different electrolytes. The produced TiO2 nanotube arrays possess large surface area and good uniformity and are ready for enzyme immobilization [172-174], which can be used as biosensors. Furthermore, different length of TiO2 nanotube arrays fabricated by anodic oxidation in different electrolytes were studied for their sensitivities to hydrogen peroxide after co-immobilized horseradish peroxidase (HRP) and thionine chloride. The nanotube arrays fabricated in potassium fluoride solution has the best sensitivity to H2O2 with a detection range from 10-5 to 3×10-3 M [163]. With the use of the template of an AAO, TiO2 nanowires can be obtained by cathodic electrodeposition [175;176] where the metallic ions are attracted to the AAO cathode electrode and reduced to metallic form. For example, in a typical process, the electrodeposition is carried out in 0.2 M TiCl3 solution with pH=2 with a pulsed electrodeposition approach, and titanium and/or its compound are deposited into the pores of the AAO. By heating the above deposited template at 500 °C for 4 hours and removing the template, pure anatase TiO2 nanowires can be obtained. Highly ordered TiO2 single crystalline (pure anatase) nanowire arrays can also be fabricated within the pores of anodic aluminum oxide (AAO) template by a cathodically induced sol-gel method [177]. During this electrochemically induced sol-gel process, both the formation of sol particles and the gelation process take place in the AAO pores. Therefore, TiO2 nanowires with very small diameters (less than 20 nm or even smaller) can be obtained by this technique. In addition, the length of the nanowires can be well controlled by varying the deposition time and potential of the working electrode.

Electrochemical Fabrication of Soft Matters in Nanoscale Nanofiber, nanospheres and other nanoscale of soft matters such as conducting polymers have been fabricated in traditional chemistry way and/or via selfassembly [178-187]. Basically, 1D conducting polymer nanostructures can also be synthesized chemically or electrochemically by using “hard” and “soft” template methods. Obviously, the hard-template method (e.g. AAO) is an effective and straightforward route for fabricating conducting polymer nanostructures with diameters determined by the diameter of the pores in the template. Controlled conducting poly(aniline) nanotubes and nanofibers have been fabricated in the AAO templates and find promising applications in lithium/poly(aniline) rechargeable batteries [188]. Nanotubes and nanowires of conducting polymers including polypyrrole (PPy), polyaniline (PANI) and poly(3,4-ethylenedioxythiophene) (PEDOT), can be synthesized by electrochemical methods using the AAO templates [189]. By changing the doping level, dopant and template-dissolving solvents, the electrical and optical properties of the nanotubes and nanowires can be controlled. The diameter of the conducting polymer nanotubes and nanowires are in the range from 100 to 200 nm, depending on the diameter if the nanoporous template used. It was found that the polymerization was initiated from the wall-side of the AAO template. The synthesized nanotubes have an open end at the top with the filled end at the bottom. As polymerization time increased, the nanotubes will be filled and nanowires will be formed with the length increased. For example, PPy nanotube can be

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synthesized by applying current of 2-3 mA for 1 min. When the time is increased to 15-40 mins, PPy nanowires will be produced. Conducting polymer nanotube and nanowires prepared by this electrochemical method using AAO templates can be applied in field emitting applications [190;191]. Fig. 21 below shows an uniform of polyaniline nanowires produced at constant potential at 1.0 V for 10 min through the AAO template [192]. PPy nanotubes and nanowires can be also electrochemically synthesized through nanoporous AAO template in ionic liquids[193]. The electrolyte consisted of pyrrole monomer, solvent, and ionic liquid dopant such as 1-butyl-3-methyl imidazolium tetrafluoroborate (BMIMBF4) or 1-butyl-3-methyl imidazolium hexafluorophosphate (BMIMPF6), which is stable in air and moisture. The length and diameter of PPy nanotubes and nanowires were determined by the synthetic conditions such as polymerization time, current, and dopant. The formation of nanotube and nanowire of PPy sample was confirmed by using field emission scanning electron microscope and transmission electron microscope. Formation of PANI nanotubules in room temperature ionic liquids by means of electrochemical polymerization without any template has also been reported [194]. PANI nanotubules were synthesized electrochemically on a modified ITO glass in BMIMPF6 containing 1M trifluoroacetic acid. Tubular structures of PANI with the diameter of ca.120 nm were shown by scanning electron microscopy.

Figure 21. SEM of electrochemically synthesized (a) PANI nanowires and (b) composite nanowires in AAO membranes (etched with NaOH) [192]. Reproduced by the kind permission from the publisher.

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Variable inorganic nanoparticles of different sizes can be combined with conducting polymers, giving rise to a novel composite material with interesting physicochemical properties and possibilities for important applications. Electrochemical methods have proved to be effective in incorporating metal nanoparticles in either pre-deposited polymers or in growing polymer films. Metals can be electrodeposited at conducting polymer electrodes [195]. Nanocables consisting of Ag nanowires sheathed by polyaniline were fabricated in porous AAO template [192]. Silver/polyaniline (Ag/PANI) nanowires were prepared by simultaneous oxidative electropolymerization of aniline and reduction of Ag+ in porous AAO from an acidic electrolyte containing Ag+ and aniline. One-step electrochemical fabrication of devices based on such inorganic-organic hybrid materials offers a new strategy to make electronic devices. Electrochemical fabrication of non-volatile memory device based on polyaniline and gold particles was reported [196]. Au nanoparticles are synthesized and embedded into the PANI polymer matrix simultaneously during the electropolymerization. The impedance states of the PANI:Au composite films are nonvolatile in nature and can be read and switched several times with minimal degradation in air. Such electrochemical fabrication method can produce multi-stable nonvolatile memory device in one step, which simplifies fabrication of the memory device significantly.

Carbon Nanotube Templates Carbon nanotubes (CNTs) can be used as templates for soft matters such as conducting polymers. Composite films of CNTs with PANI, PPy or PEDOT were prepared via electrochemical co-deposition from solutions containing acid treated CNTs and the corresponding monomer. The capacitance of such composite was studied [197].

Figure 22. Continued on next page.

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Figure 22. SEM image of (a) modified ITO glass surface (MITO); (b) purified SWNTs; (c) SWNT modified with PANI (electropolymerization for 50 cycles); (d) SWNT modified with PANI (electropolymerization for 300 cycles); (e) bare ITO glass surface [198]. Reproduced by the kind permission from the publisher.

Electrochemical functionalization of single walled carbon nanotubes (SWNTs) with PANI has also been done in ionic liquids [198]. SWNTs are covalently functionalized during the electropolymerization of aniline in ionic liquids. This methodology provides a novel way by which large amount of SWNTs (15 mg/ml) can be modified by aniline electrochemically. Fig. 22 shows the processes of coating CNTs with PANI by increasing the scan cycles during cyclic voltammetry from 50 cycles (Fig. 22c) to 300 cycles (Fig. 22d). Electrochemical synthesis of PPy/CNT nanoscale composites using well aligned carbon nanotube arrays was reported [199]. The thickness of the PPy film can be easily controlled by the value of the film-formation charge. Aligned coaxial nanowires of carbon nanotubes can be sheathed with conducting polymers shown in the following figures (Fig. 23 and Fig. 24) [200]. In addition, the aligned MWNT electrode arrays can be given additional robustness by electrodepositing conducting polymer around the tubes and then employing them as enzyme sensors. Nanostructuring electrodes with carbon nanotubes for its sensing applications have been reviewed [201]. Glucose oxidase, the most popular oxidase enzyme can be immobilized onto CNTs by electrodeposition within conducting polymers [202].

Figue 23. Continued on next page.

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Figure 23. SEM images of a) aligned nanotunes after transfer onto a gold foil (a small piece of the assynthesized aligned nanotue film is inlded at the bottom-left corner to show the amorphous carbon layeras well) and b) the CP-NT coaxial nanowires produced by cyclic voltammetry (25 mV/s) on the aligned carbon nanotube electrode in an aqueous solution of 0.1M NaClO4 containing 0.1 M pyrrole [200]. Reproduced by the kind permission from the publisher.

Figure 24. TEM images of the CP-CNT coaxal nanowire formed from the cyclic voltammetry method. The images are in the tip region (left) and on the wall (right) [200] Reproduced by the kind permission from the publisher.

Colloidal Polystyrene (PS) Latex Templates In nanoelectrodeposition, the aim is to place only a single layer or more of coverage on a surface in a very controlled way. Colloidal crystal such as polystyrene (PS) latex templates have been widely used to synthesise highly ordered macroporous ceramics [203], metals [204] and polymers [205;206] where a particular interest has been in the optical, magnetic and photonic band gap properties of the resultant structures. Studies of the magnetic properties of the macroporous films show a large coercivity enhancement in comparison to

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the corresponding plain films and it was found that the coercive field gradually increases as the diameter of the spherical voids decreases for films of a constant thickness [207]. PS latex nanospheres can be self-assembled on hydrophobic surfaces such as unoxidised silicon or gold and used as templates [207;208]. Two or three dimensional highly ordered macroporous cobalt, iron, nickel, gold, platinum and palladium films containing close packed arrays of spherical holes of uniform size (an inverse opal structure) can be prepared by such simple and versatile template technique [207-209]. The films were prepared by electrochemical reduction of metal cations (e.g. gold ions) dissolved in aqueous solution within the interstitial spaces of pre-assembled colloidal templates assembled on gold surfaces. The templates were assembled from colloidal polystyrene latex spheres assembled onto gold electrode surfaces from aqueous solution by slow evaporation. The aqueous layer was allowed to evaporate naturally so that the PS nanospheres assemble by capillary force on the substrate surface. Following the electrochemical deposition of the metal films, the polystyrene templates were removed by dissolution in toluene. The resulting films will show well-formed two or three dimensional porous structures consisting of interconnected hexagonal close packed arrays of spherical voids. The diameter of the spherical voids within the structures can be varied by changing the diameter of the PS latex spheres used to form the template. The thickness of the films is controlled by varying the charge passed in their deposition.

Figure 25. SEM images of regions of macroporous gold films grown with a thickness gradient by electrochemical deposition through templates assembled from either 750 nm-diameter polystyrene spheres [209]. Φ is the diameter of the sphere on top layer. The three dotted circle lines represent the spheres beneath. Reproduced by the kind permission from the publisher.

Electrochemical deposition is ideal for the production of thin supported layers for applications such as photonic mirrors, because the surface of the electrochemically deposited

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film can be very uniform. Electrochemical deposition occurs from the electrode surface out through the overlying template, the first layer of templated material, deposited out to a thickness comparable with the diameter of the template spheres used, has a different structure from subsequent layers. The subsequent growth of the film by electrodeposition out through the template leads to a modulation of the surface topography of the film in a regular manner that will depend on the precise choice of deposition bath and deposition conditions. This is clearly shown in the following figure if more than one layer of PS is used in templates. The electrodeposition was performed at a potential of -0.90 VSCE. The SEM image shows that the spherical voids left in the gold films after the removal of the PS spheres are arranged in well-ordered, single domain, close-packed structures. Fig. 25 gives an image for a macroporous gold film prepared through a template of 750 nm diameter spheres in a region where more than one layer of PS is selfassembled. Within each hemispherical void in the top layer there are again three smaller dark circles (diameter ca.100 nm). These correspond to the interconnections to the three spherical voids in the layer below (marked as dotted circles in Fig. 25) that are left around the regions where the original polystyrene spheres in the two layers were in contact. Nanostructured macroporous semiconductors such as PbO2 can also be electropolymerized by such template methods [211]. Other nanostructured patterns such as nanodots can also be synthesized by this template method. Highly ordered magnetic nanoscale dot arrays of Ni can be fabricated from double-templated electrodeposition [212;213]. Patterns of ordered arrays of spheres with controlled spacing can be electrochemical deposited by two steps. The double templates were firstly prepared by selfassembly of PS latex spheres on a gold coated glass substrate. This primary template was used for the electrodeposition of the conducting polymer resulting in a macroporous polymer template. After the deposition of PPy, the PS spheres were dissolved in toluene leaving a secondary polymer template. The PPy was then converted into an insulator either by over-oxidation or by undoping at a sufficiently negative potential. This insulating structure was used as the template for electrochemical deposition of magnetic material such as Ni. Electrodeposition of magnetic materials gradually fills the spherical cavities of the polypyrrole template. Ordered arrays of Ni dots with quasi-spherical geometry can be fabricated by this way and the diameter of the dots can be set from 20 nm [213]. Ordered 3D arrays of polyaniline (PANI) inverse opals can also fabricated via electrochemical methods by using colloidal crystals of polystyrene beads as sacrificial templates as shown in Fig. 26 [214].

Figure 26. Schematic illustration of the procedure used for fabricating PANI inverse opals. Reproduced by the kind permission from the publisher.

Fig. 27 shows the PANI inverse opals prepared by cyclic voltammetry in 3D PS colloidal crystal template. Compared with films obtained by chemical synthesis, the inverse opaline samples obtained by electrochemistry had a much higher structural quality.

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Figure 27. SEM images of PANI inverse opals prepared via cyclic voltammetry, at low (A) and higher (B) magnification. scan rate 20 mV/s, 10 scan cycles [214]. Reproduced by the kind permission from the publisher.

PANI inverse opals were also prepared by a galvanostatic method at a current density 0.05 mA/cm2. By adjusting the polymerization time and applied current, this method allows for the exact control over the structure formation and film thickness of the obtained PANI inverse opline films. To explore potential biosensing applications, PANI composite inverse opals were fabricated by modifying the structure with different dopants, such as poly(acrylic acid) (PAA) and poly(styrene sulfonate) (PSS). It was found that these dopants had a significant effect on the structure and the mechanical stabilities of the obtained opaline films. With selection of suitable dopants, PANI composite inverse opals could be fabricated with very high quality. The obtained films remained electroactive in buffer solutions of neutral pH. Together with their huge surface area, they would be ideal candidates for biosensing applications. Such macroporous structures were used as electrocatalysts for the oxidation of reduced β-nicotinamide adenine dinucleotide (NADH). It showed that the electrocatalytic efficiency of the inverse opline film was much higher compared with that of an unpatterned film. Using such templates is a clear example to show the significant advantages of electrochemical deposition methods. It produces a high density deposited material and no shrinkage of the material takes place when the template is removed. Also it can be used to prepare a wide range of materials and allows fine control over the thickness of the resulting macroporous film through control of the total charge passed to deposit the film.

Electrochemistry and Self-assembled Monolayers (SAMs) Electron transfer cannot occur in blocking films. However, for very thin self assembled monolayers (SAMs) of alkane thiols or oxide films, electrons can tunnel through the film and cause Faradaic reactions. Monolayer of alkane thiols can form spontaneously ordered adlayers on substrate like Au(111) due to a strong interaction between the sulphur of the thiol and the gold substrate. These SAMs have received tremendous attention in recent years [215;216]. Electrochemical deposition onto self-assembled monolayers gives new insights into nanofabrication [217]. Pattern transfer with high resolution is a frontier topic in the emerging

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field of nanotechnologies. Electrochemical molding is a possible route for nanopatterning metal, alloys and oxide surfaces with high resolution in a simple and inexpensive way. This method involves electrodeposition onto a conducting master covered by a alkanethiolate SAMs. This molecular film enables direct surface-relief pattern transfer from the conducting master to the inner face of the electrodeposit, and also allows an easy release of the electrodeposited film due their excellent anti-adherent properties. Replicas of the original conductive master can be also obtained by a simple two-step procedure. SAM quality and stability under electrodeposition conditions combined with the formation of smooth electrodeposits are crucial to obtain high-quality pattern transfer with sub-50 nm resolution. Fig. 28 demonstrated the steps involved in metal electrodeposition on SAMs covered substrates. Further in-depth investigations are required for improving SAM quality reducing the defect size and density, and accordingly increasing the lateral resolution of the method.

Figure 28. Scheme showing the steps involved in metal electrodeposition on SAMs covered substrates. a) Defective sites at SAMs (the molecules are indicated in black); b) nucleation and growth of the metal (grey) within the SAM; c) three dimensional growth outside the SAM; d) formation of a continuous metallic film on the SAM [217]. Reproduced by the kind permission from the publisher.

Electrochemical work comprised mainly copper deposition onto alkanethiol SAMs. Under UPD, Cu would go down on the bare substrate only [218]. The chain length of the alkane thiol on gold has an influence on the deposition over potential [219]. The long hain alkane thio (C18) on Au (111) is demonstrated to have a high blocking power [220]. However, it appears to remain one of the challenges for the near future to find the experimental conditions, under which metal can be deposited on top of a SAM, preferably as a 2D overlayer.

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On the other hand, electrochemistry can induce self-assembly of surface-templated (organo)silica thin films on various conducting supports, with mesopore channels oriented perpendicular to the solid surface over wide areas [221]. This method is intrinsically simple, very fast and does not require any pre-treatment of the support. It consists of combining the electrochemically driven self-assembly of surfactants at solid-liquid interfaces and an electroassisted generation methods to produce sol-gel films. The method of electrochemically assisted self-assembly of mesoporous silica thin films involves the application of a suitable cathodic potential to an electrode immersed in a surfactant-containing hydrolysed sol solution to generate the hydroxyl ions that are necessary to catalyse polycondensation of the precursors and self-assembly of hexagonally packed one-dimentional channels that grow perpendicularly to the electrode surface. The method opens the way to electrochemically driven nanolithography for designing complex patterns of widely accessible meso-structured materials.

Other Template Methods Molecular templates such as modified cyclodextrin have been used for electrochemical nanofabrication. Conducting polymer nanowires and nanorings can be electrochemically synthesized using the molecular templates (thiolated cyclodextrins) on gold [222]. The strategy is to apply electrochemical growth on gold electrodes modified with SAMs of wellseparated thiolated cyclodextrins in an alkanethiol forest. Thiolated aniline monomer is anchored to the surface within the cyclodextrin cavity and forms an initiation point for polymer wire growth. Nanosized conducting polymer wires several tens of nanometer thick and a few micrometers long can be synthesized electrochemically by this template. Even though the polymer wires appear to be made of numerous single strands, it was the first time that a single nanowire thread of a conducting polymer has been isolated. Block copolymers are a class of materials that selfassemble on macromolecular dimensions and have enormous potential in nanoscale patterning and nanofabrication as templates [223;224]. Some complicated nanoporous structures can be replicated by electrochemical deposition with the help of selfassembly of block-copolymers, for example, poly(4-fluorostyrene)-b-poly(D,L-lactide) [PFSPLA] [225]. By this means, very large arrays of possible materials can be manufactured. Free standing nanowires were obtained after removal of the block copolymer template by either dissolution or by UV irridation. Such mild etch method is generally useful to the nanostructures that are sensitive to more aggressive template removal processes. Furthermore, the electrochemical deposition method guaranties a conformable and stable interface between the electrode and the electrodeposited materials. PPy nanotubules have been chemically and electrochemically synthesized inside the pores of nanoporous polycarbonate (PC) particle track-etched membranes (nano-PTM) [226]. Other templates to produce the nano-structured conducting polymers such as DNA [227] and living neural tissue [228] have been reported for its potential applications in biosensors. Nanofabrication technologies are useful for developing highly sensitive, reproducible nanobiosensors. A nanometric system that is composed of well-oriented nanowell arrays can be used for highly sensitive electrochemical DNA detection, it obtained a two-orders-ofmagnitude enhancement in sensitivity [229].

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The step edge defects on single crystal surfaces can be also used as templates to form conducting polymer nanostructures. Polypyrrole nanostructures with diameters less than or equal to 10 nm have been electropolymerized using step and pit defects on highly ordered pyrolytic graphite (HOPG) as templates for electropolymerization [230]. Step defects were naturally occurring, and pits were formed via oxidation of freshly cleaved surfaces of an HOPG wafer by heating at similar to 640 °C. Underpotential deposition of approximately similar to 80 mV caused polypyrrole to form only on the step and pit edges of HOPG at and not on the basal plane. The size of these nanostructures could be controlled by limiting the pyrrole polymerization time at anodic potentials. Recent modeling results allow the morphology of the deposition to be inferred, and the wire-shaped growth is up to 30 s at constant potential, after which the growth changes morphology. Scanning tunneling microscopy data confirm this result. These polypyrrole nanostructures can be removed by sonication.

Others Nanoscale Electrochemistry Individual carbon nanotubes have been modified selectively on one end with metal using a bipolar electrochemical technique [231]. A stable suspension of nanotube was introduced in a capillary containing 10 mM HAuCl4 aqueous electrolyte, and a high electric field is applied to orientate and polarize the individual tubes. During their transport through the capillary under sufficient polarization (30 kV), each nanotube is the site of water oxidation on one end and the site of metal ion reduction on the other end with the size of the formed metal cluster being proportional to the potential drop along the nanotube. Bipolar electrochemistry occurs when an external electrical field polarizes an object that is not physically connected to the electrodes and thus generates an anodic and a cathodic area on the same object. The substrate can be any kind of material, but its conductivity must be higher than that of the surrounding medium. The induced potential difference between the two extremities of the object, and therefore the kinetics of the associated redox reactions, is directly proportional to the particle’s effective length [232]. When the capillary is exposed to a high electric field (10-30 kV), an electroosmotic flow is generated inside the capillary, transporting the CNT/AuCl4 suspension from the anodic capillary inlet towards the cathodic compartment. A 1 mm long carbon fiber was inserted into a glass capillary connected to two reservoirs filled with HAuCl4 electrolyte. The fiber was observed under the microscope during the application of different potential values. It was found that potentials higher than 40 V are needed to generate a visible metal deposit on the cathodic side of the fiber. After less than 5 min, a visible gold deposit was clearly formed on the negativel polarized end of the fiber. In the course of 1 h, the metal continued growing and its morphology, as revealed by SEM, was dominated by an agglomeration of small crystallites. The Au particle formed after 45min was illustrated in Fig 29.

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Figure 29. Capillary filled with an aqueous CNT/AuCl4 suspension dipping in the two reservoirs of a capillary lectrophoresis setup. A high electric field is applied, leading to the polarization of the individual nanotubes, thus triggering different electrochemical reactions on either end. Optical micrographs of a carbon fiber inside a glass capillary during dissymmetric gold deposition by bipolar electrochemistry. The applied voltage was 70 V for a capillary with a length of 10 cm, filled with 10 mM HAuCl4 aqueous electrolyte [231]. Reproduced by the kind permission from the publisher.

Looking to the future, this capillary assisted bipolar electrodeposition can be generalized to other types of nano-objects and also deposits of a very different nature such as other metals, semiconductors, or polymers. The approach therefore opens up the way to a whole new family of experiments leading to complex nano-objects with an increasingly sophisticated design allowing original applications.

Sonoelectrochemistry Single crystalline CdSe nanotubes have been successfully synthesized by a sonoelectrochemical route in aqueous solution at room temperature [233]. The sonoelectrochemical method is accomplished by applying an electric current pulse to nucleate and perform the electrodeposit, followed by a burst of ultrasonic energy that removes the products from the sonic probe cathode [234;235]. The growth progress suggested that the CdSe nanotubes were fabricated by sonication-induced rolling-up of CdSe nanosheets and the resulting CdSe products are in tubular structure. This method is simple, convenient and environmentally benign. The sonoelectrochemical synthesis of other nanomaterials is being investigated currently. In addition, traditional top-down nanofabrication methods such as focused ion beam (FIB), can be used to fabricate nanopore array electrodes [236]. FIB milling thus represents a simple and convenient method for fabrication of prototype nanopore electrode arrays. These electrode nano-arrays can be used in electrochemical nanofabrication for applications in sensing and fundamental electrochemical studies. Various organic and inorganic materials with nanoscale architectures can be fabricated by electrochemical nanofabrication. It is a versatile method for fabricating nanostructures with its simplicity, low-temperature processing, low equipment cost and precise control of the deposit thickness through control of the total charge passed. It is an exciting era to witness the

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emerging of nanotechnology. Electrochemistry will definitely contribute to its development independently and interdisciplinarily with other nanofabrication methods.

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In: Nanotechnology… Editors: C.J. Dixon and O.W. Curtines, pp. 51-69

ISBN: 978-1-60692-162-3 © 2010 Nova Science Publishers, Inc.

Chapter 2

FABRICATION AND APPLICATION OF NOVEL TWO-DIMENSIONAL NANOWEBS VIA ELECTROSPINNING Bin Ding*,1, Chunrong Li2, Dong Wang3 and Seimei Shiratori2 1

Modern Textile Institute, Donghua University, Shanghai 200051, China 2 SNT Co. Ltd., Kawasaki 212-0054, Japan 3 Fiber and Polymer Science, University of California, Davis, CA 95616, USA

Abstract This chapter reviews our recent progress on the novel two-dimensional nanowebs by the optimization of processing parameters during electrospinning. Using high applied voltage and low relative humidity in chamber, the by-product of micro-sized defect films can be splitted into nanowebs due to the fast phase separation of the charged droplets which flight with high moving speed in electric field from capillary tip to collector. The electrospun fibers act as a support for the “fishnet-like” nanowebs comprising interlinked one-dimensional nanowires. The average diameter of the nanowires contained in typical nanowebs is about one order of magnitude smaller than that of conventional electrospun fibers. Nanowebs together with common electrospun nanofibers can be assembled into a three-dimensional fibrous mat. So far, nylon-6, polyacrylic acid (PAA), poly(vinyl alcohol) (PVA)/SiO2 nanoparticles, and PVA/zinc acetate have been found to have the possibility forming nanowebs. The formation, morphology, and area density of the nanowebs in electrospun fibrous mats are strongly affected by the applied voltage, ambient relative humidity, kinds of solvents, solution concentration and conductivity, and distance between capillary tip to collector. The expanded applications of electrospun fibers are expected due to the formation of nanowebs, such as the nano-sized controllable filters, high efficient catalysts, catalyst supporter, and sensors. The preliminary data showing that the sensitivity of PAA nanowebs to ammonia is 2.5 times higher than that of electrospun PAA nanofibers. Additionally, PAA nanowebs show much quicker absorption speed and larger capacities than that of PAA nanofibers during the ammonia absorption test.

*

E-mail address: [email protected] (B. Ding) Tel & Fax: +86-21-62378392; Corresponding author.

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Introduction The process of electrospinning was first studied by Zeleny [1] in 1914 and patented by Formhals [2] in 1934. In the past decade, this technique regained a great deal of attention due to a surging interest in nanotechnology as continuous ultra-fine fibers or fibrous structures of various polymers with diameters in the range from several micrometers down to tens of nanometers [3-6]. The ultra-fine fibers can be easily fabricated under the driving force of an external electric field imposed on a polymer solution or melt. Electrospinning can be considered as a special case of electrospray process which uses electrostatic fields to form and accelerate liquid jet from the tip of a capillary [1,7]. The surface of a hemispherical liquid drop suspended in equilibrium at the end of a capillary will be distorted into a conical shape in the presence of an external electric field. This distortion is caused by a balancing of the repulsive force resulting from the induced charge distribution on surface of droplet with the surface tension of liquid [8]. Taylor showed that at a critical voltage, the equilibrium shape of the suspended meniscus was a cone with a semi-vertical angle of 49.3o [7]. A stable jet of liquid could be ejected and accelerated if the applied voltage exceeded this critical voltage. The jet breaks up into small droplets as a result of the longitudinal Rayleigh instability caused by surface tension in the case of low viscosity liquids. This process is known as electrospray for applications to obtain aerosols composed of sub-micron droplets with narrow distribution [9,10]. For high viscosity liquids, polymer solutions or melts, the jet does not break up, but travels as a jet to the grounded collector. The transverse instability or splaying of the jet into two or more smaller jets is observed due to the radial charge repulsion [11]. This process is termed as “electrospinning” and it produces the polymer fibers with diameter in sub-micron scale [8]. The experimental observation of micro-sized films among the electrospun fibers has been reported [8, 12, 13]. The sub-micron or micron sized droplets, which produced by the high electrically driven instability of the liquid droplet suspended at the tip of capillary, was considered to form micro-sized films among the electrospun fibers. These unexpected microsized films with uncompleted solvent evaporation led the melt of their covered dry fibers, and destroyed the uniformity of fibers. Therefore, they were called defect films. The variation in frequency of defect films formation, size and area density of defect films could be controlled by adjusting the processing parameters during electrospinning. The defect film density increases with increasing instability of the jet at the capillary tip. The one-dimensional (1D) electrospun fibers has a typical length > 100 μm and diameters in the range of 30-2000 nm [3]. The three-dimensional (3D) non-woven mats composed of electrospun fibers are considered to be have a larger surface-to-volume ratio and smaller pore size compared to commercial textiles, making them excellent candidates for applications in sensor system [14-16], filtration [17,18], tissue engineering [19-21], dye-sensitized solar cells [22, 23], super-hydrophobic surfaces [24-28], etc. Considerable recent progress in electrospinning were made that are expected to have lasting impact on the quality and scope of the applications, such as the alignment of electrospun fibers [29-31], fabrication of continuous carbon [32], polyoxometalate [33, 34], ceramic [35-37], coated [38-41], and hollow [34, 42, 43] fibers. Additionally, Yu et al. reported two major strategies to decrease the fiber diameter from the low concentration of polymer in solution and the high chargecarrying capacity of solution [44]. However, the reduction of electrospun fiber diameter is

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still a serious challenge. Electrospun fibers with average diameters below 50 nm are still difficult to be produced repeatedly and uniformly for most materials with electrospinnability until now. This chapter reviews our recent new observation of nanowebs, which widely distributed among electrospun fibers during electrospinning with optimized processing parameters. The nanowires within a nanoweb, not electrospun fibers, can be easily fabricated with diameters below 50 nm. Poly(acrylic acid) (PAA) and nylon-6 are selected as a template material for demonstration of the nanoweb formation process and a typical nanoweb-rich example, respectively. Additionally, other solution systems such as poly(vinyl alcohol) (PVA)/ZnO and PVA/SiO2 also formed the nanowebs after electrospinning. Preliminary comparative study of PAA nanofibers and nanowebs was carried out by ammonia gas absorption and detection.

Experimental Preparation of Polymer Solutions for Electrospinning The starting materials included PAA (Mw 250 000, Wako), Nylon-6 (Aldrich), PVA (Mn 66 000, Wako), zinc acetate (Wako), SiO2 nanoparticles (12 nm), ethanol (Wako), distilled water, and formic acid (Wako). Four kinds of polymer solutions were prepared as the following procedures. PAA powder was dissolved in H2O, ethanol, and their blends, respectively, with concentrations of 6, 8, and 10 wt%. Nylon-6 was dissolved in formic acid with concentration of 10, 15, 20, and 25 wt%. A 10 wt% PVA solution was prepared from PVA powder and distilled water at 80 oC with vigorous stirring. The electrospinning PVA/ZnO solution can be obtained by blending 10 g PVA solution, 1 g zinc acetate, and 1 g distilled water at room temperature under stirring for 6 h. 0.2 g SiO2 nanopartilces was blended with 10 g PVA solution (10 wt%) under vigorous stirring and ultrasonic treatment. The viscosity and conductivity of blend solutions were measured by a viscotester (6L/R, Hakke, USA) and electric conductivity meter (CM-40G, TOA•DKK Co., Japan), respectively.

Fabrication of Nanofibrous Membranes via Electrospinning The polymer solutions were loaded into plastic capillaries which immersed with a copper wire. The copper wire was connected to a high voltage power supply (FC30P4, Glassman High Voltage Inc., USA) that was capable of generating DC voltage up to 30 kV. The ambient relative humidity was used as 20, 45, and 75 %, respectively. The spinning distance was regulated in the range of 5-25 cm. Various samples were obtained by adjusting the processing parameters during electrospinning. The fibrous samples can be deposited on Al foil and quartz crystal microbalance (QCM, 10 MHz, AT-cut quartz crystal with Ag electrodes) electrodes. The resultant samples were conducted by 5 s of Os coating under vacuum using an Os coater (HPC-1S, Vacuum Device Ltd., Japan). The thickness of Os coating layer on samples was less than 2 nm. Morphological observations of samples were made by scanning electron

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microscopy (SEM) (S-4700, Hitachi Ltd., Japan). The diameters of samples were measured using image analyzer (Adobe Photoshop 7.0).

Results and Discussion Finding of Nanowebs in Electrospun Fibrous Membranes SEM images of PAA samples produced with concentration of 6 wt%, voltage of 30 kV, distance of 15 cm, and relative humidity of 20 % as a function of kinds of solvents are shown in Figure 1. As shown in Figures 1A and B, several discontinuous irregular defect films with maximum length up to 20 μm were found among the electrospun PAA fibers, which were spun from the solvent of H2O and H2O/ethanol = 3/1 (w/w). The same phenomenon for observation of defect films was reported with electrospinning of other materials [8,12,13]. The defect films were formed from the charged droplets which originated from the jet at the capillary tip with increased instability. However, the defect films were partly split into webs when the PAA spun from volatile solvent of H2O/ethanol = 1/1 and ethanol (Figures 1C and D). The formation of webs from films was considered due to the fast phase separation of polymer and solvent in charged droplet caused by solvent evaporation during the flight in high electric field [45,46]. PAA has a relatively faster phase separation with ethanol than with H2O because of the low boiling point of ethanol [47].

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Figure 1. SEM images of PAA fibers produced with PAA solution concentration of 6 wt%, voltage of 30 kV, distance of 15 cm, and relative humidity of 20 % as a function of kinds of solvents: (A) H2O; (B) H2O/ethanol = 3/1; (C) H2O/ethanol = 1/1; and (D) ethanol. (E) High magnification SEM image taken from the sample shown in D.

Figure 1E showed the high magnification SEM image of sample, which spun from ethanol. The strong bonding between nanowebs and electrospun fibers was found. The trace solvent left in the nanowebs caused by the uncompleted solvent evaporation led the bonding between nanowebs and electrospun fibers. Additionally, the nanowebs that covered on electrospun fibers still kept the morphology of nanoweb to form a porous surface on fibers indicating the formation of the nanowebs was performed during the flight before reaching collector.

Relative Humidity Effect on PAA Nanoweb Formation The ambient relative humidity in electrospinning chamber was important to affect the evaporation speed of solvent in charged droplets during the flight [48]. As the relative humidity increased to 45% and 75 %, the PAA fibers (Figures 2A and B) showed a relatively larger fiber diameter than that of fibers produced under a low relative humidity of 20 % (Figure 1D). And the adjacent fibers stuck together to form a porous film without the appearance of nanowebs.

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Figure 2. SEM images of PAA fibers formed with PAA solution concentration of 6 wt%, voltage of 30 kV, distance of 15 cm, solvent of ethanol, and relative humidity of 45% (A) and 75 % (B).

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The high relative humidity retarded the evaporation of solvents from jets which led an increased fiber diameter and the wet fibers linked together on collector.

Effect of Applied Voltage on PAA Nanoweb Formation Another processing parameter which affected the PAA nanoweb formation was the applied voltage. As shown in Figure 3, there were no nanowebs for the sample produced with the applied voltage of 20 kV. As reported by Deitzel [8], the applied voltage was one of the key parameters to initial the bead defect formation. In this study, the applied voltage was further proved to be one of the key parameters to affect the formation of nanowebs.

Figure 3. SEM image of PAA fibers formed with PAA solution concentration of 6 wt%, voltage of 20 kV, distance of 15 cm, relative humidity of 20 %, and solvent of ethanol.

Proposed Mechanism The forces acting on the charged droplet which flight with a high moving speed in electric field are shown in Figure 4A. The forces included electrostatic force, drag force, gravity, coulombic repulsion force, surface tension and viscoelastic force. The electrostatic force carried the charged droplet from capillary tip to collector. The drag force happened between the surrounding air and charged droplet with high moving speed. And the drag force was attributed to deforming the droplets into films. The coulombic repulsion force tried to expand the droplet. The surface tension and viscoelastic forces led the contraction of charged droplet [49]. The electric field could be increased on increasing the applied voltage with a constant distance. Consequently, the electrostatic and coulombic repulsion forces of charged droplet were reinforced with increasing of electric field. The increased electrostatic force further accelerated the moving of charged droplet, which led an increased drag force. The distortion and expand of charged droplet from spherical-like to spindle-like in electric field during electrospraying was reported by Grimm and Beauchamp [50]. The further expand could happened when the electric field increased to form thin films from droplets with the effect of increased coulombic repulsion and drag forces. Moreover, the increased radial charge repulsion force also has trend to expand the charged films. As a result, the deformation of charged droplet was strongly affected by the electric field.

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(A)

(B) Figure 4. (A) Forces acting on the charged droplet. (B) Schematic diagram illustrating the possible mechanism of nanoweb formation during electrospinning.

The schematic diagram illustrating the possible mechanism of nanoweb formation during electrospining was shown in Figure 4B. Due to the high viscosity of polymer solutions, the major jets ejected from the tip could be continuous (as in conventional electrospinning). We considered that the defect films or nanowebs were not formed from the charged droplets which were obtained by breaking jets (as in electrospraying). The defect films or nanowebs just could be regarded as a by-product caused by a high electric field induced instability of suspended charged droplets during electrospinning [8]. The microsized charged droplets [51] could be initialed together with the common electrospun fibers from the electrospining tip with a high instability. During the flight of charged droplet from capillary tip to collector, the charged droplet bore the comprehensive effects of the forces acting on it. On increasing the moving distance, the droplet was distorted and expanded into thin film by the coulombic repulsion and drag forces in the strong electric field. The splitting of thin film into nanoweb performed due to the fast phase separation between polymer and solvent which caused by the fast solvent evaporation under a low relative humidity. The fast phase separation led the

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spinodal or binodal types of phase morphologies within the fibers and the solvent rich regions were apparently transformed into pores [45]. As the electrospun fibers and nanowebs were formed at the same time, the 2D nanowebs stacked into 3D fibrous mats in a layer-by-layer structure as same as 1D electrospun fibers. The increasing of formation frequency and area density of nanowebs could be realized with increasing the instability of droplet at the electrospinning tip, such as increasing the applied voltage [8, 52].

PAA Solution Concentration Effect The morphologies of PAA nanowebs produced with various concentrations of PAA solution are shown in Figure 5. It could be observed that all the samples showed the uncompleted splitting of defect films. As shown in Figures 5A and B, the nanowebs produced with 6 wt% of PAA showed the largest average hole diameter and diameter deviation of nanowebs among three samples. As the PAA concentration increased to 8 wt% (Figures 5C and D), the hole diameter and shape of nanoweb became uniform and the average hole diameter reduced compared with the sample in Figures 5A and B. Moreover, as shown in Figures 5B and D, the shapes of holes in nanowebs included round, triangle, quadrangle, pentagon, and hexagon. Figures 5E and F showed the nanowebs produced with 10 wt% of PAA. The nanowebs showed almost only round shape of holes with the smallest average hole diameter among three samples. As a result, the shape, diameter, and uniformity of holes in PAA nanowebs could be adjusted by changing the solution concentrations.

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Figure 5. SEM images of PAA fibers produced with applied voltage of 30 kV, distance of 15 cm, relative humidity of 20 %, and solvent of ethanol as a function of PAA concentrations: (A) 6 wt%; (C) 8 wt%; and (E) 10 wt%. (B, D, and F) High magnification SEM images taken from the sample shown in A, C, and E, respectively.

Effect of Voltage and Relative Humidity on Nylon-6 Nanoweb Formation The example for relatively completed splitting of films into nanowebs was demonstrated using nylon-6 dissolved in formic acid. As the key parameter for the formation of nanowebs, the influence of applied voltage was investigated during fabrication of nylon-6 nanowebs. When the applied voltage was 10 kV (Figure 6A), a web comprising several ultrathin nylon-6 nanowires appeared in a small region. As the voltage increased gradually from 15 to 25 kV (Figures 6B to D), the 2D web comprising many interlinked 1D nanowires appeared. The nanowebs almost covered all the regions in SEM image when the applied voltage was 30 kV (Figure 6E). It was found that the area density of nanowebs in fabric was sharply increased with increasing the applied voltage. This result was identical to the increase of bead defect density [8] with increasing the applied voltage. (A)

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Figure 6. SEM images of nylon-6 fibers formed with nylon-6 solution concentration of 15 wt%, distance of 15 cm, and relative humidity of 20 % as a function of applied voltage: (A) 10 kV; (B) 15 kV; (C) 20 kV; (D) 25 kV; and (E) 30 kV.

Unlike the PAA, the nylon-6 nanowebs could be fabricated under a high relative humidity of 75 % (Figure 7). This was attributed to the nylon-6 has faster solidification [53] speed than that of PAA [45]. Despite the observation of nylon-6 nanowebs under the high relative humidity of 75%, the area density of nylon-6 nanowebs was largely decreased compared with the sample fabricated under the low relative humidity of 20 % (Figure 6E).

Figure 7. SEM image of nylon-6 fibers formed with nylon-6 solution concentration of 15 wt%, voltage of 30 kV, distance of 15 cm, and relative humidity of 75 %.

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Effect of Distance Figure 8 showed the influence of electrospinning distance to the morphology of nylon-6 nanowebs. As shown in Figure 8A (distance of 5 cm) and 8C (distance of 25 cm), the 2D nanowebs distributed randomly among electrospun fibers as the way that electrospun fibers stacked. The electrospun fibers acted as supporter to support the nanowebs. From the high magnification SEM image (Figure 8B), the short distance (5 cm) led a relatively larger wire diameter about 30-50 nm in nanowebs because of the uncompleted expand of the nanowebs. Meanwhile, the strong bonding between electrospun fibers and nanowebs was found due to the uncompleted solvent evaporation in this short spinning distance. As the distance increased to 25 cm (Figure 8D), the thinner wire diameter about 10-20 nm in nanowebs was found due to the relatively sufficient expand of nanowebs in this long spinning distance. As a result, the nanowire diameter in nanowebs and bonding between fibers and nanowebs were strongly affected by spinning distance.

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Figure 8. SEM images of nylon-6 fibers formed with nylon-6 solution concentration of 15 wt%, voltage of 25 kV, relative humidity of 20 %, and distance of (A) 5 and (C) 25 cm. (B and D) High magnification SEM images taken from the samples shown in A and B, respectively.

Nylon-6 Solution Concentration Effect Figure 9 showed the SEM images of nylon-6 samples produced from various concentrations. At 10 wt% (Figure 9A), the electrospun fibers have the thinnest mean fiber

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diameter of 66 nm (Figure 10A) and a few nanowebs appeared. On increasing the concentration from 10 to 20 wt% (Figures 9A to C), the fiber diameter increased gradually up to 184 nm (Figures 10A to C) as well as increased the area density of nanowebs. However, at 25 wt%, the fiber diameter was sharply increased to 1 μm and the area density of nanowebs decreased (Figures 9D and 10D).

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Figure 9. SEM images of nylon-6 fibers produced with applied voltage of 25 kV, distance of 15 cm, and relative humidity of 20 % as a function of nylon-6 concentrations: (A) 10 wt%; (B) 15wt%; (C) 20 wt%; and (D) 25 wt%. 100

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The typical nanoweb-rich example of nylon-6 (shown in Figure 9C) formed from 20 wt% of nylon-6 solution were found like fishing net to cover the electrospun fibers completely and to form strong bonding with electrospun fibers. The high magnification of same sample was shown in Figure 11A. The dense nylon-6 nanowebs with uniform pore were formed. The pore diameter of nanowebs was in the range of 10-80 nm, which was much less than that of pores among electrospun fibers. The nanowire diameter distribution of nylon-6 nanowebs was shown in Figure 11B. The major distribution region (over 80 %) of nanowire diameters was 10-20 nm with average diameter of 17 nm. The standard deviation of the wire diameters in nanowebs was 5 nm. The average diameter of nanowires was one order of magnitude smaller than that of common electrospun fibers. Therefore, the surface-to-volume ratio of fibrous mats was expected to be increased with the appearance of nanowebs [36]. (A)

Figure 11. (A) High magnification SEM images of nylon-6 nanowebs formed with solution concentration of 20 wt%, voltage of 25 kV, distance of 15 cm, and relative humidity of 20 %. (B) Histogram showing the diameter distribution of nanowires shown in A.

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Gas Absorption and Sensing Properties of PAA Nanofibers and Nanowebs The gas absorption test was carried out in a closed box (9 L) which installed a fan (Figure 12A). Each 0.25 g fibrous PAA sample was examined with 10.5 ppm ammonia gas. The remaining concentration of ammonia was determined by ammonia gas testing tube. Figure 12B presents the ammonia gas absorption by PAA fibers and nanowebs. Due to the higher surface area of PAA nanowebs, the nanowebs showed the faster absorption speed and higher absorption ability than those of nanofibers. After 55 min absorption, the concentration of ammonia decreased from 10.5 ppm to 3.5 and 1.0 ppm for nanofibers and nanowebs, respectively.

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(B) Figure 12. (A) Schematic of gas absorption system. (B) Gas absorption efficiency of PAA fibers and PAA webs.

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(B) Figure 13. (A) Schematic of a gas testing system for NH3 detection. (B) Response of sensors containing PAA nanofibers and nanowebs exposed to 1 ppm NH3 at the relative humidity of 30 %.

Nanofibrous membranes were incontinuously collected on the surface of QCM until the required frequency shift (coating load) was got. The resonance frequencies were measured by a frequency counter (Hewlett Packard 53131 A). Then, the nanofibrous membranes coated QCMs were dried at 80 oC in vacuum for 2 h to remove the trace solvent prior to the subsequent characterizations. The QCM frequency shifts caused by the deposition of PAA fiber and nanowebs were 28 000 and 22 000 Hz, respectively. In the current work, the Sauerbrey equation for used QCM can be drawn as Δf = —Δm/0.96×10-9. It means that the frequency is decreased for 1 Hz in the case of 0.96 ng of gas molecules were absorbed. The flow-type experimental setup for measuring the sensing properties of sensors is shown in Figure 13A. The sensor was installed in the chamber which kept with constant temperature (25 oC) and relative humidity (30%). The N2 was used as carrier gas. The flow rates of dry N2, wet N2, and target gases were kept constant by mass flow controllers (MFCs, Estec, SEC-

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400 MK3). During the measurement, the concentration of NH3 was regulated at 1 ppm and the absorption time was 30 min. The sensor responses to target gases were examined by measuring the resonance frequency shifts of QCM which due to the additional mass loading. The resonance frequencies were measured by the frequency counter. The data from the sensors were recorded by a personal computer. After stabilization in N2 at relative humidity of 30%, both samples showed the response the 1 ppm ammonia gas (Figure 13B). The PAA nanofibers showed the maximum 6 Hz frequency shift for 30 min ammonia absorption and returned to the original frequency after the ammonia gas desorption. For nanoweb sample, it showed the faster response speed and higher maximum frequency shift of 15 Hz to 1 ppm ammonia gas, 2.5 times higher than that of nanofibers. However, it needs more time to get the maximum frequency shift and desorb the ammonia gas. As the nanofiber sensor already showed much higher sensitivity than flat sensing films [15,16], the nanoweb structure would be the better candidate for the future highly sensitive material mode.

Formation of Nanowebs in Other Solution Systems Figure 14 showed the morphology of the electrospun pure PVA fibrous films. As a typical electrospun fibrous film, the PVA fibers were randomly oriented as a porous film with a wide fiber diameter distribution. The PVA fiber was straight with an average fiber diameter of 239 nm.

(A)

(B)

Figure 14. SEM images of pure PVA electrospun nanofibers.

The SEM images of composite PVA/ZnO fibrous membranes are shown in Figure 15. It can be observed that the composite fibers have many junctions among the fibers, showing a poor fiber uniformity compared with the pure PVA fibers. The average diameter of PVA/ZnO fibers (258 nm) was larger than that (239 nm) of pure PVA fibers due to its increased viscosity (Table 1). It is well known that the morphology and properties of electrospun nanofibers are strongly influenced by the solution properties such as viscosity and conductivity [54]. Table 1 shows the viscosity and conductivity of the electrospinning solutions. Additionally, in Figure 15, the formation of nanowebs was observed among the fibers. The electrospun PVA/ZnO fibers acted as a support for the “fishnet-like” nanowebs comprising interlinked one-dimensional nanowires. The average diameter of the PVA/ZnO

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nanowires (10 nm) contained in this nanoweb was about one order magnitude less than that of conventional electrospun fibers. The formation of the nanowebs was considered to be due to the electrically forced fast phase separation of the charged droplets which move at high speed between capillary tip and the collector. The phenomenon of nanoweb formation has been reported in our previous study [55]. Here, the observation of PVA/ZnO nanowebs was ascribed to the sharply increased conductivity from 0.019 to 0.712 S/m (Table 1).

(A)

(B)

Figure 15. SEM images of composite PVA/ZnO electrospun nanofibers.

Table 1. Properties of PVA and PVA/zinc acetate solutions Sample PVA PVA/zinc acetate

Viscosity (centipoises) 420 600

Conductivity (S/m) 0.019 0.712

Another example for nanoweb formation is the PVA/SiO2 nanowebs by electrospinning the suspension of PVA and SiO2 nanoparticles. SEM images of PVA/SiO2 nanowebs are shown in Figure 16. The formation of nanowebs probably was due to the increased instability by blending the high content nanoparticles in polymer solution.

(A)

(B)

Figure 16. SEM images of composite PVA/SiO2 electrospun nanofibers.

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Concusion In this chapter, we have reviewed the fabrication of PAA, nylon-6, PVA/ZnO, and PVA/SiO2 nanowebs that stacked as layer-by-layer and widely distributed in the 3D structure of fibrous mats. The development of PAA nanowebs indicated the formation process of nanowebs. The splitting of defect films into nanowebs occurred under the extreme processing conditions, such as the high applied voltage, low relative humidity, and fast phase separation between polymers and solvents. Meanwhile, the morphology and area density of nanowebs were found to be the comprehensive effect of various electrospinning processing parameters of relative humidity, applied voltage, kinds of solvents, distance, and solution concentration. The typical nanoweb-rich sample nylon-6 showed that the nanowire diameter in nanowebs was one tenth of common electrospun fibers. As the formation of defect films was usual phenomenon during electrospinning, other materials with electrospinability also have the possibilities to form nanowebs. Additionally, the PAA nanowebs showed the larger gas absorption capacity and higher sensitivity to ammonia compared with PAA nanofibers due to its relatively high specific surface. This review collected some data from our previous publication of reference 55.

References [1] [2] [3] [4] [5] [6] [7] [8] [9] [10] [11] [12] [13] [14] [15] [16] [17] [18] [19] [20] [21]

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[22] K. Onozuka, B. Ding, Y. Tsuge, T. Naka, M. Yamazaki, S. Sugi, S. Ohno, M. Yoshikawa and S. Shiratori, Nanotechnology, 17 (2006) 1026. [23] H. Kokubo, B. Ding, T. Naka, H. Tsuchihira and S. Shiratori, Nanotechnology 18 (2007) 165604. [24] B. Ding, C.R. Li, Y. Hotta, O. Kuwaki and S. Shiratori, Nanotechnoligy 17 (2006) 4332. [25] Y. Miyauchi, B. Ding and S. Shiratori, Nanotechnology 17 (2006) 5151. [26] T. Ogawa, B. Ding, Y. Sone and S. Shiratori, Nanotechnology 18 (2007) 165607. [27] B. Ding, T. Ogawa, J. Kim, K. Fujimoto and S. Shiratori, Thin Solid Films 516 (2008) 2495. [28] M. Kanehata, B. Ding and S. Shiratori, Nanotechnology, 18 (2007) 315602. [29] D. Li, Y. Wang and Y. Xia, Adv. Mater. 16 (2004) 361. [30] P. Katta, M. Alessandro, R.D. Ramsier and G.G. Chase, Nano Lett. 4 (2004) 2215. [31] Theron, E. Zussman and A. Yarin, Nanotechnology 12 (2001) 384. [32] K. Yang, D. Edie and D. Lim, Carbon 41 (2003) 2039. [33] J. Gong, C. Shao, G. Yang, Y. Pan and L. Qu, Inorg. Chem. Commun. 6 (2003) 916. [34] B. Ding, J. Gong, J. Kim and S. Shiratori, Nanotechnology, 16 (2005) 785. [35] D. Li and Y. Xia, (2003) Nano Lett. 3 555. [36] B. Ding, H. Kim, C. Kim, M. Khil and S. Park, Nanotechnology 14 (2003) 532. [37] B. Ding, C. Kim, H. Kim, M. Seo and S. Park, Fiber. Polym. 5 (2004) 105. [38] R. Caruso, J. Schattka and A. Greiner, Adv. Mater. 13 (2001) 1577. [39] B. Ding, J. Kim, E. Kimura and S. Shiratori, Nanotechnology 15 (2004) 913. [40] B. Ding, K. Fujimoto and S. Shiratori, Thin Solid Films, 491 (2005) 23. [41] B. Ding, C.R. Li, S. Fujita and S. Shiratori, Colloids and Surfaces A: Physicochem. Eng. Aspects 284-285 (2006) 257. [42] D. Li, J. Mccann and Y. Xia, Small, 1 (2005) 83. [43] D. Li and Y. Xia, Nano Lett. 4 (2004) 933. [44] J.H. Yu, S.V. Fridrikh and G.C. Rutledge, Adv. Mater. 16 (2004) 1562. [45] M. Bognitzki, W. Czado, T. Frese, A. Schaper, M. Hellwig, M. Steinhart, A. Greiner and J. Wendorff, Adv. Mater. 13 (2001) 70. [46] S. Megelski, J. Stephens, D. Chase and J. Rabolt, Marcromolecules, 35 (2002) 8456. [47] B. Ding, M. Yamazaki and S. Shiraori, Sensor. Actuat. B 106 (2005) 477. [48] C. Casper, J. Stephens, N. Tassi and J. Rabolt, Marcromolecules, 37 (2004) 573. [49] C. Mit-uppatham, M. Nithitanakul and P. Supaphol, Macromol. Chem. Phys. 205 (2004) 2327. [50] R.L. Grimm and J.L. Beauchamp, J. Phys. Chem. B 109 (2005) 8244. [51] Z. Olumee, J. Callahan and A. Vertes, J. Phys. Chem. A 102 (1998) 9154. [52] S. Hong, J. Moon, J. Lim, S. Kim and S. Yang, Langmui. 21 (2005) 10416. [53] T. Young, D. Lin, J. Gau, W. Chuang and L. Cheng, Polymer. 40 (1999) 5011. [54] H. Fong, I. Chun and D. Reneker, Polymer. 40 (1999) 4582. [55] B. Ding, C.R. Li, Y. Miyauchi, O. Kuwaki and S. Shiratori, Nanotechnology, 17 (2006) 3685.

In: Nanotechnology… Editors: C.J. Dixon and O.W. Curtines, pp. 71-105

ISBN: 978-1-60692-162-3 © 2010 Nova Science Publishers, Inc.

Chapter 3

NANO-SCALE CHARACTERIZATION AND SPECTROSCOPY OF STRAINED SILICON Norihiko Hayazawa* and Alvarado Tarun Nanophotonics Laboratory, RIKEN, 2-1 Hirosawa, Wako, Saitama 351-0198, Japan

Abstract Strained silicon (ε-Si), the fundamental material of integrated circuit, is finding tremendous attention because it boosts the speed and reduces the power consumptions of electronic devices. However, poor homogeneity distribution of strain in ε-Si layers can degrade performance of electronic devices. Raman spectroscopy is used to study strain fluctuations in silicon because the optical phonons in Raman spectra are strongly influenced by strain. Though silicon are Raman active devices, the Raman efficiency of a nanometer layer of ε-Si is extremely weak and is often eclipsed under the Raman scattering of underlying buffer substrates. Micro Raman measurements show only uniform features in the nano-scale because of averaging effect from diffraction-limited spatial resolution. Here, we utilized surface enhancement in Raman scattering to overcome weak emission problems and to suppress averaging effect. Thin ε-Si layers were covered with thin Ag layer to invoke surface enhanced Raman spectroscopy (SERS). Results show that SERS effectively enhanced the Raman signal from ε-Si layer and it stands distinctly apart from the Raman signal originating from the buffer layer. This technique is promising but it lacks the spatial resolution in the nano-scale due to diffraction limit from the probing light. In order to achieve nano-scale spectroscopy, point-surface-enhancement was used, rather than a large surface enhancement. We used a silver-coated sharp tip, just like SERS, but only the sample region very close to the tip apex is characterized. This technique, known as the tip-enhanced Raman spectroscopy (TERS), provides nanometric resolution in our measurement. We observed localized strains by employing TERS. The TERS spectra revealed clear nano-scale variation in Raman frequency. Now that we can distinctively separate ε-Si from underlying buffer layer, signal-to-noise ration (SNR) needs further improvement. We improve TERS SNR in two ways: optical field enhancement using different metallic tip and background signals reduction arising from bulk materials. The tip-enhancement is more important for homogenous nano-materials or for samples with very weak signals whereas the background signal reduction is indispensable for nano-materials that consist of different thin layers with strong signals such as ε-Si or samples *

E-mail address: [email protected]. phone: +81-48-467-9339, fax: +81-48-467-9170. (Author for Correspondence)

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Norihiko Hayazawa and Alvarado Tarun with strong signal level. Accordingly, we introduce several approaches mainly for the suppression of background signals arising from other bulk materials. We will discuss the utilization of UV light source, specialized tip, sample orientation relative to probing polarization, and depolarization configuration to obtain high contrast Raman signal. The characterization techniques describe above is applicable to other nano-materials.

1. Introduction Strained Si (ε-Si) on silicon germanium (Si1-xGex) substrate has now been widely used in state of the art electronic devices such as Intel dual core products primarily due to the mobility enhancement of electrons and holes in the ε-Si layer [1,2]. As ε-Si technologies become more complicated and transistors become smaller, in the order of nanometer (~45nm), issues involving strain defects and the characterization of key properties in silicon substrates are expected to increase. Thus, development of advanced technique for nanoscale characterization and investigation of localized strain in the ε-Si substrate are important in the current semiconductor technology. Raman spectroscopy has been widely used for the investigations of strain in semiconductors [3-5] because of its nondestructive capability to observe lattice vibrations sensitive to strains. However, due to the low Raman scattering efficiency, it becomes difficult to observe the localized strain in nanometer scale. Moreover, in case of thin ε-Si layer fabricated on Si1-xGex substrate, the Raman scattering of Si-Si vibrational mode from Si1-xGex buffer layer can easily eclipse the vibrational signals originating from the ε-Si layer. Because of these two difficulties, conventional Raman scattering is not suitable for strain measurements in these thin samples. Nanometer scale and surface selective sensitivities are required, and hence surface enhanced Raman scattering (SERS) [6] and tip-enhanced Raman spectroscopy (TERS) [7-13] turn out to be exceptionally promising spectroscopic techniques. In this chapter, we demonstrate how effectively and selectively enhance the Raman signal from the thin ε-Si layer on a thick underlying Si1-xGex layer using SERS. We show that the proposed technique can be straightforwardly used for TERS microscopy to investigate localized nano-scaled strain in ε-Si. In section 2, we discuss SERS from the thin ε-Si covered with thin Ag layer. Results show that SERS effectively enhanced the Raman signal originating from ε-Si layer and it stands distinctly apart from the Raman signal originating from the buffer layer. This technique is promising but it lacks the spatial resolution in the nano-scale due to diffraction limit from the probing light. In section 3, in order to achieve nano-scale spectroscopy, point-surface-enhancement was used, rather than a large surface enhancement. We used a sharp Ag-coated tip placed over the ε-Si and only the sample region very close to the tip apex is characterized. This technique, known as the tip-enhanced Raman spectroscopy (TERS), provides nanometric resolution in our measurement. We observed localized strains by employing TERS. The TERS spectra revealed clear nano-scale variation in the Raman frequency. Now that we can distinctively separate ε-Si from underlying buffer layer, signal-to-noise ration (SNR) needs further improvement. We improve TERS SNR in two ways: optical field enhancement using different metallic tip and background signals reduction arising from bulk materials. The tipenhancement is more important for homogenous nano-materials or for samples with very weak signals whereas the background signal reduction is indispensable for nano-materials that consist of different thin layers with strong signals such as ε-Si or samples with strong

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Raman signal level. Accordingly, we introduce several approaches mainly for the suppression of background signals arising from other bulk materials. We will discuss the utilization of UV light source and specialized tip in section 4 and then introduce depolarization configuration in TERS for the efficient suppression of background signals in section 5.

2. Surface Sensitive Detection by Surface Enhanced Raman Scattering (SERS) Recent developments in surface enhanced Raman scattering (SERS) has gained attention not only due to enormous (1014) level of Raman enhancement capable of single-molecule detection but also for surface sensitive characterization of nanometer structured surfaces. It was observed that huge increase in the Raman cross section was result of excitation of surface plasmon in metals. The most popular metal used in SERS studies is Ag due to intense plasmon resonance excitation produced by nanometer sized Ag in the visible wavelength. In this section, we utilize SERS to obtain strain information in a thin layer, ~30nm, of ε-Si over silicon germanium substrate.

2.1. Experimental Configuration for SERS Figure 1 illustrates the concept of the SERS experimental system for highly sensitive detection of strain in the ε-Si surface. The substrate used for growing ε-Si layer was prepared by doping Ge in a pure 〈100〉 oriented silicon substrate. The concentration of Ge in the host Si lattice was gradually increased up to 25 % to expand the lattice constant of the host Si lattice. In order to minimize inhomogeneous strain, a layer with 1.0 μm thick Si1-xGex buffer layer was grown on the prepared substrate at constant 25% Ge concentration. Finally, a 30-nmthick Si layer was epitaxially grown on the buffer layer to fabricate ε-Si. The crystal face for both the buffer and silicon layer was maintained at 〈100〉. Due to lattice parameter mismatch at the interface, the thin Si layer was under a constant strain. The thickness of this ε-Si layer has to be low otherwise the induced strain could be gradually relaxed with increasing the thickness. To detect Raman from a very thin layer, we invoked SERS effects by evaporating Ag with 8-10 nm thickness on the ε-Si layer under vacuum (~10-6 Torr). The topographic image is shown in the inset of Fig. 1. This image was obtained by contact-mode atomic force microscope (AFM) using a silicon cantilever (Nanosensor ATEC-CONT). The Ag film consists of plenty of Ag grains, with typical grain size of 20~40 nm in diameter and 10 nm in height. These fractal-like "pancake” structures of Ag grains provide almost uniform surface enhancement factor averaged within the diffraction limited focused spot of micro-Raman experiments. A CW blue laser (wavelength: 488 nm, power: 9 mW) was focused by an oilimmersion high numerical aperture (NA) objective lens (NA=1.4; 100x). The gap of Ag film layer and the sample was filled with oil that has a matching refractive index (1.515) as the objective lens. The Raman and SERS signals from the sample were collected by the same objective lens and detected by confocal Raman microscope (pinhole: 30 μm), in which the spectrometer (focal length: 560 mm, 1800 grooves/mm, blaze wavelength: 500 nm) was equipped with both, a cooled charge coupled device (CCD) camera for Raman spectra

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measurements (such as the results in Fig. 2), and an avalanche photodiode for Raman image acquisitions (such as in Fig. 3).

Figure 1. Schematic of SERS microscopy for ε-Si and AFM topographic image of the Ag island film coated on the ε-Si.

2.2. Comparison of Raman and SERS Spectra Vibrational modes located at the center of the Brillouin zone can only be observed in the first-order Raman scattering due to the momentum conservation. Further, the selection rules for the 〈100〉 oriented Si crystal in backscattering geometry allow only LO phonon to be Raman active. Therefore, we expect to observe only LO phonon in first-order Raman scattering experiments, which should appear at 520 cm-1 for an unstrained or pure Si crystal. Since the sample consists of a ε-Si layer above Ge-doped Si buffer layer, we expect shift from the original frequency of LO phonon in both layers. The LO phonons in doped crystals can be easily modified depending on the atomic mass as well as the amount of the dopant. In Gedoped Si crystal, the LO phonon can have three different vibrations, namely the Si-Si mode arising from the vibrations of neighboring Si atoms, the Ge-Ge mode arising from the vibrations of neighboring Ge atoms, and the Si-Ge mode arising from the vibrations involving the bonds between a Si and a Ge atom. Since the amount of Ge is much smaller than the amount of Si in our samples, the LO mode corresponding to the Si-Si vibration is dominant. We need to concentrate only on Si-Si mode because it interferes with the Si-Si vibration coming from the thin ε-Si layer in the experiments and we will be discussed throughout the chapter. The vibrational frequency of this mode decreases almost linearly with increasing Ge concentration [14]. The bonds between two neighboring Si atoms are stretched when Ge is doped introducing strain in the lattice. This strain is responsible for the shift of LO mode frequency originating from the Si-Si vibration in the buffer layer. Similarly, when a thin Si layer is grown on Si1-xGex buffer layer, the Si-Si bonds in Si layer are stretched due to the lattice mismatch at the interface. If the thickness of the layer is not too large, the strain produced in the Si layer is transferred vertically throughout the surface. The amount of

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developed strain at surface can be different from that of the strain in the buffer layer even if the strain in both layers is caused by the presence of Ge atoms in the buffer layer. The penetration depth of blue-green probing light in Si is much larger than the thickness of the ε-Si layer in our sample [15]. Therefore, in Raman scattering experiments, the probing laser penetrates deep through the buffer layer. As a result, Raman scattering in the buffer layer is strongly excited. The intensity of LO phonon mode coming from the buffer layer is much more strong than the one coming from the ε-Si layer. This is a problem if the vibrational frequencies of the two phonons are close to each other. It becomes difficult to clearly observe the LO phonon corresponding to the ε-Si layer. To overcome this problem, we used the SERS technique to selectively enhance Raman at the surface and compared the results from those obtained by normal Raman spectroscopy. Figure 2 shows SERS (with a Ag film deposited on ε-Si layer) and Raman spectra (without a Ag film) of the sample. Both spectra were fitted by double Lorentzian functions as shown by dashed lines. The lower and higher Raman modes in both spectra are the LO phonon modes arising from the Si-Si vibration in Si1-xGex (x=0.25) and in ε-Si, respectively. The Raman spectrum is dominated by a strong peak at 504.9 cm-1, corresponding to the background LO phonon signal from Si1-xGex substrate, together with a very weak shoulder at 513.8 cm-1 that arise from the LO phonon in ε-Si [16-20]. Since the Raman signal from the ε-Si is very close to the Si1-xGex buffer layer and considering 30 nm thickness of the ε-Si layer, the Raman signal from ε-Si is almost overwhelmed by the signal from the Si1-xGex layer. On the other hand, in SERS spectrum, the Raman signal from ε-Si is remarkably enhanced. The signal becomes even stronger than the Si1-xGex buffer layer due to the SERS effect. Even though the ε-Si Raman signal in SERS is enhanced relative to the signal from the buffer layer, we noticed reduction in the scattering efficiency in SERS with a signal bias of ~20000 counts/60s compared to the Raman spectrum. The reduction in overall scattering efficiency is due to the lower transmission of the probing light caused by Ag island film whereas the bias signal is due to the white continuum emission from the Ag island film, which is often accompanied in SERS measurements [21].

Figure 2. SERS (with Ag film) and Raman (without Ag film) spectra of ε-Si. Both spectra are fitted by double Lorentzian functions representing Raman modes of Si1-xGex (lower frequency) and ε-Si (higher frequency). The meshed areas are used for Raman imaging in Fig. 3.

As seen from Fig. 2, we resolved Raman signal from the ε-Si in SERS. This suggests that SERS effect is very powerful to observe the surface condition of the thin ε-Si layer. Both SERS and Raman measurements were done in the same condition except the existence of Ag

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island film to invoke SERS. It should be noted here that we did not use any analyzer for polarization dependent detection in either SERS or Raman detection. The ratio of the intensities between Si1-xGex and ε-Si would change when measured through an analyzer, which will be discussed later in Section 5.

2.3. Comparison of Raman and SERS Images Next, we carried out confocal micro-Raman imaging of the same sample at specific vibrational modes from Si1-xGex and ε-Si. Figure 3(a) shows the reflection confocal image of the sample from a region close to the edge of the evaporated Ag film. As expected, the Ag covered layer appears to be brighter in the reflection image. Figure 3(b) was obtained by setting the spectrometer at the lower frequency in Fig. 2, so that the avalanche diode detects only the Raman signal from Si1-xGex buffer layer. Since the intensities level of SERS and Raman from Si1-xGex buffer layer is almost similar from Fig. 2, Raman image in Fig. 3(b) also reflects low contrast between Ag coated and uncoated area. However, when the spectrometer is set to detect only the higher frequency shift arising from ε-Si (see Fig. 2), the enhanced Raman signal from Ag coated ε-Si has high contrast as depicted in Fig. 3(c). Fig. 3(c) does not reflect any particular information about the possible nanometric strain distribution in ε-Si, which could be varying within areas smaller than the focal spot. This is because the enhanced scattering signal in this experiment is averaged within the diffraction

Figure 3. (a) A reflection confocal and Raman images obtained at Raman shift of (b) Si1-xGex and (c) εSi as indicated by meshed area in Fig. 2. (d) is obtained at Raman shift of ε-Si, the same as (c) but with different polarization of excitation light. An arrow in the figure indicates polarization of an excitation light in each image. The dimension of each image is 45 μm by 45 μm consisting of 100 pixels by 100 pixels.

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limited focused spot. Accordingly, when polarization of the excitation light is rotated 90 from the previous polarization as indicated in the Fig 3(d), no significant change in the contrast could be observed. The entire signal level in Fig. 3 (d) is decreased compared to that in Fig 3(c) due to the polarization dependence of crystalline silicon 〈100〉 facet [22,23] and also due to the different detection efficiency of the experimental system. Here, we note that the temperature dependence of Raman scattering in silicon [24] is not observed in our experiments under room temperature conditions. However, the temperature dependence is also an important factor that can influences the properties of ε-Si substrates and could be observed under the experimental conditions of SERS.

3. Surface Sensitive and High Spatially Resolved Detection by Tip-Enhanced Raman Scattering (TERS) In SERS measurement, high surface sensitivity and signal enhancement was successfully achieved, however, the spatial resolution is still diffraction limited. Moreover, SERS is an irreversible or destructive detection process because it requires Ag coating on the sample. For non-destructive and higher spatial resolution measurement capable of detecting localized nanometric strain, a sharpened metallic tip [25-27] can be utilized instead of Ag island film. In this section we introduce the use of “tip-enhancement” instead of “surface enhancement” to go beyond the diffraction limit of probing light. The concept of TERS is rather straightforward that can be understood by considering the sharp apex of a metallic tip as a point source surface enhancer [7-10].

3.1. Experimental Configuration of Reflection-Mode TERS Figure 4 shows the concept of reflection mode TERS microscopy [28-31] for opaque sample such as ε-Si layer assembled on thick substrate. There are generally two configurations for TERS, transmission mode and reflection mode. In transmission mode [713,32], the metallic probe tip is illuminated through the sample using objective lens. This configuration has been widely used in TERS because of its high efficiency in both illumination and collection, especially when using a high numerical aperture (NA) objective lens. However, transmission mode TERS cannot be used for opaque or thick samples. In reflection mode TERS [28-31], the tip is illuminated from the same side in which it is approached towards the sample (Fig. 4). This is advantageous and promising for observing opaque samples such as silicon substrates. A schematic of the experimental configuration is shown in Fig. 5 [33,34]. An incident CW YAG laser (532 nm, JDS Uniphase) is directed to the optical set-up through a single mode optical fiber. The light coming out of fiber is set to p-polarization, parallel to the tip axis, for optimum tip-enhancement [28] from the combination of a half-waveplate and a polarizer. The beam diameter of the laser is then expanded to ~10 mm. This expanded and polarized light is focused on the tip and sample using a long-working-distance (LWD) objective lens (Mitsutoyo, NA=0.28, 20x, WD=30.5 mm). The illumination optics mentioned above, including the LWD objective lens, is precisely positioned by three-axis actuators (New Focus Inc. Picomotor, accuracy <30nm). The apex of the cantilever (Nanosensor ATEC-

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Figure 4. Schematic of reflection-mode TERS.

Figure 5. Experimental setup of reflection-mode TERS. The inset shows the structure of ε-Si on Si1(x = 0.25).

xGex

CONT) is adjusted onto the focused spot using the same principle as a contact-mode operation of AFM. The power of the incident laser light is set to deliver 1 mW on the sample. The same LWD objective lens collects the TERS signal. A dichroic mirror (or a nonpolarized beam splitter cube) directs the output Raman signal via a multimode optical fiber to a spectrometer (Photon Design PPDT3-640). The spectrum is detected using a liquid nitrogen-cooled CCD camera (Princeton Instruments). Raman signal was obtained with an

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integration time of 120 seconds. To attain an efficient tip-enhancement effect in the reflection-mode TERS, the probe tip is prepared by depositing Ag on commercially available silicon cantilever used for AFM. This is done by two-step thermal evaporation process. First, a 5-nm-thick gold palladium alloy (AuPd) is deposited on a tip to modify the surface of silicon, and then Ag is evaporated with the thickness of 30~50 nm at a rate of 0.5 Å/sec.

3.2. Comparison of Raman and TERS Spectra The reflection mode TERS described above has been successfully applied to nanoscale localization of ε-Si using dichroic mirror and Ag coated Si-tips with an excitation laser wavelength of 532nm [33]. Figure 6(a) shows the Raman signal from 532nm probing light under Ag coated Si-tip (Nanosensor, ATEC-CONT-20). Figure 6(b) shows the background corrected signal, which is the difference between the TERS and far-field (tip withdrawn)

Figure 6. (a) TERS (red) and far-field (blue) spectra of ε-Si obtained using 532nm probing light with Ag coated Si-tip. (b) Background corrected near-field spectrum that is broken down into three Lorentzian curves for band peak identification. The peaks at 503, 514, and 520 cm-1 are assigned to phonon modes from SiGe layer, ε-Si layer and Si-tip, respectively.

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Raman signal. In order to determine the Raman peaks, we decomposed the corrected spectrum in Fig. 6(b), into three Lorentzian functions. The three peaks are assigned to Si-Si vibrations in the SiGe (503 cm-1) bulk layer, Si-Si phonon mode in ε-Si (514 cm-1) layer and Si-Si in Si-tip (520 cm-1). The corrected spectrum does not contain the peaks around 480 cm1 , which was later found to be coming from PC monitor. This background corrected TERS spectrum shown by the spectrum in Fig. 6(b), contains nanometer scale information of localized strains. In the far-field spectrum, the ε-Si peak is recognized only as a shoulder of the strong background signal generated from the underlying Si1-xGex (x=0.25) layer. The TERS spectrum, however, shows a clear and distinct peak owing to tip-enhancement. Since the ε-Si layer is in the order of several tens of nanometers, a sufficient amount of incident light penetrates into the Si1-xGex buffer layer, and hence the background signal from the underlying buffer layer is still significant. The shift of LO phonon mode originating from the ε-Si with respect to the original position of the LO phonon for an unstrained Si (520 cm-1) [3] provides us with plenty of information, such as the local strain in the layer [35-38]. The estimated enhancement factor for ε-Si is 2.2 x 104[33], which is calculated by considering the size of the diffraction limited focused spot (φ: 3 μm) and the enhanced electric field corresponding to the tip diameter (φ: 40 nm).

3.3. TERS Spectra Mapping Nanoscale lateral heterogeneity of strain in ε-Si / Si1-xGex structures is a serious problem in fabrication of high-mobility electronic devices. Therefore, it is necessary to investigate the heterogeneity of strain at nanoscale. A position dependent TERS measurement can be used for a detailed characterization. Figure 7(a) shows the bright-field microscopic image of the εSi surface. The crosshatch patterns induced by lattice-mismatch are seen on the epitaxial surface of the ε-Si. Fig. 7(b) illustrates the contact AFM image (3 μm x 3 μm) of the crosshatch pattern. The six red crosses indicate positions where the tip-enhanced measurements were carried out. Fig. 7(c) shows near-field (TERS) spectra of ε-Si extracted by Lorentzian curve fitting. From Fig. 7(c), we can observe that the peak position and the intensity of the ε-Si change. The changes are due to variations in thickness of the ε-Si layer based on an underlying surface topography in Fig. 7(b). Fig. 7(d) corresponds to the Si1-xGex Raman spectra. The spectral pattern in the Si1-xGex under layer depends both on the Ge concentration and presence of any residual defects in the Si1-xGex due to lattice mismatch during heteroepitaxial growth[36,37]. The dark areas, c and b have lower ε-Si intensity compared to bright areas, a, d, e, and f, in the AFM image. If the intensity of ε-Si (Fig. 7(c)) is weak, the intensity of Si1-xGex (Fig. 7(d)) is strong. From the cross sectional view of Fig. 7(b), the step height between the bright to the dark area is ~16 nm. Figure 7(c) also illustrates the slight shift (~ 2 cm-1) of the ε-Si peaks. While the average deviation of the peak of the silicon tip stays within a range 520 ± 0.04 cm-1 (data not shown), its value in the ε-Si spectrum is 514 ± 0.8 cm-1. Raman shift of Si-Si vibration basically provides a measure of the inter-atomic Si-Si spacing in the ε-Si layer. This spacing causes fluctuation of channel strains, and subsequently affects the performance of the device. It has been reported that lateral strain variations is mostly associated with the crosshatched pattern and its variation is in the order of

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~1 cm-1 [36]. Calculated from the shift in Raman peaks, the strain imposed on the capped ε-Si layer is about 1.5 ± 0.2 GPa [37].

Figure 7. (a) Reflection microscopic image of the ε-Si surface. (b) AFM image of the ε-Si surface. Red crosses indicate the tip positions where TERS mapping was carried out. (c) and (d) are Lorentzian fitting results of position dependent TERS spectra of ε-Si and Si1-xGex, respectively. a-f correspond to the positions shown in (b).

Figure 7(d) shows the subtracted near-field spectrum for the LO phonon mode arising from Si1-xGex buffer layer. The LO phonon mode in the underlying Si1-xGex layer did not show much spectral shift during the spectra mapping except a slight shift observed in the spectrum d. This can be attributed to the local density of doped Ge, which could be higher around the position d. This is also evident from Fig 7(c) point d, which exhibits the highest spectral shift for the LO phonon mode in ε-Si layer. The nearby spectra obtained at c and e also have slight spectral shifts. This means that, at position d, we have high concentration of defects associated to the heterogeneity in the Ge-concentration in the underlying layer. This is the direct evidence of the fact that the Si-Si bond length in ε-Si is affected by the crystalline structures and concentration of Ge in the underlying Si1-xGex layer on nanoscale. However, the observed spectral shifts in the ε-Si (~2 cm-1) are higher than those in the Si1-xGex layer (~1 cm-1). This is understandable, because even though the strains in both layers are caused by the presence of Ge atoms in the buffer layer, the amount of strain may not be the same. All these observed nanometer scale information on localized strain cannot be detectable by micro SERS experiments due to averaging effect within a diffraction limited focused spot

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[34]. Using TERS, however, we could successfully reveal the localized strain hidden in nanometer scale.

4. Managing the Excitation Wavelength and the Tip Material for TERS So far, we have demonstrated selective enhancement of LO phonon signals arising from the thin ε-Si layer by SERS and TERS in the presence of strong background scattering from Si1-xGex buffer layer to distinctly observe Raman signal and successfully revealed localized strain in nanometer scale using TERS. However, as can be seen from Fig. 6, the LO phonon from the ε-Si layer is sandwiched between two comparatively strong unwanted peaks arising from the Si1-xGex buffer layer and Si-tip material. In order to have higher sensitivity and to obtain high contrast image of the strain distribution in silicon layer by TERS, it is important to get rid of these two unwanted background scattering modes. We address this by using shorter wavelength and using different tip material.

4.1. Ultraviolet Light Excitation for Background Suppression from Underlayers Since the strong scattering from the buffer layer is excited due to the longer penetration depth of the probing light, this problem can be overcome only if the penetration depth of the probing light in silicon can be reduced. It has been reported [15] that by varying the excitation wavelength from 532 to 351 nm, the penetration depth of the probing light in silicon can be changed from 400 to 5 nm. For example, the penetration depth is 5 nm, 47 nm, 108 nm, 241 nm, and 406 nm for the excitation wavelength at 351 nm, 405 nm, 442 nm, 488 nm, and 532 nm, respectively [39]. Moreover, resonant Raman scattering in Si has been reported [40] for excitation around 365 nm (3.4 eV). Thus, an increase in the phonon intensity due to the resonance effect can also be observed for lower excitation wavelengths, which is an additional advantage in using the lower excitation wavelengths. However, the use of Agcoated tip in TERS limits the wavelength that can be effectively used for enhancement. This is because the field enhancement near the tip apex comes from the plasmon resonance within the Ag layer coated on the tip [41]. By reducing the excitation wavelength, we move away from the plasmon resonance of the tip, and hence lose the enhancement of the field at the tip apex. Therefore, for an efficient TERS detection from the thin layer of ε-Si, we need to find a compromise between the two opposite effects of reducing the excitation wavelength. One, a positive effect by reducing the penetration depth together with increasing Raman efficiency due to resonance Raman effect, and the other, a negative effect by reducing the field enhancement at the tip apex. To compromise for the negative effect of using shorter wavelength, we performed far-field Raman measurements in different excitation wavelength. Figure 8(a) shows the far-field Raman spectra (without tip) obtained from sample. Each spectrum is normalized and then decomposed into two Lorentzian function (dashed lines) for peak intensity analysis. The relative intensity of the Si-Si in ε-Si relative the Si-Si in Si1-xGex buffer layer increases with decreasing wavelength. The total intensity for each phonon mode

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Figure 8. (a) Far-field Raman intensity of ε-Si layer grown in the SiGe layer at different wavelength. The Lorentzian functions under Si-Si in Si1-xGex (x=0.25) peaks at 502.6, 503.4 and 502.6 cm-1, which corresponds to phonon modes from 442nm, 488nm and 532nm, respectively. (b) Total Raman intensity ratio (blue) between ε-Si and SiGe and penetration depth (red) as a function of excitation wavelength. The Raman intensity ratios were obtained by calculating the area under the Lorentzian curves.

in Si1-xGex and ε-Si were obtained by calculating the area under the two Lorenztian curves. The relative intensity of the LO phonon mode arising from ε-Si layer is increased by shortening the excitation wavelength. Figure 8(b) shows the total Raman intensity ratio of εSi over the Si1-xGex at varying wavelengths. The secondary y-axis in Fig. 8(b) is the calculated penetration depth for each excitation wavelength used. As the penetration depth decreases with decreasing wavelength, the total Raman intensity ratio between ε-Si and Si1xGex increases. At 442nm excitation wavelength, the penetration is ~108nm and still bigger than the 30nm thickness of ε-Si sample. A laser with 405nm excitation wavelength may be a good candidate because the penetration is ~47nm and near the 30nm ε-Si thickness. At 442nm, the Raman intensity from ε-Si is increased to half the intensity of Si1-xGex. This is a significant improvement compared to 532nm light where only a small ε-Si Raman signal is observed (see Fig. 6). The incurred loss in enhancement due to shorter wavelength was offset

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by the gain in the ratio between ε-Si and SiGe signal. Reducing the excitation wavelength allows us to obtain effectively background-free scattering signal from the ε-Si layer, leading to a feasible imaging of strain in ε-Si substrates. The field enhancement at the tip apex is still reasonable for the He-Cd laser with excitation wavelength of 441.6 nm (data not shown here), so we selected this wavelength in further TERS aimed towards higher sensitivity detections.

4.2. Selection of Tips for Background Suppression from Silicon Tips In order to further improve the sensitivity in TERS, we still need to address the other problem arising from the presence of the LO phonon mode from the tip itself. Since the tip is made of silicon and coated with Ag, it works as a perfect enhancer for the Si-Si scattering in the tip itself. This is because all the silicon atoms are adsorbed in the enhanced field created at the tip apex. Hence, it is impossible to suppress the scattering from the tip. The best way address the unwanted peak arising from Si-Si vibration mode in the Si-tip that interferes with the Raman signal from ε-Si layer, is to change the tip material. We need to select a material that will not interfere with sample vibrations. Unfortunately, most of the commercially available AFM cantilever tips are made of silicon, which is not a good tip-material for detecting TERS from silicon samples. However, cantilever tips made of silicon nitride (Si3N4) are now commercially available. The Raman vibrational modes of Si3N4 are distinctly separated from those of Si, which is a good tip-material for observing TERS from Si samples. Aside from Si3N4 tips, oxidized silicon tips [42,43] or silicon dioxide (SiO2) tips such as a tapered optical fiber probes can be also useful for the base material for metallized tips. If all Si atoms of the tip are oxidized, Raman vibration of Si from the tip can be completely suppressed, and hence this tip can be effectively use. It has been recently reported [44] that tips made of materials with lower refractive index, such as Si3N4 (n=2.05) or SiO2 (n=1.5) could show higher enhancement factors especially in shorter wavelength region.

Figure 9. A TEM image of an oxidized silicon cantilever tip.

Figure 9 shows the transmission electron microscope (TEM) image of the thermally oxidized silicon cantilever probes at 1100°C temperature under steam atmosphere for 3

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minutes. As seen in the figure, SiO2 layer with a thickness of ~100 nm is formed on a siliconbased tip without degrading the original sharpness. This oxidized silicon tips is also promising for tuning the plasmon resonance of metals coated on the tip by controlling the thickness of the SiO2 layer. Instead, for our experiments, we utilized the commercially available Si3N4 cantilever tips, after coating them with Ag. Figure 10 shows the SEM image of the Si3N4 tip (Budget Sensors, BS-SiNi) coated with 50-nm-thick Ag layer by thermal evaporation process. The diameter of the tip apex after Ag coating was about 30 nm, which is relatively smaller than the 50nm thickness of coated Ag because of the very sharp tip shape. Fig. 10 (b) is the SEM image (bottom view) of the cantilever. In order to illuminate the tip by a LWD objective lens with the angle of 30 degrees, part of the cantilever is removed by focused ion beam (FIB), otherwise it creates an unwanted shadow on the sample [28]. This was not necessary for the previous silicon tip, because the shapes of the bottom structures for the Si and Si3N4 tips are different. Even after chopping off a part of the tip, some shadow of the tip still existed for Si3N4 tips. It was impossible to reduce this shadow any further, so the overall intensity of Raman scattering was slightly reduced for TERS compared to the far-field measurements in the TERS experiments involving Si3N4. For reflection mode TERS, the geometrical accessibility of the tip for illumination should also be taken into consideration to minimize excitation power loss and fake scattering signals from the protruding parts of most commercially available tips.

Figure 10. SEM images of Ag coated Si3N4 tip. The area surrounded by dashed lines is removed by FIB for an efficient light delivery onto the tip by a LWD objective lens.

4.3. TERS Imaging In this section, we demonstrate TERS imaging on ε-Si layer using He-Cd laser with 441.6 nm wavelength and Ag coated Si3N4 cantilever tip. Figure 11 (a) shows the comparison between TERS (with a tip) and far-field Raman spectrum (without a tip). As expected, the SiSi background signal seen from the tip at 520 cm-1 in Fig. 6 is completely eradicated. Furthermore, the penetration depth of the excitation light at 441.6 nm in the underlying Si1xGex layer is drastically reduced compared with that at 532 nm. The far field Raman signal

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Figure 11. (a) Comparison of TERS and far-field Raman spectra, (b) Comparison of SERS and far-field Raman spectra.

from the thin ε-Si layer is relatively increased to the half of the intensity of the background signal from the Si1-xGex layer. This is a significant change compared to the Fig. 6(a), which uses 532nm, where only a small shoulder corresponding to the LO phonon arising from the εSi layer was observed. In the TERS spectrum in Fig. 11 (a), the background signal from Si1xGex layer is slightly reduced because of the shadowing effect of the tip. Even with shadowing effect by the tip, the signal from the ε-Si layer is increased because of the efficient tip-enhancement effect. In addition, Fig. 11 (b) shows the comparison between SERS (with Ag film; thickness ~ 8nm) and far-field Raman spectrum (without Ag film). The Raman spectrum without Ag film is the same without Ag tip in Fig. 11 (a). The presence of thin Ag layer reduces the intensity of excitation light reaching the sample due to absorption. This explains the reduction of the overall intensity of the Raman scattering in SERS spectrum compared to the far-field Raman spectrum. However, in the SERS spectrum, the relative intensity of the LO phonon arising from the ε-Si layer compared to the LO phonon intensity of the buffer layer is increased. This is once again a good demonstration of selective

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enhancement of the LO phonon vibration arising from the ε-Si layer. Comparing the experimental results of TERS and SERS in Fig. 11, the enhancement in the relative intensity of ε-Si with respect to the scattering from Si1-xGex buffer layer achieves almost same value. But, the scattering volume involved in TERS (about the same size as tip apex) is much smaller than that in SERS (same as the diffraction limited focused spot), which suggests that tip-enhancement with a Ag coated Si3N4 tip at 442 nm excitation works much more efficiently compared with surface enhancement at the same wavelength. Since the combination of a Si3N4 tip and 442 nm excitation shows much higher sensitivity in TERS measurements compared to the previous experiments using 532nm and Si tip, we employed this combination to image the nano-scale strain distribution on the surface of the ε-Si layer through near-field Raman imaging. In order to realize the nano distribution of the strain, the near-field Raman image was also compared with a far-field image from the same area of the sample. Figure 12(a) shows a topographic image (128 pixels by 128 pixels) and Figs. 12(b), and 12(c) show the strain images extracted from near-field and far-field Raman spectra, respectively, obtained from the same position (32 pixels by 32 pixels). In order to obtain these Raman images, the LO phonon modes from each position where fitted with Lorentzian functions, and then the peak positions were estimated from the fittings, which were plotted with respect to the position. The near-field image was constructed by the abovementioned process applied to the subtracted near-field spectra obtained by moving the tip at different positions. The far-field image on the other hand was constructed by applying the process to the far-field spectra by moving the center of the focused laser spot at different positions. Since the shift in the frequency position of the LO phonon mode is a direct reflection of the local strain, the images shown in Fig. 12(b) and 12(c) correspond to the distribution of the local strain in the sample. The near-field image clearly reveals strain embedded in nanoscale, which cannot be seen in the far-field image that shows almost uniform contrast. This confirms that there is nanometric distribution of the local strain in the sample, which could not be seen by the far-field imaging. Figures 12(d) and 12(e) show the plots of the number of pixels as a function of Raman peak position for the near-field image in Fig. 12(b) and for the far-field image in Fig. 12(c), respectively. Narrow Gaussian distribution produced in the far-field image is due to averaging effect within the diffraction limit while the widely broadened distribution in the near-field image is due to the observation of nano-scale strain distribution. This meets our expectation simply because near-field image in Fig. 12(b) shows large fluctuation of contrast compared to Fig. 12(c). It would be important to point out that the LO phonon peak positions in Fig. 12(d) is distributed in lower Raman shifts compared to Fig. 12(e), as can be seen from the shift of the Gaussian distributions in the two figures. There are two possible mechanisms for this effect; one is related to the chemical effects taking place between the Ag atoms of the coated tip and silicon atoms of the strained layer, and the other, due to the local pressure applied on the sample by the tip. Further investigation is required for confirmation. A careful observation of SERS spectrum in Fig. 11 reveals that there was a small downshift of Raman mode in SERS compared to far-field spectrum. It is also well know that Raman modes in SERS measurements are often downshifted [45-48]. The primary reason for such a shift is the chemical bonding between the Ag atoms and the sample atoms, which effectively modifies Raman vibrations. This can also happen for TERS, when Ag molecules of a coated tip come in close contact with the sample molecules. Since we used the contact mode operation of AFM to control the tip position in

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TERS experiment, the tip can apply a unidirectional force on the sample at the contact point. Even though the value of this force is not large, the resultant pressure per atom could be large enough due to the nano-sized of the tip apex. This effect was recently observed in various nanomaterials [45,49,50] under controlled application of pressure using tip. In our case, the tip applied pressure works in a direction perpendicular to sample surface. In the equilibrium condition, we can expect that the pressure to increase the bond lengths in the direction parallel to the sample surface. As a result, it adds to the strain present in the silicon layer due to the lattice mismatch between the silicon and the buffer layers. Hence, Raman modes are further shifted to the lower frequencies. For detailed characterization of strain distribution, some analytical investigations such as finite element method (FEM) are also very important factors together with the development of experimental system.

Figure 12. (a) topographic image and strain images of (b) near-field and (c) far-field, respectively. (d) and (e) are Raman peak distributions in (b) and (c), respectively.

5. Controlling the Polarization in Detection At this point, we tried to improve sensitivity by modifying the tip material, using shorter wavelength to reduce penetration depth and focus spot, and manipulating the well-defined polarization of crystalline sample along with the polarization conditions of probing light. The key issue is how to improve the sensitivity of the weak near-field signal at the nanoscale level to be detectable from the strong background coming from the total illuminated area. Most of the research works done have focused on electric field enhancement, either by Ag coated Si and Si3N4 tip or by pure metallic tip [11,29,32,51,52]. Various arguments were made

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regarding tip form and material [53-55] and the control of polarized excitation light [56,57]. Recently, researchers started working on reducing the far-field background signal without comprising the field enhancement to obtain high contrast near-field Raman image. One way is manipulating the well-defined polarization properties of crystalline sample relative to the polarization and propagation of the incident light. Ossikovski et al, proposed a simple model that describe polarized TERS to characterize tip-enhancement process [58]. The model is based on tip-amplification tensor describing the interaction of tip on both the incident and scattered optical field. From their results, by appropriately setting the polarizer and analyzer relative to the material crystalline axis, far-field signal can be reduced. Poborchii et al, proposed far-field suppression by depolarization configuration [59]. In this configuration, the tip depolarizes the incident beam from p-polarized light to s-polarized light or vise versa. The resulting depolarized light will incur depolarized Raman signal that can be detected through an analyzer while blocking the Raman signal from the originally polarized light to minimize the far-field Raman signals [60]. Sokolov’s group optimizes this depolarization configuration to obtain high contrasts Raman image on silicon material with lateral resolution of 20nm [61]. However, the authors noted that the optimum polarizer and analyzer condition, which gives highest contrast, changes with the morphology and geometry at the tip end. In this section, we used the silicon Raman tensors to estimate at which combination of incident polarization and sample orientation will offer minimum far-field Raman intensity. We also introduce different analyzer setting to further reduce the far-field background by taking into account the depolarization effect mentioned above. We performed SERS experiment to validate the numerical results and we found good agreement with the Raman tensor calculation. The configuration that gives minimum far-field background signals with high contrast in SERS was utilized in TERS experiment to obtain high contrast near-field Raman signal. In this approach, we fixed the polarizer and analyzer along with sample orientation based on SERS to ease the difficulty of optimizing the polarizer and analyzer setting for different tip geometry. Both the field enhancement effect and depolarized detection were considered to obtain a high signal-to-noise TERS signal. We found out that for Raman active thin samples assembled in bulk crystalline materials, depolarization effect outweighs the field enhancement effect in getting high-contrast Raman signal.

5.1. Theoretical Calculation of Raman Intensity by Side Illumination Optics The geometrical model used in the Raman intensity calculation is shown in Figure 13. This model was formulated for reflection-mode TERS system where a metallic tip is placed in contact with the ε-Si sample. There are two Cartesian coordinate systems, the sample (xyz) and the light (x’y’z’). The laser beam is introduced at φ=60 degrees relative to the sample normal (z-axis). The sample rotation angle (azimuth) is denoted as θ. The sample azimuth is varied to obtain the maximum and minimum Raman intensity level for both p- and spolarization of incident light.

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Figure 13. Schematic of reflection-mode TERS. TERS were defined with two Cartesian coordinate systems, the sample (xyz) and the light (x’y’z’) reference frames. Laser beam is introduced at φ relative to the sample’s normal (z-axis) and θ denotes the rotation angle (azimuth). The thickness of the strained Si is ~30nm. Electron micrograph of the Ag coated tip shows a tip-diameter of ~30nm.

The far-field Raman scattering intensity I is generally expressed as [22,62]:

I ∝ ∑ e Ts R j e i

2

(eq.1)

j

where R is Raman tensor for any j active phonon, and ei, es are the unit vector of incident and scattered electric field, respectively. The superscript T denotes transpose of the scattered electric field vector. The derivation of Raman tensor R for c-silicon in the (100) plane by back-scattering geometry (normal incidence) is discussed in detail by R. Loudon [22]. The csilicon Raman tensor for the phonon modes in the sample reference frame (xyz) is:

⎛0 0 ⎜ TO1 ⇒ R1 = ⎜ 0 0 ⎜0 d ⎝

⎛ 0 0 d⎞ ⎛0 0⎞ ⎟ ⎟ ⎜ ⎜ d ⎟ , TO 2 ⇒ R2 = ⎜ 0 0 0⎟ , LO ⇒ R3 = ⎜ d ⎜ d 0 0⎟ ⎜0 0 ⎟⎠ ⎠ ⎝ ⎝

d 0⎞ ⎟ 0 0⎟ (eq.2) 0 0⎟⎠

in which R1, R2, and R3 are respectively, the TO1, TO2, and LO, phonon modes of c-silicon and d is the constant that is dependent on the Raman polarizability of the sample. Note that in Raman spectra of Si there is only one very strong one-phonon peak at 519 cm-1 (at 305 K), corresponding to the zone-center optical phonons in Si, in which the TO and LO phonons are degenerate at zone center in diamond-type crystals [63]. The d value for the c-silicon can be different from ε-Si but for qualitative Raman intensity analysis, using c-silicon Raman tensor can provide a good approximation. Nevertheless, the Raman tensor of unstrained silicon expressed in eq. 2 is different from ε-Si because of the inhomogeneous tensile stress induced in the ε-Si layer. The amount of strain is maximized at the boundary of SiGe layer and ε-Si layer, then and gradually decreases and minimized at the surface layer of ε-Si. In order to

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obtain the qualitative comparison of Raman intensity in ε-Si at different incident and scattered polarizations with sample azimuth, we simply employ the Raman tensor of unstrained Si. To utilize the Raman tensors and scattered Raman intensity in eq.1 in reflection-mode TERS studies, the Raman tensors must be transformed to the laser illumination reference frame (x’y’z’) at angle φ. This is because the Raman tensors in eq. 2 were obtained in normal incidence. First, we need to rotate the sample azimuth about its normal (along the z-axis) in the xyz coordinate, and then formalize the z-component of the transformation matrix, Tz(θ). The Raman polarizability tensor of the sample rotated by an angle θ would become,

Rj ′( θ ) = TzT ( θ )RjTz( θ )

(eq.3)

where, Tz(θ) is the rotational matrix for a sample and given by,

⎛ cos(θ ) −sin(θ ) 0⎞ ⎟ ⎜ Tz (θ ) = ⎜ sin(θ ) cos(θ ) 0⎟ ⎜ 0 0 1⎟⎠ ⎝

(eq.4)

Next step is to transform the Raman tensor in the x’y’z’ reference frame. In this case, rotation is done along the z’-axis. The transformed Raman tensor R''(φ, θ) of the sample in the (x’y’z’) reference frame at an angle φ is given by

Rj ′′( φ ,θ ) = TxT ( φ )Rj ′( θ )Tx( φ )

(eq.5)

where, Tx(φ) is the x-component of the rotational matrix converting (x’y’z’) to (xyz) and expressed as

⎛1 0 0 ⎞ ⎟ ⎜ Tx ( φ ) = ⎜ 0 cos(φ ) −sin(φ )⎟ ⎜ 0 sin(φ ) cos(φ ) ⎟ ⎠ ⎝

(eq.6)

Consequently, the transformed scattering intensity I’ for the transformed Raman tensor in far-field takes the form similar to (eq.1) as

I ′ ∝ ∑ e R j ′′ ( φ , θ )e i T s

2

(eq.7)

j

Another factor, that we need to consider, is the change in amplitude of the electric field at the surface, such as air/ε-Si and strained Si/SiGe interfaces. The transmission and reflection

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coefficients for the p- and s-polarized light at different media of known refractive index can be easily be derived from Snell’s law and Fresnel equations given as [64]:

n 1sinφ 1 = n 2 sinφ 2

(eq.8)

n1 and n2 are refractive indexes. φand φ are the angle of incidence and refraction of light, respectively.

Tp =

n 2 cosφ 2 2 tp n 1cosφ 1

R p = rp

Ts =

2

n 2 cosφ 2 2 ts n 1cosφ 1

R s = rs

2

tp =

2n 1cosφ 1 n 2 cosφ 1 + n 1cosφ 2

(eq.9)

rp =

n 1cosφ 1 − n 2 cosφ 1 n 1cosφ 2 + n 2 cosφ 1

(eq.10)

ts =

2n1cosφ1 n1cosφ1 + n2cosφ 2

(eq.11)

rs =

n1cosφ 2 − n2cosφ 2 n1cosφ1 + n2cosφ 2

(eq.12)

Figure 14. Schematic of transmittance and reflectance ratio in ε-Si. Tp, Ts are transmission coefficient for p and s polarization at φ=60º. Rp, Rs are reflection coefficient for p and s polarization at φ=60º. The index of refraction n1, n2 and n3 are for air, ε-Si and SiGe, respectively. Index of refraction values were obtained in Ref.65,66.

The summary of the transmittance (Tp, Ts) and reflectance (Rp, Rs) obtained from the calculation are illustrated in Figure 14. Refractive indexes of n2 (Si) and n3 (SiGe) were quoted from reference [65,66]. However, the strain is gradually relaxed from the boundary of

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strained Si/SiGe to strained Si/air. Thus, the strain at the surface is minimum and is very close to ordinary silicon, which means that the refractive index of ε-Si approximately equal to unstrained Si. Detailed characterization of the refractive index can be analyzed by ellipsometry. It should be noted that when the incident angle φ1 is 60 degrees, the angle of incidence inside underneath the ε-Si layer, φ2, is 10.5 degrees. This fact should be taken into account when calculating the far-field Raman contributions inside the material. Taking into account all the factors above, the dependence of the far-field Raman intensities with sample azimuth was calculated at four combinations of polarizer and analyzer. The four combinations are P-P-θº, P-S-θº, S-P-θº, and S-S-θº, which represent the polarizer-analyzer-sample azimuth in degrees, respectively. In the next section, numerical results will be compared with SERS experimental data.

5.2. Sample Azimuth Dependence of Polarization Raman Measurement : Depolarization Effect in SERS The left most part in Figure 15(a)P-P-θº (b)P-S-θº, (c)S-P-θº and (d)S-S-θº shows the calculated normalized far-field Raman intensities as a function of sample azimuth relative to y-axis using c-silicon Raman tensor at different polarization states of polarizer and analyzer. Fig. 15(e)P-P-0º, (i)P-P-45º, (f)P-S-0º (j)P-S-45º, (g)S-P-0º, (k)S-P-45º, (h)S-S-0º, and (l)S-S45º are the experimental results of the far-field Raman and SERS spectra at 0 and 45 degrees sample azimuth, which indicate the minimum or maximum far-field Raman intensity for (a)PP-θº, (b)P-S-θº, (c)S-P-θº and (d)S-S-θº. All the spectra in Fig. 15 were obtained with the same exposure time of 10 minutes and laser power of 9.8 mW before the objective lens (NA:0.28) in a side illumination optics oriented at an angle φ1=60 degrees from the sample normal. The CCD dark counts are subtracted from all the spectra. For SERS spectroscopy, a 10 nm thick Ag film is evaporated at the rate of ~0.7Åm/s on the ε-Si surface under the ~10-6 Torr vacuum pressure. The SiGe Raman peak intensities at ~503 cm-1 in SERS (with Ag) are reduced to almost half of the far-field (without Ag) Raman intensity. This is because the illumination light has to transmit through the Ag films, in turn losing some probing power. Besides, no surface enhancement effect is expected to reach the underlying SiGe layers because the enhanced field decays strongly and is localized within ~20 nm from the surface [30,67]. From the calculation results, the amplitude of sinusoidal Raman intensities in (a)P-Pθº is to found to be greater than (b)P-S-θº, (c)S-P-θº and (d)S-S-θº. Fig. 15 (b)P-S-θº and (c)S-P-θº have the same amplitude. This amplitude of the Raman intensity is dependent on the excitation transmittance at the sample surface, e.g., the detection efficiency of P-P-θº is 67 % (0.82 x 0.82: incident x scattered), P-S-θºand S-P-θº is 28% (0.82 x 0.34=0.34 x 0.82), and S-S-θº is 12% (0.34 x 0.34). The detection efficiency was summarized in Fig. 14. In the P-P polarizer-analyzer setting in Fig. 15, the far-field Raman intensity level of (e)P-P-0º can be suppressed down to 1/4 when compared to the (i)P-P-45º where sample is rotated 45 degrees. This observation is consistent with the calculation result wherein the ratio of the maximum amplitude intensity over the minimum amplitude intensity is eight because the offset intensity marked in the red circle in Fig. 15(a) is ~0.12. Note that this discrepancy between experiments and calculations is because experimental results are obtained as an average of wide variety of incident angles using an objective lens. The offset of ~0.12, marked in the red

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circle in Fig. 15(a), is the sample azimuth where Raman intensity was obtained in (e)P-P-0º setting. However, in (j)P-S-45º, (k)S-P-45º and (h)S-S-0º cases, the far-field Raman intensity offset available for detections is suppressed down to almost zero. The reason why there exists a strong Raman background signals in (e)P-P-0º measured in side illumination, can be justified from the contributions of R1 (TO1), R2 (TO2), and R3 (LO) to the total detected Raman intensity. However, when measured by backscattering illumination optics with normal incidence (φ=0; normal incidence optics), only R3 (LO) becomes Raman active and can be totally suppressed by an analyzer [23]. On the other hand, at incident light angle of φ≠0 (side illumination optics), both R1 (TO1) and R2 (TO2) also becomes Raman active and contributes to the total Raman intensity because it’s not suppressed by the analyzer. In other words, Raman signal of (e)P-P-0º are mainly coming from TO phonon modes but not from the LO phonon modes.

Figure 15. (a-d) Calculated Raman intensity as a function of sample azimuth at a fixed polarizer and analyzer setting using the transformed silicon Raman tensor. (e-l) Comparison of far field and SERS Raman spectra at various combinations of polarizer, analyzer and sample azimuth. The black line spectra are far field Raman spectra, and the red line spectra are SERS spectra. Dashed lines are fitted Lorentz function spectra for Raman intensity analysis. SERS contrast value at each combination was obtained and indicated. (i) and (k) shaded with gray are the two conditions used in the TERS experiment. Red circles in the (a-d) indicate the sample angle at which minimum Raman intensity is observed.

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Strictly, the TO modes may not be overlapped with the stretched LO phonon mode. For example, in Fig. 15(e), Raman peak of ε-Si is observed at 511.1 cm-1, which is mainly due to the TO modes, while SERS peak of ε-Si is slightly shifted and observed at 509.4 cm-1, which is due to stretched LO mode. These facts suggest that the TO modes are considered as background signal for LO mode in our experiment and should be suppressed for higher sensitivity. Total intensities from TO and LO phonon modes are plotted in Fig. 15, (a)P-P-θº, (b)P-S-θº, (c)S-P-θº and (d)S-S-θº. To clearly understand the origin of background Raman signal, Figure 16 shows the calculated Raman intensities at each LO, TO1, and TO2 phonon mode in both normal and side illumination optics. Fig. 16(a) shows the normal incidence and Fig. 16(b), (c), and (d) correspond to P-P-θº, P-S-θº (or S-P-θº), and S-S-θº in the side illumination optics tilted at φ1=60 degrees incidence, respectively. When the incident angle is 0 degree (normal incidence), the minimum intensity of LO phonon mode is zero at the sample azimuth of 0 and multiples of 90 degrees (see inset in Fig. 16(a)), while the maximum intensity occurs at 45, 135, 225 and 315 degrees because the TO phonon mode is not active at all. However, in cases of (b) and (c), the TO phonon modes at the minimum LO phonon mode is not zero. This is because the two TO phonon modes are 90 degrees out of phase, making the sum a non-zero constant value at minimum LO phonon mode sample azimuth. Only in the case of (d), the TO phonon modes are not active. This suggests that selective LO phonon mode measurement is feasible. The issue is that (d)S-S-θº in Fig. 15 is not optimal situation for spectroscopy. In particular, for microscopic image acquisition the Raman signal level is extremely weak. According to the individual Raman intensities assigned to LO, TO1 and TO2 phonon modes shown in Fig. 16, the residual (offset) far-field Raman intensities of (e)P-P-0º, (j)P-S-45º and (k)S-P-45º, red circles in Fig. 15, are all due to TO phonon modes because LO phonon modes are completely rejected by the analyzer. On the contrary, far-field Raman intensity of S-S-0º should not have a contribution from the TO phonon modes, however, the far-field Raman spectrum in Fig. 15 (h)S-S-0º has the signal intensity similar to the (j)P-S-45º and (k)S-P-45º. This residual far-field Raman signal observed in (h)S-S-0º is experimentally due to the wide variation of the incident angle and depolarization effect of the objective lens itself; the background signal level becomes almost the same as (k)S-P-45º. In this sense, S-S0º is not advantageous over S-P-45º. Using this polarization dependence, we suggest a method to collect the near-field Raman signals with high efficiency in P-P-0º, P-S-45º, and SP-45º by tip-depolarization effect [59-61]. The far-field Raman signal that comes from the large focused spot is linearly polarized while the near-field Raman signal that comes from the proximity of the tip can be depolarized by the tip geometry. The depolarized near-field Raman signal can pass through the analyzer by choosing proper conditions of incident angle, polarizer, analyzer, and sample azimuth. Therefore, SERS spectra were studied first by coating 8nm thickness Ag film on surface of ε-Si before TERS. The depolarization effect in SERS occurs at the entire focused spot with the aid of roughened Ag thin films. In TERS, the depolarization is localized and takes place at the tip apex while in SERS, the depolarization occurs at the entire focus. Hence, the main difference in the depolarization effect is the scattering volume. The SERS and far field Raman spectra are shown in Fig. 15(e)P-P-0º, (i)PP-45º, (f)P-S-0º, (j)P-S-45º, (g)S-P-0º, (k)S-P-45º, (h)S-S-0º, and (l)S-S-45º which correspond to either minimum or maximum conditions for each polarizer-analyzer set.

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Figure 16. Raman intensity dependence of different Raman modes, LO (red) and TO (blue), as a function of sample azimuth at different polarizer and analyzer setting. (a) Normal incidence at P-P, P-S, S-P and S-S polarizer-analyzer polarization setting and side illumination (oblique incidence) (b)P-P, (c)P-S and S-P, (d)S-S polarizer-analyzer polarization configuration. Solid and dashed blue lines are Raman signal intensity of TO1 and TO2 phonon modes, respectively. Red solid line is Raman signal intensity from LO phonon mode. Black line is total Raman signal intensity. (inset: zoomed plot showing the zero Raman intensity)

Based on the SERS spectra, the Raman intensities of ε-Si is higher compared to their farfield Raman intensities at (e)P-P-0º, (j)P-S-45º, (k)S-P-45º, and (h)S-S-0º conditions. The reason is that, at these conditions, the far-field Raman is at minimum intensity as indicated by red circles in the calculation results of Fig. 15 and both depolarization and surface enhancement effect works efficiently with Ag film. To quantify the improvement, we calculated the parameter, contrast that is defined as the ratio of near-field signal intensity Inear to the far-field signal intensity Ifar expressed as [61]:

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I near ⎛ I total ⎞ =⎜ −1⎟⎟ I far ⎜⎝ I far ⎠

Itotal is the measured intensity that is the sum of far-field and near-field Raman intensities. Since the Raman signal from the sample consists of both the target ε-Si layer and the underlying SiGe layer, which generates strong Raman signal as background component, we focused on the relative intensity (I(ε-Si /SiGe)) between ε-Si and SiGe layer. In this case, the contrast can be expressed as:

⎛ ⎞ ε - Si ) ⎟ ⎞ ⎜ I SERS( ⎛I SiGe −1 Contrast= ⎜⎜ total −1⎟⎟ = ⎜ ⎟ ε ⎠ ⎜ I far ( - Si ) ⎝ I far ⎟ ⎠ ⎝ SiGe Since the Raman signal from the sample consists of both the target ε-Si layer and the underlying SiGe layer, which generates strong Raman signal as background component, we focused on the relative intensity (I(ε-Si /SiGe)) between ε-Si and SiGe layer. The Raman spectra were break down into Lorentz function spectra for peak intensity analysis. The total intensity for each peak, ε-Si and SiGe, were obtained from the area under the fitted Lorentz function spectrum. The calculated contrasts for each combination of polarizer, analyzer and sample azimuth are labeled in the experimental results of Fig. 15. The highest contrast (~3.7) is obtained in the (j)P-S-45º and (k)S-P-45º conditions. We can expect the same effect in TERS experiment. Following the SERS experiments, in order to show that the background suppression is more important than electromagnetic field enhancement for contrast improvement in Raman active and bulk crystalline materials such as silicon, we performed TERS measurement under the conditions in P-P-45º and S-P-45º (shaded with gray in Fig. 15). In the former case, P-P-45º, Raman signal level is strongest because tip-enhancement effect is expected in both for the incident and scattered light. In the latter case, S-P-45º, background signal is well suppressed due to the s-polarized incident light. Moreover, tipenhancement effect is not expected to happen in S-P-45º configuration because the incident light is s-polarized. However, tip-enhancement is expected for the scattered light. It is important to point out that the only difference between P-P-45º and S-P-45º is the incident polarization and the other parameters, analyzer and sample azimuth, are the same. These facts suggest that S-P-45º benefits on the tip-enhancement effect from the scattered light whereas P-P-45º utilizes tip-enhancement effect also from the incident light. This is because it is generally accepted that p-polarized illumination is better for the efficient incident tipenhancement effect [28,55]. Besides, the S-S-0º condition actually showed the best background suppression in the calculation, however, we exclude the use of S-S-0º in TERS because as discussed above the observed residual far-field Raman signal in S-S-0º is mainly due to the wide variation of the incident angle and depolarization effect caused by the objective lens itself. In this context, S-S-0º is not advantageous over S-P-45º.

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5.3. Experimental Observation of Depolarization Effect in TERS In order to compare contributions of field enhancement and depolarization effect included in P-P-45º and S-P-45º spectra, TERS measurements were carried out with and without Ag coated Si3N4 tip. In pure Si3N4 tip (without Ag), field enhancement is lower when compared to Ag coated Si3N4 tip that is expected to have higher field enhancement. The TERS result obtained from the Ag coated Si3N4 tip is shown in Figure 17. In the P-P-45º condition, the far-field Raman signals from LO phonon mode can efficiently pass through the analyzer. The decrease in the peak Raman intensity of SiGe when Ag coated tip is in contact can be attributed to the depolarization effect and the incident intensity loss caused by shadowing effect [28,61] in the tip. The influence of the shadowing due to tip insertion is almost negligible because tip diameter (~30nm) is much smaller than the area of an ellipsoidal focused spot (~2.6 x 1.5μm2) generated by LWD objective lens. The Raman intensity loss in SiGe associated with depolarization effect is the dominant factor because the Raman signal from LO phonon mode arising from the depolarized incident light, s-polarized in the P-P-45º case, is blocked by the analyzer. In order to accurately analyze the field enhancement and the depolarization effect, the discussion below are based on the following assumptions: (1) the diameter of the near-field scattering volume resulting from field enhancement is comparable to the tip diameter (~30nm) and the same as the thickness of the ε-Si layer (~30 nm), (2) the tip-enhancement factor for SiGe layer is negligible [68] because the thickness of ε-Si is 30 nm and set to 1x. Next, the 0.89x TERS depolarization effect was derived from simple intensity ratio between with Ag and without Ag coated tip in Fig. 17(a). This is because the reductions in SiGe Raman signal level is mainly attributed to the depolarization effect. These values are summarized in Fig. 17(c). On the other hand, the ε-Si Raman intensity exhibits a signal increment. Since depolarization effect was observed for SiGe layer, it should be also applicable to ε-Si layer in the same manner because both layers have the same facet (100). The depolarization effect factor on ε-Si is set to 0.89x in Fig. 17(c). The ratio of the total TERS intensity from ε-Si with Ag tip is 1.64 times higher than without Ag coated tip. The tip-enhancement factor is set to 1.84x in Fig. 17(c), and was calculated from 1.64/0.89. This fact supports that the tip actually enhanced the incident electric field intensity. Due to the existence of high background Raman signals from the entire focused spot, the estimated tip-enhancement contribution is small, ~1.84x. It should be noted that only the surface (ε-Si) is selectively enhanced and SiGe signal is reduced. This signifies that the result is not based on the signal enhancement from the multiple reflections, which is necessary in evaluating field enhancement on multilayered sample as discussed in C. Georgi et al, studies [69]. In the case of S-P-45º, the obtained depolarization effect value is 1.34x. The reason why the value turned-out to be more than one (signal is increased) is based on the fact that the polarized LO phonon (background) signals were rejected by the analyzer and the depolarized LO phonon signals by the tip passed through the analyzer efficiently. Using the value of the depolarization effect, tip-enhancement value can be also calculated in the same way we calculate for P-P-45º. Results are tabulated in Fig. 17(d). In S-P-45º, the tipenhancement is very large, ~11.7x, compared to P-P-45º, ~1.84x, despite the s-polarized incident polarization and p-polarized scattered light detection. In this case, tip-enhancement is expected to occur only on the Raman signal at the detection. These results strongly suggest that for bulk sample and Raman active crystals with well-defined crystalline polarization such

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as silicon, far-field (background) Raman signal suppression, which is likely to happen in spolarized illumination, is more important and should be carefully considered in improving the signal-to-noise ratio (SNR) in TERS experiment. For microscopic application and for faster data acquisitions, P-P-45º is recommended because of higher Raman signal level. The Raman signal from thin ε-Si layer was detected with much higher SNR in S-P-45º than the case in PP-45º in the presence of strong background signals from the thick underlying SiGe layer.

Figure 17. TERS (red) and far field (blue) Raman spectra with Ag coated silicon nitride (Si3N4) tip at (a)P-P-45º and (b)S-P-45º polarizer-analyzer-sample azimuth setting. (c) and (d) tabulates the calculated values for tip enhancement and depolarized effect contribution coming from P-P-45º and SP-45º setting, respectively.

To further demonstrate the importance of depolarization effect, TERS using pure Si3N4 tips (without Ag) was carried out and the results are shown in Figure 18. The far-field Raman spectra in Fig. 18(a) and (b) and the ones of Fig. 17 (a) and (b) are the same. The summary of enhancement factors induced by depolarization effect and tip-enhancement effect are listed in Fig. 18(c) and (d). Depolarization effect by the Si3N4 tip in P-P-45º case also caused reduction in the SiGe Raman intensity. This time, the depolarization effect is a bit higher, ~0.97x in Fig. 18(c) compared to 0.89x from the previous TERS experiment using Ag coated tip in Fig. 17(c). This could be due to the slight difference in the tip geometry between Ag coated Si3N4 tips and uncoated Si3N4 tips. On the other hand, the field enhancement value at the ε-Si is 1.24x, smaller than TERS with Ag coated tip. This means that tip-enhancement is still present even without the Ag coating. This can be due to the large contribution of an image dipole (the polarizability of image dipole: β=(εs-1)/(εs+1), εs: dielectric constant of material [70]) and efficient gap mode formed by a high refractive-index sample, in this study, ε-Si. Furthermore, in S-P-45º, the depolarization effect value (2.60x) is larger than Ag coated tip. This can be again due to difference in the tip geometry discussed above. The tipenhancement value is 2.61x and smaller than the case with Ag coated tip, but it is more than

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twice as high as Fig. 18 (a)P-P-45º with an uncoated tip. Even with uncoated Si3N4 tip and spolarization incidence, the ε-Si signal level got almost the same signal level as SiGe by the depolarization effect, and the SNR is much better than Fig. 17 (a) P-P-45º with a Ag coated tip.

Figure 18. TERS (red) and far field (blue) Raman spectra with pure silicon nitride (Si3N4) tip at (a)P-P45º and (b)S-P-45º polarizer-analyzer-sample azimuth setting. The tip enhancement and depolarized effect contributions in the P-P-45º and S-P-45º setting are listed in (c) and (d), respectively.

6. Conclusion We introduced the utilization SERS to selectively enhance the LO phonon vibrations originating from a thin silicon layer, which is under strain due to lattice mismatch with underlying Ge-doped silicon buffer layer. The selective enhancement makes it possible to efficiently detect the weak Raman signal originating from a thin layer, in the presence of a strong background signal originating from comparatively thick buffer layer. This selective enhancement could be localized to nanometric scale by utilizing TERS, where a metal-coated sharp tip is used for point-surface enhancement, which turned out to be an effective tool to investigate the nanometric variation of strain within the ε-Si layer. Further, we suggested some optimization and improvements in TERS for a higher sensitivity of the measurements. This included the selection of shorter wavelength for excitation, proper selection of tip material, and consideration of tip-originated effects, such as the tip-applied pressure that can add up in the local strain of the sample. In order to further improve the signal-to-noise ratio, we calculated the far-field Raman intensity of ε-Si (100) based on Raman tensor, Snell’s law, and Fresnel’s equations with the incident angle of 0 (normal incidence) and 60 degrees. Numerical results suggest that for normal incidence, only LO phonon mode is observed while in side illumination, incidence angle of 60 degrees, both TO and LO phonon modes are also

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observed. Because of TO phonon modes contributions, background Raman signal cannot be completely suppressed in cross-polarized configuration. However, LO phonon mode can be suppressed. SERS experiments suggest that P-P-45º gives the highest signal intensity and good for faster data acquisition in microscopic Raman studies. The P-S-45º and S-P-45º show the highest contrast between ε-Si and SiGe layers. In order to demonstrate the importance of background signal reduction for TERS, we compared experimentally the TERS spectra between P-P-45º and S-P-45º whose only difference is the incident polarization. Results show that S-P-45º condition strongly reduces the background signal but also reduces the tipenhancement effect. However, much higher signal contrast is observed between near field and far field compared to P-P-45º condition where high tip-enhancement effect is higher but also with strong far field background. Hence, for imaging Raman active and bulk crystalline materials such as silicon, background signal suppression (s-illumination) needs to be carefully considered as well as the field enhancement with strong far field Raman signal level (ppolarization). Accordingly, depending on the specific geometry of each tip, which affects both depolarization effect and field enhancement effect, the proper combination of s- and pilluminations show the highest contrast in the experiments while minimizing the background contributions from the sum of TO1, TO2, and LO modes through optimize analyzer and sample azimuth setting.

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In: Nanotechnology… Editors: C.J. Dixon and O.W. Curtines, pp. 107-150

ISBN: 978-1-60692-162-3 © 2010 Nova Science Publishers, Inc.

Chapter 4

NANOTECHNOLOGIES FOR CANCER DIAGNOSTICS AND TREATMENT Phong Tran1 and Thomas J. Webster2,* 1

Physics Department, Brown University, Providence, RI 02912 Division of Engineering and Department of Orthopedic Surgery, Brown University, Providence, RI 02912

2

Abstract Cancer treatment usually uses drugs (chemotherapy) to reduce tumor size, followed by surgery to remove the tumor (if possible). Then, more chemotherapy and radiation therapy is used to kill as many tumor cells as possible. The goal of this collective treatment is to target and kill cancerous tissue while minimizing side effects on healthy cells. Due to their non specificity, current cancer therapies have poor therapeutic efficacy and can also have severe side effects on normal tissues and cells. In addition, cancer is often diagnosed and treated too late, i.e., when the cancer cells have already invaded and metastasized (i.e. spread) to other parts of the body. At this stage, treatment methods are highly limited in their effectiveness. Thus, scientists have been focusing efforts into finding alternative methods to detect cancer at earlier stages and kill such cancerous tissues more effectively. Nanoparticles (that is, particles with at least one dimension less than 100 nm) have become very attractive for improving cancer diagnosis and treatment due to their novel optical, magnetic and structural properties not available in conventional (or micron) particles or bulk solids. Nanoparticles have been extensively studied for various applications including delivering anti-cancer drugs to tumorous tissues and/or enhancing imaging capabilities to better diagnose and treat cancer. In this review, recent work related to the improved targeted therapy for specific cancers (whether by developing more specific anti-cancer agents or by altering delivery methods) are summarized. Discussions on the advantages and disadvantages of the most widely studied nanoparticles (i.e., liposome nanoparticles, polymer-based nanoparticles, quantum dots, nanoshells, and superparamagnetic particles) in cancer imaging followed by anti-cancer drug delivery are highlighted. Lastly, bone cancer and current research in using nanoparticles for treating bone cancer, with an emphasis on the novel use of selenium (a natural anti-cancer element found in our bodies), are addressed.

* E-mail address: [email protected] (Webster TJ). Tel.: +1-401-863- 2318. (Corresponding author)

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1. Introduction Nanotechnology is an exciting multidisciplinary field that involves designing and engineering of materials and systems whose structures and components are less than 100 nm in at least one dimension [Ferrari 2005]. At this nanometer scale, the properties of objects (such as electrical, optical, mechanical, etc.) significantly differ from those at larger scales such as the micrometer scale [Klabunde 1996, Wu 1996, Baraton 1997 and Siegel 1995]. Having the ability to create new products with new characteristics and properties which were not available before, nanotechnology has shown great potential in a wide range of applications (such as information and communication technology [Heath 1998, Akyildiz 2008], biology [Giaever 2006, Saini 2006], and medicine [Silva 2004, Jain 2003, and West 2003], etc.). One of the most exciting contributions nanotechnology has made is in medicine and medical research to improve human health. Novel nanometer scale drug delivery systems and nano-tools have been developed to aid in disease detection of higher accuracy at earlier stages to treat such diseases more effectively [Park 2007, Orive 2005, Kong 2005, Jain 2003, and Shantesh 2006]. Of the tools that nanotechnology has created, nanoparticles have emerged over the last decade as the most promising for disease treatment and management, especially for diagnosing and treating cancer. Nanoparticles have unique properties that allow them to specifically target and deeply penetrate tumors. In addition, some nanoparticles (such as quantum dots, superparamagnetic nanoparticles) also have superior imaging capabilities allowing for a highly sensitive diagnosis of cancer [Shantesh 2006, Brannon-Peppas 2004, and Brigger 2002]. (See Table 1 for a summary of the advantages and disadvantages of the nanoparticles used in cancer diagnosis and treatment that will be discussed in the chapter). Among many kinds of cancer, bone cancer has been the subject of numerous studies due to its complexity. It is estimated that 2,380 individuals (1,270 men and 1,110 women) will be diagnosed with bone and joint cancers and 1,470 individuals will die from primary bone and joint cancers in 2008 in the United States [ACS 2008]. Primary bone cancer is rare as usually bone cancer is a result of the spread of cancer from other organs (such as the lungs, breasts and the prostate [Millier 2007]). Because many deaths are officially attributed to the original cancer source, the true numbers of bone cancer related deaths have been underreported. A common technique to treat bone cancer is the surgical removal of the cancerous tissue followed by insertion of an orthopedic implant to restore patient function. However, it is not always possible to remove all cancerous cells and therefore the remaining tumor cells can redevelop cancer. In these cases, it would be beneficial to have implants specifically designed to prevent the reoccurrence of bone cancer and simultaneously promote healthy bone tissue growth. This same argument can be made for numerous tissues (such as the lung, breast, etc.). This can be achieved by imparting conventional implant materials (such as titanium, stainless steel, etc.) with an anti-cancer chemistry. To promote healthy bone tissue growth, nano-feature surfacing can be used since studies have shown that osteoblast (bone-forming cells) function (including adhesion, proliferation, and the deposition of calcium containing minerals) is greater on nanostructured compared to current implant surfaces (which are micron-scale rough and nanoscale smooth) [Webster 1999, Webster 2000a, Webster 2000b, Webster 2001, Webster 2004, Perla 2005].

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This chapter will first give a brief introduction to the development and characteristics of cancerous tumors, different cancer targeting strategies, and then discuss the emerging roles of nanoparticles (including liposomes, quantum dots, nanoshells, superparamagnetic nanoparticles, and polymeric nanoparticles) in cancer imaging and therapeutics. Finally, this chapter will focus on the novel use of selenium nanoclusters as anti-cancer coatings for orthopedic applications. Table 1. Summary of advantages and disadvantages of the use of nanoparticles for cancer treatment to be discussed Type of nanoparticle Liposomes

Quantum dots

Advantages

Disadvantages

References

Biocompatible; biodegradable; non- Poor control over leakage of Torchilin 2005 immunogenic; amphiphilic; size, drugs; low encapsulation Tiwari 2006 charge and surface properties of efficacy; poor stability Soppimath 2001 liposomes can be easily changed during storage; poor Hans 2002 manufacturability at the industrial scale Fluorescently bright; large extinction coefficients; high quantum yields; absorption coefficients across a wide spectral range; highly resistant to photobleaching

Composition includes heavy Gao 2004a metals which are toxic Dubertret 2002 Ballou 2004 Reiss 2002 Niemeyer 2001 Alivisatos 1996

Fine-tunable optical response in a broad region of the spectrum from the near-UV to the mid-infrared; can be designed to strongly absorb or strongly scatter light in the NIR region; gold shell is biocompatible and rigid Superparamagnetic Controllable nano size; nanoparticles magnetic properties; easy to be directed using external magnetic fields; have controllable, specific Curie temperatures which allow for self regulated hyperthermia

Little-known fate following Loo 2004 introduction to human Loo 2005 bodies Oldenburg 1999 Hirsch 2003

Nanoshells

Polymeric nanoparticles

Hard to be directed to tumors which have a large distance to possible position of magnets; occlusion of blood vessels can occur in the target regions; possible toxic responses to human bodies Tailorability of polymer; Some preparation methods biocompatibility;degradability; use toxic organic solvents; various nanoparticle synthesis poor drug encapsulation for methods; versatility of drug loading certain hydrophilic drugs techniques; controllable drug and possibility of drug release characteristics leakage

Mornet 2004, Wang 2001, Jordan 1996, Wust 2002, Falk 2001, Kapp 2000, Gerner 2000, Alexiou 2000 Hans 2002 Anderson 1997 Brigger 2002 Soppimath 2001

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2. Growth and Other Characteristics of Tumors 2.1. Introduction to Tumors A tumor starts with a single or a couple of mutated cancerous cells surrounded by healthy, normal cells. As the cancerous cells replicate, due to its modified DNA, they generally develop at a higher speed than normal cells [Brannon-Peppas2004]. The term “in situ cancer” refers to the earliest detectable cancer cell mass [Rudon 1987]. In situ cancers are small lesions (a few millimeters in diameter) localized within normal tissues. These small tumors are avascular, i.e., they do not have their own network of blood vessels to supply oxygen and nutrients. Avascular tumors depend on diffusion of nutrition from pre-existing blood vessels and therefore grow slowly. As the cancerous cells develop further, surrounding normal tissue will not be able to compete for nutrition. Because of the insufficient nutrient supply, some tumor cells also die especially those located deep inside the tumors. Compared to cancer cells at the edge of the tumor, deep-located cancer cells rely entirely on diffusion to receive nutrients and to eliminate waste products. The cancerous cells will continue to duplicate and displace surrounding healthy cells until they reach a “diffusion-limited maximal size” where tumor cell growth rate is equal tumor cell death rate [Brannon-Peppas2004]. Unless there is more nutrients supplied from blood, tumors can not grow beyond this diffusion-limited maximal size (which is around 2 millimeters in most tumors [BrannonPeppas2004, Grossfeld 2002, and Jones 1998]). Tumors stay in this stage for years until they initiate the formation of blood vessels (the process known as angiogenesis). In order to continue growing, tumors need to be establish their own blood vessels through a process known as vascularization which is stimulated by angiogenesis factors (such as the vascular endothelial growth factor family, transforming growth factor β, ephrins, cadherin 5type 2, etc.). The vascularized tumors begin to grow more as nutrition supply is established. Clinically detectable tumors (approximately 109 cells, about one gram mass) can be found within a few months or years, depending on cell type [Ruddon 1987]. By that time, the tumor will have already gone through about 30 doublings, approximately two-thirds of its lifetime. If the tumor is not detected, a tumor mass of 32 gram can be achieved after five more population doublings (approximately 2 days for bone cancer cells (osteosarcoma)). More seriously, the tumor compresses surrounding tissues, invades basement membranes, and spread to other organs in the bodies (the process called metastasis) [Ruddon 1987]. This exponential growth of tumors and their abilities to metastasize are the reason why early cancer detection is extremely important.

2.2. Vascularization Process in Tumors (Angiogenesis) Angiogenesis is the process in which new blood vessels grow from pre-existing blood vessels. Angiogenesis can come from circulating endothelial precursors, shed from the vessel wall or mobilized from the bone marrow [Carmeliet 2000]. Angiogenesis starts with retraction of pericytes (vascular cells that wrap around blood capillaries) from the abluminal surface of capillaries. After that, endothelial cells release and activate proteases (such as urokinase (uPA), progelatinase A, progelatinase B, etc.) that degrade the extra-cellular matrix surrounding the existing capillaries. Endothelial cells then migrate, proliferate and align to

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form tube-like structures, which eventually join together to form new capillaries [Eatock 2000]. All of these stages can be the targets for cancer diagnosis and treatment.

2.3. Characteristics of Tumor Vascular Structures Blood vessels of tumors are very different from those of normal tissues [Carmeliet 2000]. Tumor vessels have uneven diameters, more branching and especially, a lot of openings. The reason for this irregular structure may be the imbalance of factors involved in angiogenesis, such as vascular endothelial growth factor (VEGF) and angiopoietins. Tumor vessels may also not have mechanisms to protect themselves from changes in oxygen, hormonal or metabolic balance. Abnormal vascular structure of tumors also have “chaotic” blood flow that results in insufficient blood supply in some areas. There areas, having not enough oxygen supplied from blood, become hypoxic (oxygen-starved) and acidic [Helmlinger 1997]. Hypoxic tumors are difficult to treat for traditional treatment such as drug delivery through blood stream or radiation therapy. Tumorous tissues not only have higher vascular density than normal tissues, but the structure of tumor vessels is also different [Carmeliet 2000]. For example, there are many “openings” on the walls of the vessels, the inter-endothelial junctions are wider, the basement membrane is discontinuous or absent, and the shapes of endothelial cells are also abnormal. Tumor vessels, therefore, are “leakier” than those of healthy, normal tissues. Rapid vascularization also leads to impaired lymphatic drainage systems in tumor tissue. The high vascular density of tumors, their leaky vessels and impaired lymphatic drainage in the tumor tissues creates the so-called “enhanced permeation and retention” (EPR) effects which the anti-cancer treatment methods can exploit [Sledge 2003, Teicher 2000].

3. Tumor Targeting Methods 3.1. Passive Targeting Passive targeting is a methodology that increases the target/non-target ratio of the amounts of delivered drugs primarily by minimizing non-specific interactions with non-target organs, tissues, and cells [Yokoyama 2005]. Nanoparticles are good candidates for this targeting strategy as they can be designed to have various size, surface charge, and hydrophobicity/hodrophilicity properties that subsequently affect the uptake of the particles by organs, tissues and cells. For example, particles with more hydrophobic surfaces (such as poly(propylene oxide), poly(methyl methacrylate), etc.) are taken up more by the liver than by the spleen and lungs [Brigger 2002]. To maximize circulation time and targeting ability, the nanoparticle size should be less than 100 nm (in diameter) and the surface should be hydrophilic to avoid clearance by macrophages. This goal can be achieved by coating nanoparticles with hydrophilic coatings (such as poly(ethylene glycol) (PEG), poloxamines, or poloxamers [Storm 1995, Brannon-Peppas 2004]. The enhanced permeability and retention (EPR) effect creates opportunities to increase the transport of drugs from blood vessels into tumor tissues and retain the drugs inside tumor tissues. Specifically, drugs can be designed to have sizes larger than the pore size of blood vessels in healthy tissues (from 2 nm-6 nm) but

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smaller than the pore size of blood vessels in tumor tissues (ranging from 100 nm to 780 nm) [Yuan 1995, Hobbs 1998]. The drugs therefore, will accumulate preferably at the tumor site and side effects of the drug on healthy tissue will be minimized. Moreover, due to the impaired drainage lymphatic system at tumor sites, the drugs can stay in the tumor longer, therefore increasing treatment effectiveness.

3.2. Active Targeting Active targeting is the method in which the therapeutic agent is delivered to tumors by attaching it to a ligand that binds to specific receptors that are over-expressed on target cells [Marcucci 2004]. Upon binding to the receptors on the target, the particle-ligand conjugates could be internalized into the cell (internalizing ligand) or only the particle is internalized (non-internalizing ligand) [Marcucci 2004]. The receptors that are over-expressed on target cells can be divided into the following groups [Marcucci 2004]: (i) receptors that are over expressed on endothelial cells in tumor blood vessels (such as integrin-vβ3 and negatively charged phospholipids) [Hood 2002, Ran 2002] and (ii) receptors that are over-expressed on tumor cells (such as HER2 and disialoganglioside (GD2)) [Park 2002, Pastorino 2003]).

Figure 1. EPR effects and targeting nanoparticles to the tumor. Nanoparticles with diameters between 50 nm and 100 nm can avoid entering normal tissue but can extravagate into tumors and target the tumor cells by ligand-receptor methods and release anti-cancer drugs to kill the tumor cells.

In active targeting strategy, sizes of drug carriers are also very critical. One needs the drugs to be transported by blood to the desired site (i.e. tumor) to be taken up by the tumor cells. As mentioned in section 3.1, blood vessels in most healthy tissues have a pore size from 2 nm to 6 nm while tumor vasculature has pore sizes ranging usually from 100 nm to 780 nm.

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Therefore, as illustrated in Figure 1, drug-nanoparticle conjugates of 50 nm to 100 nm size can enter the tumor from tumor blood vessels but cannot get into healthy tissue from healthy blood vessels [Marcucci 2004, Drummond 1999]. Figure 2 gives an example of doxorubicin-loaded liposomes with phospholipid-anchored folic acid–PEG conjugates (FTL-Dox) accumulated inside KB-HiFR tumor cells in vitro [Gabizon 2004]. Phospholipid-anchored folic acid is the ligand to target folate receptors over expressed on KB-HiFR tumor cells. The doxorubicin fluorescence (orange) is clearly seen in the nucleus.

Reprinted with permission from [Gabizon 2004].

Figure 2. Confocal fluorescence microscope picture of KB-HiFR tumor cells after two hours of in vitro exposure to 10 μM FTL-Dox.

Clearly, the above choice of nanoparticle size is part of passive targeting when one tries to avoid the uptake of drugs from healthy tissues. The passive component of drug targeting is important in active targeting systems and should not be over looked because the majority of a living body is non-target sites. For example, liver which is one of the largest drug targets comprises only 2% of the body weight. The rest of the body, which is 98% of the weight, can be considered to be a non-target site in this case [Yokoyama 1996]. Furthermore, before specific ligand-receptor interactions take place in active targeting, the drug-carrier-ligand complex has to be transferred through many other tissues since most targets are located in the extravascular space [Yokoyama 1996].

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After reaching the tumor site, the drugs will be internalized (i.e., enter tumor cells) with or without internalization of the carrier, i.e. nanoparticles. In the latter, nanoparticles will stay localized in the interstitium surrounding the tumor cells, while the drug, after being released from the drug-nanoparticle conjugate degradation, will enter tumor cells through diffusion or active transport [Marcucci 2004]. According to Marcucci and co-workers [Marcucci 2004], the factors that accelerate this degradation to release drugs are possibly the irregular condition of the tumor environment (such as acidic pH, enzymes and oxidizing agents [Drummond 1999]). Many methods have exploited these factors to design particulates that preferentially disintegrate at acidic pH or elevated temperature [Kirpotin 1996]. If the nanoparticles are also internalized along with the drugs, the internal environment of lysosomes will be the factors that disintegrate the nanoparticles and the drug can then diffuse out of the lysosomes. Between these two strategies of internalizing drugs, the methods that also internalize carrier particles have higher delivery efficacy.

4. Liposome Nanoparticles 4.1. Liposomes and their Advantages in Drug Delivery Liposomes are spherical vesicles formed by one or several concentric lipid bilayers with an aqueous phase inside and between the lipid bilayers (Figure 3A) [Torchilin]. Therefore, liposomes usually have hydrophilic outer surfaces, hydrophilic inner cores and hydrophobic matrices in between. Figure 3B gives an example of a transmission electron microscopy (TEM) image of liposomes (adapted with permission from [Ran 2003]).

Adapted with permission from [Ran 2003].

Figure 3. Schematic structure of a liposome with both water-soluble and water-insoluble encapsulated drugs (A) and a TEM image of the liposome (B).

Liposomes have several advantages in drug delivery for cancer treatment [Torchilin 2005]. They are highly biocompatible, biodegradable and non-immunogenic. These properties are attributed to their composition of naturally occurring lipids. Second, the size, charge and surface properties of liposomes can be easily changed by adding new ingredients to the lipid mixture before their preparation. For example, small unilamellar liposomes formed by a single bilayer can be around 100 nm in size, larger unilamellar vesicles have sizes ranging

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from 200 nm to 800 nm and multilamellar vesicles can be as large as 5,000 nm and consist of several concentric bilayers. Maybe the biggest advantage of using liposomes for drug delivery is their amphiphilicity (i.e., having both hydrophilic and hydrophobic properties) that allows for conjugations with both hydrophilic and hydrophobic therapeutic agents. Water-soluble (hydrophilic) drugs can be encapsulated in the aqueous core of liposomes and water-insoluble (hydrophobic) drugs can be entrapped within the lipid membrane [Torchilin 2005, Tiwari 2006]. Liposomes can also fuse with cell membranes and transfer drugs in the liposome to inside the cells [Chatterjee 2007]. There are two main mechanisms of liposomes releasing encapsulated drugs [Jesorka 2008]. One mechanism is based on the development of an affinity reaction [GómezHens 2006]. In this release mechanism, liposomes are fabricated to contain some component of affinity interaction (such as antigens, antibodies, enzymes, receptors) which interact with complementary affinity component at the targeted site and release the loaded drugs in situ. In the other mechanism, the triggered-release mechanism, liposomes are synthesized to contain some ingredients that are sensitive to external stimulus leading to changes in the membrane when an appropriate external stimulus is applied. Examples of this triggered-release liposomes include pH sensitive [Hafez 2000, Mizoue 2002, Reddy 2000, Yamada 2005], thermosensitive [Anyarambhatla 1999, Liu 2004], light sensitive [Benkoski 2006] or ultrasound sensitive[Huang 2004] liposomes. Membrane integrity of liposomes is lost under effects of the external stimulus leading to the release of entrapped compounds.

4.2. Passive Targeting Liposomes with PEG Coatings Liposome particles have several disadvantages when used as carriers for drug delivery. They are very quickly captured by the RES (half-life is less than 30 minutes). They are also instable, therefore, can not be used as a delivery agent without modification [Torchilin 2005]. PEG coatings, as seen in Figure 4, can prolong the circulation time of liposomes in the bloodstream (half-life can be increased to 5 hours [Klibanov 1990]) and PEG-coated liposomes can also be targeted to tumors by the EPR effect [Yokoyama 2005]. PEG is a hydrophilic, nonionic polymer commonly used as coatings on nanoparticles. PEG molecules have good biocompatibility properties and can be covalent bonded, absorbed to surface of particles or can be mixed with other ingredients to form particles during preparation process [Hans 2002]. It has been proposed that there are two main factors that affect the affinity of liposomes to the RES [Lee 1999]. Nonspecific hydrophobic interactions of liposomes with RES cells, and a specific opsonization reaction involving some blood component(s) such as immunoglobulin, complement proteins, apolipoproteins, and fetuin [Moghimi 2001]. Hydrophilic PEG coating can make liposomes become more hydrophilic therefore their nonspecific interactions with RES cells decrease. Specific interactions with opsonizing proteins are also reduced with PEG coatings because these coating layers may act as a brush shielding liposomes from the proteins [Drummond 1999]. More interestingly, coating liposomes with PEG can also be designed so that after the PEG coated liposomes reach the tumor site through EPR effect, the local conditions of tumor tissues (e.g. acidic pH in tumors) will make the PEG coating detached, promoting intracellular delivery of drugs [Immordino 2006].

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Figure 4. Schematic illustration of a PEG-coated, drug loaded liposome nanoparticle.

4.3. Active Targeting with Liposomes Besides passive targeting of liposomes to the solid tumor site as mentioned above, liposomes can also be conjugated with ligands for active targeting. One of the examples is the incorporation of a ligand of folate receptors (FR) to liposomes. FR is over-expressed on many cancer cells compared to their healthy, normal counterparts. For instance, the folate receptor is over-expressed in “ovarian (52 of 56 cases tested), endometrial (10 of 11), colorectal (6 of 27), breast (11 of 53), lung (6 of 18), renal cell (9 of 18) carcinomas, brain metastases derived from epithelial cancers (4 of 5), and neuroendocrine carcinomas (3 of 21)” [Sudimack 2000, Garin-Chesa 1993] In experiments by Pan and co-workers [Pan 2003], nanometer liposome particles with diameters of 100 nm, loaded with doxorubicin (DOX) and covalently attached with folate, were effectively delivered to KB xenograft tumors in mice and inhibited tumor growth, resulting in a 31% higher (p <0.01) increase in the lifespan of the tumor bearing mice compared to those that received liposomes loaded with DOX only.

4.4. Disadvantages of Liposomes Despite many advantages of liposomes, they still have some disadvantages in drug delivery applications [Bhojani 2007]. One of the disadvantages is leakage of the encapsulated drugs from liposomes into the blood before the drug-liposome conjugates can reach the targeted sites. Other limitations include low drug encapsulation efficacy and poor stability of liposomes during storage and manufacturability at the industrial scale [Soppimath 2001, Hans 2002].

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5. Quantum Dots 5.1. Properties of Quantum Dots Quantum dots (QDs) are nanocrystalline semiconductors consisted of an inorganic core (e.g., cadmium, mercury, cadmium selenite, etc.) with a metal shell (e.g., ZnS) that shields the core and provides biocompatibilities. QDs have diameters ranging from 2-10 nm (for coreshell QDs) and 5-20 nm after surface modifications (core-shell-conjugate QDs) (Figure 5A, 5B). QDs have unique optical and electrical properties (such as their fluorescence emission) and can be tuned from visible to infrared wavelengths depending on their size and composition. Especially, QDs can have large absorption coefficients across a wide range of spectrum and are very bright and photostable, making them ideal candidates for various imaging applications in vivo [Gao 2004a, Niemeyer 2001, Alivisatos 1996].

Adapted with permission from [Kim 2003].

Figure 5. Schematic quantum dot (QD) structures (A) and a TEM image of CdTe core - CdSe shell QDs in water (B).

The light-emitting properties of QDs are attributed to the quantum effects that are due to (i) their nanoscale structures and (ii) the quantum confinement phenomenon. Quantum confinement is a quantum effect in which the energy levels of a small nanocrystal (smaller than the Bohr exciton radius, about a few nanometers) are quantized with values directly related to the size of the crystal [Alivisatos 1996]. When exposed to light sources, the cores of QDs absorb incident photons and generate electron-hole pairs (characterized by a long lifetime, greater than 10 nanoseconds [Michalet 2005, Efros 2000]). The pair then recombines and emits a photon in a narrow, symmetric energy band (full width at half maximum is typically from 30 to 50 nm) [Michalet 2005]. The emitted light has wavelength ranging from 400 to 1350 nm (for QDs with sizes from 2 to 9.5 nm not including functional layers) [Michalet 2005]. In comparison with current organic fluorophores, QD nanoparticles have more desirable properties, therefore offering exciting new opportunities for in vivo imaging, especially in cancer diagnostics and management [Michalte 2005, Larson 2003]. First, fluorescence emitted by QDs is brighter than that of current organic dyes. Single QDs is from 10 to 20 times brighter than organic dyes [Gao 2004a]. QDs have very large

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extinction coefficients (around an order of magnitude larger than those of most dyes) that make them absorb more incoming light per unit concentration of dye [Michalte 2005, Dubertret 2002, Ballou 2004]. They also have high quantum yields (close to 90% [Reiss 2002, Bailey 2003]), meaning a high amount of light emitted over that absorbed. Second, QDs have properties that allow for multiple and simultaneous labeling of biological targets [Lee 2007]. Large absorption coefficients of QDs in a wide range of wavelength [Niemeyer 2001, Alivisatos 1996, Gao 2004a] enable one to simultaneously excite multiple QDs of different emissions with a single excitation wavelength. Furthermore, their emission spectra have very distinct and narrow wavelengths, therefore, allowing independent labeling and identification of numerous biological targets [Han 2001, Gao 2003, Gao 2004b]. This feature is very useful in studying tumor physiology which requires one be able to distinguish and monitor each component of the tumor microenvironment under dynamic conditions [Stroh 2005]. For example, researchers have used two-photon microscopy to image blood vessels within a tumorous tissue using PEG coated QDs [Stroh 2005]. As seen in Figure 6B, compared to Figure 6A, there was a good contrast between cells, matrix and the leaky vascular.

Reprinted with permission from [Stroh 2005].

Figure 6. (A) Concurrent imaging of both QDs with a 470 nm emission maximum and green fluorescent protein (GFP) (indicated in green) provides clear separation of the vessel from GFPexpressing perivascular cells and (B) and vessels highlighted with a deep red (QDs with a 660 nm emission maximum) micelle preparation were imaged simultaneously with the second harmonic generation signal (indicated in blue); the image represents a projection of a stack of 20 images at an interval of 2 micrometers per slice. Bars = 50µm.

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This experiment demonstrates the potential of QDs as a fluorescence contrast agent for noninvasive diagnostics and imaging of human tumors. Another advantage of QDs is that they have very high resistance to photobleaching. This feature is extremely important for in vivo imaging where fluorescent agents are constantly photobleached during image acquisition. For example, Dubertret and colleagues showed that, after 80 minutes of constant illumination, QDs fluorescence intensity remained unchanged while the dextran molecules were photobleached [Dubertret 2002].

5.2. Quantum Dots in Cancer Imaging and Treatment 5.2.1. Active and Passive Targeting for QDs Having superior light emitting properties, QDs are excellent candidates for tumor imaging if they can be effectively delivered to tumor sites. Both passive and active targeting mechanisms are being aggressively pursued in order to achieve this goal. Researchers have been working to increase circulation time of QDs in the blood and target QDs to cancerous tissues [Gao 2004a, Åkerman 2002, Rosenthal 2002]. Coating QDs with polymers (such as PEG) to avoid uptake by the RES, thereby, improving circulation time is an attractive approach being actively studied. For example, PEG coated CdSe-ZnS (core-shell) QDs have shown prolonged circulation in the mouse blood stream (half life more than 3 hours) compared to organic dyes which were cleared from circulation within minutes after injection [Gao 2005, Ballou 2004]. These PEG coated QDs were also shown to remain fluorescent after at least four months in vivo. It is believed that these PEG coated QDs are in an intermediate size range in which they are small and hydrophilic enough to reduce opsonization (that is, the alteration of a particle’s surface either by the attachment of complement proteins or antibodies specific for the antigen, so that the particle can be ingested (phagocytosed) by phagocytes, macrophages and/or neutrophils) and decrease uptake of RES, but they are still large enough so that renal filtration is minimized [Gao 2005]. Taking this design a step further, some researchers have conjugated PEG-coated QDs with ligands that can recognize cell membrane receptors (such as HER2, FR) that are overexpressed on cancer cells or on endothelia cells in cancer blood vessels. Specifically, a new class of QD conjugates containing an amphiphilic triblock copolymer for in vivo protection, targeting-ligands for tumor antigen recognition and multiple PEG molecules for improved biocompatibility and circulation have been developed [Gao 2004a]. In this study, researchers prepared CdSe-ZnS (Core-Shell) QDs encapsulated in a copolymer layer, tri-noctylphosphine oxide (TOPO), coated with PEG and conjugated to a prostate-specific membrane antigen monoclonal antibody (PSMA). PSMA is a cell surface marker for prostate epithelial cells and neovascular endothelial cells [Chang 2001] and has become a target for both imaging and targeted treating of prostate cancer [Gao 2005]. The QD conjugates have been successfully targeted to tumor sites in mice through both active and passive targeting mechanisms and have enhanced fluorescent imaging ability.

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5.2.2. QDs in Drug Delivery and Therapy for Cancers In addition to using QDs in imaging, there is a growing interest in using QDs for drug delivery and therapy. The size of QDs can be readily controlled, allowing for the evaluation of the size effect on the delivery efficiency and specificity, and therefore, the optimal size of drug carriers can be determined. High surface-to-volume ratios of QDs allow scientists to impart multiple functionalities on a single QD to develop a multi functional vehicle while keeping the overall size within the optimal range. For example, by coating a QD with an amphiphilic polymer layer (such as poly(maleic anhydride-alt-1-octadecene) (PMAO), octylamine-modified poly(acrylic acid), etc.), hydrophilic therapeutic agents and targeting biomolecules (such as antibodies, peptides), can be immobilized onto the hydrophilic side of the amphiphilic polymer and small molecule hydrophobic drugs can be entrapped in the matrix between the inorganic core and the amphiphilic polymer coating layer as illustrated in Figure 7A. Figure 7B shows an image of fixed breast cancer SK-BR-3 cells incubated with monoclonal anti-Her2 antibody (which was used to bind to the external domain of Her2) and goat anti-mouse IgG conjugated to QDs with an emission maximum at 630 nm (QD 630IgG). Her2 was clearly labeled with QD 630−IgG (shown in the red color).

Adapted with permission from [Wu 2002].

Figure 7. (A) Schematic of QDs fabricated to carry drugs and targeting ligand and (B) Fixed breast cancer SK-BR-3 cells incubated with monoclonal anti-Her2 antibody and goat anti-mouse IgG conjugated to QDs with an emission maximum at 630 nm (QD 630-IgG). Her2 was clearly labeled with QD 630−IgG (shown in the red color). The nuclei were counterstained with Hoechst 33342 (blue). Bar = 10 µm.

This integrated nanoparticle may serve as a multi-functional vector that not only specifically bind to and destroy tumor cells, but also emit detectable signals for real-time monitoring of its path.

5.3. Disadvantages of QDs The main disadvantage of QDs is due to their composition of heavy metals in the cores (e.g. Cd, Pb and Se which are known to be toxic to vertebrate systems at parts-per-million

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concentrations [Hardman 2006]) and the instability of uncoated QDs when exposed to UV radiation which leads to the release of heavy metal ions. For example, free cadmium in solution can bind to sulfhydryl groups of mitochondrial proteins [Wang 2008]. Thiol group inactivation then leads to oxidative stress and mitochondrial dysfunction [Notas 2006]. The instability of QDs under UV exposure is due to the fact that the energy of UV radiation is close to that of covalent chemical bonds [Smith 2006]. Therefore, UV exposure can dissolve the semiconductor particles through a photolysis process and toxic ions (such as cadmium ions) are released [Gao 2005]. For example, under exposure to air (30 min) or long exposure to UV (2–8 h), QDs at a concentration of 0.0625 mg/mL are already found to be extremely toxic [Ozkan 2004, Derfus 2004]. Generally, QD toxicity depends on the affects of multiple factors from both individual QD physicochemical properties and environmental conditions [Hardman 2006]. It has been shown that toxicity of QD depends on size, charge, concentration, outer coating bioactivity (i.e., capping material, functional groups, etc.), and oxidative, photolytic as well as mechanical stability of QDs [Hardman 2006]. For example, Derfus and co-workers found that the toxicity of CdSe QDs in liver cultures was influenced by different factors such as synthesis parameters, exposure to ultraviolet (UV) light and surface coatings [Derfus 2004]. They found CdSe (core only, no shell) QDs non cytotoxic under standard conditions of synthesis and water-solubilization with mercaptoacetic acid (MAA). However, TOPO-coated QDs which were first exposed to air for 30 min then modified with MAA, showed a dosedependent cytotoxicity toward primary hepatocyte cells (viability decreased from 98% to 21% at a QD concentration of 62.5 microgram/mL). They also coated CdSe QDs with layers of ZnS to study the effect of capping on cytotoxicity. They found ZnS capping eliminated cytotoxicity caused by air oxidation but ZnS capping did not eliminate cytotoxicity induced by UV exposure. To overcome the cytotoxicity problems, researchers have coated QDs with PEG or encapsulated them in micelles (that is, vesicles formed by the aggregation of surfactant molecules dispersed in a liquid colloid) which can limit the release of toxic metals induced by UV exposure [Cuenca 2006, Gao 2004a, Dubertret 2002, Stroh 2005, Gao 2005]. More work is still needed to develop optimal coatings and, thus, encapsulation that prevents the release of heavy ions from QDs.

6. Nanoshells 6.1. Structure of Nanoshells Metal nanoshells are spherical nanoparticles consisting of a dielectric core (such as silica) coated with a thin layer of metal (typically gold) (Figure 8). Gold nanoshells have optical properties similar to colloidal gold nanoparticles (such as strong optical absorption due to the collective electronic response of metal to light) [Loo 2004]. However, gold nanoshells show a strong dependence of optical response on the relative size of the core and the shell’s thickness. By varying the core size and the thickness of the gold layer, the color of these nanoparticles can be fine tuned across a wide spectral range spanning the visible and near infrared regions. Therefore, nanoshells show great promise in biomedical imaging and therapeutic applications in general and cancer imaging and treatment in particular [Loo 2005,

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Hirsch 2003, Chen 2001]. In addition, as with other types of nanoparticles, especially gold nanoparticles, drugs can be conjugated to nanoshells for drug delivery to tumors as seen in Figure 8.

Reprinted with permission from [Oldenburg 1998].

Figure 8. (Left) Schematic illustration of a nanoshell. Polymers (PEG) can be coated onto a nanoshell surface and drug molecules can be entrapped within the polymer matrix. (Middle and Right) (a)–(f) TEM images of nanoshell growth on 120 nm diameter silica dielectric nanoparticle. (a) Initial gold colloid-decorated silica nanoparticle. (b)–(e) Gradual growth and coalescence of gold colloid on silica nanoparticle surface. (f) Completed growth of metallic nanoshell.

The optical properties of nanoshells are attributed to the plasmon resonance at the dielectric-metal interface [Hirsch 2003]. Plasmon resonance is the phenomenon in which light induces collective oscillations of conductive metal electrons at the dielectric-metal interface. The absorbing and scattering properties of the particle will be influenced by the particle’s plasmon resonance. Depending on the relative thickness of the core and shell layers of a nanoshell, its plasmon resonance and hence, the optical absorption can be tuned across a wide range of the spectrum from the near-UV to the mid-infrared. Of particular interest in this spectral range is the near infrared region (NIR) [Oldenburg 1999] in which the transmittivity of tissues is high and adsorption is low. The NIR is very important in tissue imaging [Weissleder 2001]. The sensitive dependence of optical properties on the core diameter-shell thickness ratio can be understood by the hybridization model of Prodan et al. [Prodan 2003] in which plasmon resonance of a nanoshell is considered as the result of the interaction between the nanosphere plasmon and cavity plasmon which are “electromagnetic excitations that induce surface charges at the outer and inner interfaces of the metal shell”, respectively [Prodan 2003]. The sphere plasmon depends on the diameter of the sphere and the cavity plasmon depends on the inner and outer radius of the metallic shell [Aden 1951]. In addition, the

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interaction of the sphere and cavity plasmons depends on the thickness of the shell layer [Prodan 2003].

6.2. Optical Properties of Gold Nanoshells As previously stated, the optical response of nanoshells strongly depends on the relative size of the dielectric core and thickness of the metal layer. The ability to fine tune nanoshells to have optical resonance which varies over a broad region ranging from the near-UV to the mid-infrared is very important for tissue imaging. It has been demonstrated that nanoshells can be developed to highly scatter light within this spectral region as would be desired for many imaging applications; alternatively, nanoshells may be engineered to function as highly effective NIR absorbers permitting photothermal-based therapy applications as well [Loo 2004]; dual scattering/absorbing nanoshells can also be fabricated [Loo 2005]. For instance, the conventional NIR dye indocyanine green has an absorption cross section of 10-20 m2 at ~800 nm while the cross-section of absorbing nanoshells can be as high as 4 × 10-14 m2 [Hirsch 2003]. This comparison indicates that those nanoshells are over a million fold more (compared to the conventional dye) likely to encounter an absorbing event and convert that light into thermal energy which can be used to kill cancer cells or improve chemotherapy and radiotherapy efficiency [Hirsch 2003].

6.3. Nanoshells in Cancer Diagnostics and Treatment Gold nanoshells, having attractive optical properties as discussed above, offer promises for biomedical sensing and therapeutic applications, especially in cancer diagnosis and treatment [Loo 2005]. In addition to attractive optical properties, nanoshells also have rigid structures due to metallic shells, therefore offering improved stability. Therapeutic agents and other targeting molecules can be readily conjugated to the surface of nanoshells to specifically target and destroy cancer cells [Loo 2004]. To enhance biocompatibility and improve blood circulation, nanoshells are usually coated with “stealthing” polymers (such as PEG) even though the gold surfaces of nanoshells are generally considered to be biocompatible [Tang 1998]. Researchers have developed a novel approach to combine cancer diagnostics and therapeutics based on the use of gold nanoshells as near-infrared (NIR) absorbers. Loo and co-workers [Loo 2004] fabricated silica-gold nanoshells of 3 different sizes (core radius/shell thickness): 120nm/35nm, 100nm/20nm and 60nm/10nm. These nanoshells have strong optical scattering and absorption coefficients in the NIR region (wavelength around 800nm) and were used for imaging (by both dark field scattering and optical coherence tomography) and killing (by photothermal therapy) cancer cells (SKBr3 breast cancer cells). In order to specifically target cancer cells, anti-HER2 was conjugated onto the nanoshells to deliver the bioconjugates (i.e., anti-HER2- nanoshells) to SKBr3 breast adenocarcinoma cells which overexpress HER2. Compared to the control groups (cells cultured with nanoshells alone), cells cultured with “targeted” nanoshells were imaged with increased contrast. Cells were exposed to NIR irradiation (820 nm, 0.008 W/m2 for 7 min) to test photothermal therapy concepts. Cell death was observed only in the SKBr3 breast cancer cells incubated with anti-

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HER2 nanoshells. Cell death was not observed in cells treated with either nanoshells alone or exposed to NIR light alone. In vivo studies have also shown promising initial results for the use of gold nanoshells for cancer detection and treatment. In one study [Hirsch 2003], investigators fabricated nanoshells of 110 +/- 11-nm diameter core and a 10 nm thick gold shell. These gold nanoshells had a peak absorbance at 820 nm. A monolayer of PEG was assembled onto the gold nanoshell surfaces to stabilize them. The nanoshells were injected interstitially ~5 mm into a solid tumor volume in immunodeficiency (SCID) mice. From 5 to 30 minutes after nanoshell injection, tumors were exposed to NIR light (820 nm, 4 W/cm2, 5 mm spot diameter, <6 min). It was demonstrated that tumors injected with nanoshells had an average temperature elevation of ΔT = 37.4 ± 6.6º C within 4-6 minutes after light radiation. The increased temperature resulted in tumorous tissue damage indicated by coagulation, cell shrinkage, and loss of nuclear staining. In contrast, control cells which were injected with saline had significantly reduced average temperature increase after exposure to the same NIR treatment (ΔT < 10º C) and showed no tissue damage.

6.4. Disadvantages of Nanoshells Although nanoshells with attractive optical properties which can be readily fine tuned offer great promises in cancer detection and treatment, limitations still exist. The biggest disadvantage of nanoshells is probably their little-known fate following introduction to human bodies. These particles, having a very small size might present some potential undesired effects on the organs and tissues in our bodies. More intensive in vivo animal studies are currently being completed in order to investigate both applications and limitations of nanoshells in cancer imaging and treatment.

7. Superparamagnetic Nanoparticles Superparamagnetic nanoparticles (SPMNPs) have been extensively studied in biomedicine applications and, in particular, cancer diagnostics and treatment [Mornet 2004, Wang 2001, Jordan 1996, Wust 2002, Falk 2001, Kapp 2000, Gerner 2000, Alexiou 2000]. They have attractive properties that are desired in tumor imaging and treatment. First, they have a controllable nano size. Second, they possess magnetic properties that make them excellent candidates in magnetic-based imaging techniques and magnetic-based treatment methods. For example, SPMNPs can be conjugated with drugs, targeted to tumors using magnets. They can then absorb energy from a time-varying magnetic field to be heated up and result in the release of drugs while the heat can also make chemotherapy and radiation therapy more effective or can directly destroy tumor tissues. In this section, we will review the use of SPMNPs in magnetic resonance imaging (MRI), hyperthermia treatment and magnetic-based drug delivery.

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7.1. SPMNPs used as Magnetic Contrast Agents in MRI MRI works by the dual applications of an external magnetic field Bo (up to 2 Tesla) and a transverse (i.e. transverse to Bo) radio frequency (RF) pulse on protons which are present in large amounts in biological tissues (for example in water molecules, membrane lipids, proteins, etc.) [Mornet 2004]. Each proton has a small magnetic moment that can be aligned by applying an external magnetic field. The external magnetic field first applies to align a magnetic moment of each proton. The RF pulse is then directed to the region of interest and the protons absorb energy from the pulse and spin in a direction different than the direction of the external field. The RF pulse then will be turned off, the protons now begin to return to their original alignment with the external magnetic field and release excess stored energy through a process called relaxation. When this happens, a signal will be given off and the MRI system can detect this signal and process it to transform it into images of the tissue area of interest. Relaxation can be divided into two independent processes [Mornet 2004]: (i) longitudinal relaxation in which the longitudinal component of the proton’s magnetic moment returns to align with an external field Bo (called T1-recovery); and (ii) transverse relaxation, in which the transverse (i.e. transverse to Bo) component of the proton’s magnetic moment vanishes (called T2-recovery). Different tissues have different T1 and T2 values, therefore, they can be differentiated by MRI scanning. Both T1 and T2 can be shortened by the use of magnetic contrast agents. Magnetic contrast agents work by altering local magnetic moments in the region of interest, therefore, relaxation signals can be enhanced. For example, T2 can be shortened according to the equation (1) [Pankhurst 2003]: 1/T2* = 1/T2 + γ∆Bo/2

(1)

Where ∆Bo is the change in the magnetic field caused by variations in the local magnetic field inhomogenity or local magnetic susceptibility of the system due to the presence of magnetic enhance agents [Leach 1988, Browne 1999]. Figure 9 shows TEM images of Fe3O4 SPMNPs and schematically illustrates how SPMNPs can affect T2 recovery time. Compared to traditional MR contrast agents (such as gadolinium), SPMNPs have several advantages. First, SPMNPs are in the size range in which they can be delivered to the desired tissues more effectively. They are often composed of a superparamagnetic core composed of iron oxides such as magnetite Fe3O4, maghemite (γ-Fe2O3) or other insoluble ferrites (generally from 3 to 10 nm in diameter) with or without polymer coatings such as dextran [Sahoo 2003]. In particular, in the case of MRI for tumor imaging, SPMNPs can be delivered to tumors using passive targeting methods based on EPR effects. Second, they have much greater magnetic susceptibility (more than three orders of magnitude greater than traditional gadolinium superparamagnetic particles [Kraichman 2001]), meaning they are more responsive to an external magnetic field, than traditional MR contrast agents. For example, superparamagnetic iron oxides are effective as T2 relaxation enhancers due to the induction of local magnetic field inhomogenity in the surrounding tissues leading to a stronger shortening of the recovery time as in equation (1) [Pankhurst 2003].

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Adapted with permission from [Sun 2004].

Figure 9. TEM bright field images of (A) 6 nm and (B) 12 nm Fe3O4 SPMNPs deposited from their hexane dispersion on an amorphous carbon-coated copper grid and dried at room temperature. (C) Schematic illustration of SPMNPs bound to a tissue of interest, thereby, shortening T2 recovery time [Pankhurst 2003].

7.2. SPMNPs in Hyperthermia Treatment for Cancer Hyperthermia is a cancer treatment method in which temperatures of certain organs or tissues are raised to 41oC- 46oC. If the temperature is raised to above 47oC, tissue destruction occurs and this process is called thermoablation [Jordan 1999, Chen 2007]. Thermoablation is characterized by acute necrosis, coagulation or carbonization of tissue. It is clearly undesirable in a clinical situation due to systemic side effects and further clinical complications [Chen 2007]. Modern hyperthermia trials focus mainly on the optimization of thermal homogeneity at moderate temperatures (42oC- 43oC) in the target volume. SPMNPs have the potential to solve this optimizing problem. As it is well known, tumor cells are more susceptible to heat increases than healthy, normal cells [Chen 2007]. Cancer cells are more vulnerable to elevated temperatures compared to normal, healthy cells because [Chen 2007]: (i) normal cells locate close to normal blood streams which can dissipate heat more effectively compared to cancerous cells which have abnormal blood flow and (ii) cancerous cells have a more acidic surrounding environment, therefore, they are more susceptible to hyperthermia. Hyperthermia treatment for cancer works by raising the tumor temperature resulting in damages in the plasma membrane, the cytoskeleton and the cell nucleus [Shellman 2008]. In addition, some regulatory proteins, cytokine or kinases are affected by elevated temperature, which leads to changes in the cell cycle and can even induce apoptosis (that is, programmed, natural cell death driven by the cell regulatory system itself) [Chen 2007, Fairbairn 1995, Sellins 1991].

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There are several different methods of creating local heat increase (such as using microwave radiation, capacitive or inductive coupling of radiofrequency fields) by implanting electrodes, by ultrasound, or by lasers [Glockl 2006]. Magnetic hyperthermia uses the losses during the magnetization reversal process of the particles when exposed to alternating magnetic fields to convert such losses to heat. More specifically, SPMNPs are first directed and concentrated at the desired location using passive or active targeting methods or both. An alternating magnetic field is then applied resulting in the magnetization reversal process of the particles. Energy lost in this process is converted to heat. There are many advantages of using SPMNPs in hyperthermia. First, they can be effectively directed to tumor tissues using magnetic fields and hence can avoid heating up normal tissues. For example, in one in vivo study, superparamagnetic magnetite (Fe2O3/Fe3O4) particles (core size 3.1 +/- 0.7 nm) were coated with modified dextran and intratumouraly injected and delivered using magnets to C3H mammary carcinoma in the leg of mice. With an applied magnetic field of 50mT, the iron concentration at tumor increased 2.5 fold compared to the control in which no magnets were used for targeting [Jordan 1996]. Second, SPMNPs can be readily conjugated with therapeutic agents and delivered to tumors. Upon reaching the tumor site, an alternating magnetic field will be applied causing the particles to be heated, melting the thermosensitive coatings, leaking the drug at the tumor site and, at the same time, heating the tumor [Chen 2007]. This strategy has a dual effect on cancer treatment as researchers have shown increased chemotherapy efficacy when combined with hyperthermia [Wust 2002, Falk 2001, Kapp 2000, Gerner 2000]. Another advantage of SPMNPs in hyperthermia is that they have an appropriate and specific Curie temperature (TC) (i.e., the temperature above which SPMNPs lose their magnetic properties and, hence, their coupling with the external magnetic field) that limits the hyperthermia at a predetermined temperature. For example, PEG-coated copper nickel (Cu/Ni) alloy magnetic nanoparticles were synthesized to have a Curie temperature in the range of 43–46°C [Chatterjee 2005]. This range of Curie temperature of the Cu/Ni magnetic nanoparticles is very desirable in hyperthermia applications. Once the temperature of the magnetic nanoparticles increases to above TC (from 43-46oC in this case), the external magnetic field will stop interacting with the nanoparticles, therefore, further tissue heating will not occur.

7.3. Magnetic Targeting of SPMNP–Drug Conjugates A clear advantage of using SPMNPs is that they can be directed to tumor sites using an external magnetic field. Using this method, SPMNPs conjugated with anti-cancer drugs not only can be delivered specifically to tumors to release therapeutic agents but also the magnetic field can keep the conjugate localized in the tumor site leading to increased drug efficiency. For example, investigators have used 100 nm (hydrodynamic diameter) iron oxide SPMNPs bound to mitoxantrone (that is, an agent used in the treatment of certain types of cancer including metastasis breast cancer and acute myeloid leukemia) to target squamous cell carcinoma in rabbits [Alexiou 2000]. A magnetic field of up to 1.7 Tesla was used to attract SPMNPs to the tumor site. They observed a very high concentration of SPMNPs within the tumor after intra-arterial infusion. Importantly, this treatment resulted in a “complete and permanent remission” of the cancer cells compared with the control group (no treatment) with no signs of toxicity.

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In another example, investigators used a ferrofluid (i.e. colloidal solutions of iron oxide SPMNPs surrounded by anhydroglucose polymers) chemically bound to epirubicin (an antibiotic drug commonly used in chemotherapy) to treat cancer in animals [Lubbe 1996a]. The conjugating method allowed for a reversible drug-nanoparticle binding that enables drug desorption under physiological conditions of tumors (such as osmolality, pH, temperature). Their experiments showed for the first time, high tolerance and efficacy in mice and rats in which no LD50 (i.e. a dose at which 50% of subjects will die) could be found for the ferrofluid. This success led to clinical trials in humans. The ferrofluid was directed and concentrated in the tumor site using permanent magnets that generated a magnetic field of 0.8 T in tumors located near the body surface [Lubbe 1996b]. The investigators showed that the ferrofluid could be successfully targeted to the tumors in about half of the patients.

7.4. Disadvantages of SPMNPs In spite of the very exciting properties SPMNPs can offer, there are still several problems associated with using these nanoparticles in cancer diagnostics and therapeutics especially for magnetically targeted drug delivery. The first limitation is the use of magnets for attracting SPMNPs. When dealing with the tumors which have a large distance to possibly position the magnets, magnetic field strength at the tumor site will not be enough to attract the SPMNPs. Secondly, occlusion of blood vessels can occur in the target regions as due to high accumulation of magnetic particles [Timko 2005]. In addition, toxic responses of the body to the SPMNPs can also limit the application of such particles. All of these potential disadvantages need to be carefully considered.

8. Polymeric Nanoparticles Polymers have become increasingly attractive in formulating drug-loaded nanoparticles for treating cancers due to the tailorability of polymer, biocompatibility and biodegradability properties, various nanoparticle synthesis methods, versatility of drug loading techniques and controllable drug release characteristics. In this section, common fabrication methods will be briefly summarized followed by drug loading techniques to demonstrate how drug release characteristics of polymeric nanoparticles can be manipulated.

8.1. Polymeric Nanoparticle Preparation Methods Polymeric nanoparticles can be formulated from either preformed polymers or monomers by a variety of methods. However, the most popular methods to prepare polymeric nanoparticles start from biocompatible preformed polymers [Hans 2002].

8.1.1. Emulsification-Solvent Evaporation Method The polymer is dissolved in an organic solvent first, and then the drug is dissolved into the polymer solution to create a mixture of drugs dispersed in a polymer solution. The polymer solution is then added with water and a surfactant to create the emulsion. The

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organic solvent is finally evaporated and the polymeric nanoparticles are collected by centrifugation or lyophilization [Hans 2002, Torche 2000, Suh 1998, Song 1997, Cheng 1998, Feng 2001]. In this method, both water soluble and water non-soluble drugs can be incorporated into polymeric nanoparticles, however, hydrophilic drugs have showed a lower incorporating efficiency [Hans 2002].

8.1.2. Emulsification–Diffusion Method In this synthesis method, the solvent is chosen as a partially water soluble solvent (e.g. acetone). The polymer and drug are dissolved in the solvent and emulsified in the aqueous phase containing the stabilizer. The role of the stabilizer is to prevent emulsion droplets from aggregating. Then, water is added to the emulsion so that the solvent can diffuse into the water. The solution is stirred, resulting in the precipitation of nanoparticles which can be collected by centrifugation [Hans 2002, Kwon 2001, Takeuchi 2001]. The problem with this method is that water soluble drugs tend to leak out during the diffusion of the solvent [Hans 2002].

8.1.3. Nanoprecipitation Method In this method, the polymer and drugs are dissolved in solvents (such as acetone to prepare PLGAs) and followed by dispersing the mixture into a solution containing Pluronic F68 which is a difunctional block copolymer surfactant terminating in primary hydroxyl groups, a nonionic surfactant that is 100% active and relatively nontoxic. The solvent is then evaporated under reduced pressure leaving the nanoparticles remained in the suspension [Hans 2002].

8.1.4. Salting-Out Process The principle of salting out process is that solubility of a certain material is affected by concentration of salt in the solution. For example, proteins are less soluble in the more “salty” solution. In salting-out method to prepare polymeric nanoparticles, an emulsion is formed containing the polymer, acetone, magnesium acetate tetrahydrate, stabilizer, and the drugs [Hans 2002]. Then, water is added until the volume is sufficient to allow for diffusion of the acetone into the water, which results in the formation of nanoparticles. One limitation of this procedure is the incompatibilities of the salts with many bioactive compounds. As mentioned, the above methods are the most popular techniques to prepare nanoparticles from biocompatible, preformed polymers (such as poly(lactic-co-glycolic acid) (PLGA), poly(lactic acid), poly(hexadecylcyanoacrylate, etc.) [Hans 2002]. However, biodegradable nanoparticles can also be made from monomer polymerization methods [Sakuma 2001, Behan 2001]. Hydrophilic polysaccharides (like chitosan (CS)) can also be used to make polymeric nanoparticles by spontaneous ionic gelatin process [Vila 2002, Janes 2001]. This technique is becoming attractive because it does not use harmful organic solvents and it can formulate particles of small sizes and positive surface potentials [Hans 2002].

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8.2. Control the Properties of Polymeric Nanoparticles Along with various methods for forming polymeric nanoparticles is the firm control over the properties of the nanoparticles [Hans 2002]. First, the type of polymer and its concentration used in nanoparticle fabrications strongly affect the particle properties. The molecular weights of polymers can influence the particle size and drug loading capacity, specifically, smaller nanoparticles can be formulated from lower molecular weight polymers but also have lower drug encapsulation efficiency in general [Hans 2002]. Higher polymer concentrations and higher polymer molecular weights lead to higher encapsulation efficiency and larger sizes of the nanoparticles [Blanco 1997, Song 1997, Kwon 2001]. For example, Table 2 demonstrates the dependence of bovine serum albumin (BSA) capture efficiency on molecular weights and concentrations of PLGA (PLGA nanoparticles were prepared by emulsion-solvent evaporation). Table 2. Effect of PLGA molecular weight (MW) and concentration on bovine serum albumin (BSA) incorporation MW of PLGA 58000 58000 102000 102000

Concentration of PLGA (%) 3.0 6.0 3.0 6.0

Concentration of BSA (%) 10.0 14.0 10.0 14.0

Efficiency of BSA capture (%) 24.8 36.8 68.0 74.8

Adapted with permission from [Songa 1997].

As seen in Table 2, lower concentration of PLGA or lower molecular weight PLGA resulted in particles that had lower BSA entrapment than those made from higher molecular weight PLGA or higher PLGA concentration. Figure 10 shows the effect of PLGA concentrations on the size of PLGA nanoparticles prepared by emulsification–diffusion method using poly(vinyl alcohol) (PVA) as a stabilizer [Kwon 2001]. Larger size PLGA nanoparticles were formed with higher PLGA concentrations. Second, the type and amount of surfactant/stabilizer are also factors that affect properties of nanoparticles as well as drug encapsulation characteristics [Hans 2002]. For example, dipalmitoyl-phosphatidylcholine (DPPC), a phospholipid found in cell membrane, when used as an emulsifier to synthesis PLGA nanoparticles provided a more dense coating on the surface of the nanoparticles resulting in a smoother surface than particles made with traditional emulsifier, PVA [Feng 2001]. The drug loading efficiency of the PLGA nanoparticles made with DPPC also increased compared to the particles made with PVA using the emulsification solvent evaporation method. Another example is the case of PLGA nanoparticle preparations. It was shown that PLGA nanoparticle sizes were smaller when prepared using didodecyl dimethyl ammonium bromide (DMAB) than those prepared with PVA [Kwon 2001]. In another study, PLGA nanoparticles were formulated using PVA, chitosan or PVA-chitosan blend as different stabilizers [Ravi 2004]. It was shown (Figure 11) in this study that the sizes and morphologies of PLGA nanoparticles were different for different stabilizers.

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Reprinted with permission from [Kwon 2001].

Figure 10. Effect of PLGA concentration on the mean particle size of PLGA nanoparticles (at the concentration of 2.5 (weigh/volume %) of poly(vinyl alcohol) (PVA)).

Adapted with permission from [Ravi 2004].

Figure 11. Atomic force microscopy images of PLGA nanoparticles: (A) PLGA nanoparticles with PVA alone as a stabilizer, (B) with chitosan alone, and (C) with a PVA-chitosan blend. Bars represent 150 nm.

Third, the Zeta potential of polymeric nanoparticles can be controlled [Hans 2002]. Zeta potential is an electric potential that is directly related to, and an indication of, the total charge of a particle in a solution. Particles with more charges will have larger Zeta potential. Zeta potential has a strong influence on the stability of nanoparticles. Strong electronic interaction between charged particles with high Zeta potential (repulsive forces) will result in more stable particles (leading to less aggregation) and more uniform size distribution. Surface coatings (such as PEG) can make Zeta potential more negative (meaning, absolute value of Zeta potential increases), increasing the nanoparticle stability [Vila 2002].

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8.3. Drug Loading Methods There are many places and different methods therapeutic drugs can be loaded into polymeric nanoparticles. The drugs can be dispersed in a polymer matrix, entrapped in a nanoparticle core, surrounded by a polymer membrane, chemically bond to the polymer, or physically absorbed onto the particle’s surface [Hans 2002]. The method of drug loading can be divided into two classes [Soppimath 2001]: (i) incorporation of the drugs during the formulation of polymeric nanoparticles and (ii) absorption of the drugs into nanoparticles after the formulation of the particles. It was shown that the absorption method to load drugs after forming polymeric nanoparticles has a lower loading efficiency than incorporating the drugs during the process of forming nanoparticles [Alenso 1991, Ueda 1998]. In the incorporation method (i.e., drug is incorporated during the synthesis of nanoparticles), several factors can affect drug loading efficiency. The concentration of monomer and the concentration of drugs can both determine the drug entrapment efficiency. For example, the drug loading of doxorubicin into poly(butylcyanoacrylate) (PBCA) nanoparticles was shown to strongly depend on the concentration of doxorubicin [Yang 2000]. Increasing the concentration of doxorubicin from 5% to 20% resulted in a decrease in the entrapment efficiency from 24.9% to 10.5%. In the absorption method, the surface properties of nanoparticles (such as hydrophilicity) have a strong impact on drug loading. These surface properties, in turn, are functions of the properties of the polymers themselves and the methods used during preparation.

8.4. Drug Release Characteristics and Drug Biodistribution Profiles Since polymeric nanoparticles can be fabricated from various polymers by many different methods, the release characteristics of the nanoparticles loaded with drugs can be readily controlled [Bhojani 2007]. Prolonged release of drugs is desirable and can be achieved by changing polymer composition of the particles and/or their surface properties [Anderson 1997]. For example, by changing the composition of polymers, the degradation of nanoparticles varied from days to months [Brigger 2002]. For instance, PLGA nanoparticles were formulated from PLGA- monomethoxypoly(ethyleneglycol) (mPEG) copolymers of different compositions (i.e., different PLGA-mPEG molar ratios) and loaded with cisplatin (that is, a drug used in treatment of various types of cancers, including sarcomas, some carcinomas and lymphomas) [Avgoustakis 2002]. The composition of the PLGA–mPEG nanoparticles was shown to strongly affect the drug release profile, i.e. the rate of release increased at higher content of the mPEG in composition. The Zeta potentials and biodegradability of the PLGA-mPEG particles were also influenced by the content of mPEG in the copolymer. Not only does the polymeric composition of the nanoparticle determine the release profile of the drug, but properties of the drug itself (such as molecular weight, charge, localization in the nanospheres, drug loading method) can also strongly affect the drug distribution pattern in the RES [Brigger 2002, Couvreur 1980]. For example, it was also demonstrated that the incorporation of cytostatic drugs (drugs that stop cancer cells from multiplying in contrast to cytotoxic drugs that kill cancer cells) into polymeric nanoparticles resulted in a change in the drug distribution in animal bodies [Brigger 2002]. In one experiment in mice treated with

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doxorubicin–loaded poly(isohexylcyanoacrylate) (PIHCA) nanoparticles, more doxorubicin were found in the liver, spleen and lungs, as compared to the controls (i.e., mice treated with free doxorubicin) [Verdun 1990]. In another experiment, it was demonstrated that when actinomycin D was adsorbed on poly(methylcyanoacrylate) (PMCA) nanoparticles and injected to rats, most of the drug was found in the lungs of the rats [Brasseur 1980]. In contrast, actinomycin D incorporated into a more slowly biodegradable poly(ethylcyanoacrylate) (PECA) nanoparticles, the drug was found mainly in the small intestine of rats [Couvreur 1980]. Even with the same PECA nanoparticles, when other drug (vinblastine, an anti-cancer drug used to treat several kinds of cancer such as Hodgkin's lymphoma, non-small cell lung cancer, breast cancer and testicular cancer) was loaded, the drug highly concentrated in the spleen of rats [Couvreur 1980].

8.5. Disadvantages of Polymeric Nanoparticles Although polymeric nanoparticles have many advantages (such as various fabrication methods, versatile drug loading and drug release methods), they still have some drawbacks. The main limitation is that some preparation methods use toxic organic solvents that could degrade certain drugs when they come into contact during the formulation process or could be toxic to the environment as well as to the physiological system [Birnbaum 2000]. Other disadvantages include poor drug encapsulation for certain hydrophilic drugs and the possible drug leakage before reaching its target tissues and/or cells.

9. The Use of Selenium Nanoclusters as Anti-cancer Coatings for Orthopedic Applications 9.1. Selenium as a Chemopreventive Agent Natural materials, most notably, selenium are also being used to fight cancer. Selenium belongs to the group of metalloids, which are neither fully metals nor non-metals but share characteristics of both [Styblo 2001]. Selenium is naturally found in humans and animals and selenoproteins are important in antioxidant defense systems, thyroid hormone metabolism and redox control of cell reactions [McDowell 2003]. Animal studies have shown that selenium intake in excess of the nutritional requirements can inhibit and/or retard carcinogenesis [Combs 1986]. Moreover, studies have shown that people in areas of low soil selenium (bellow 0.05 ppm) and people with decreased plasma selenium levels (bellow 128 ng/ml) have higher cancer incidence and/or cancer mortality rates [Combs 1998, Clark 1991]. High levels of selenium in the blood (up to 112 µg/ml) have been correlated with reduced numbers of cancers including pancreatic, gastric, lung, nasopharyngeal, breast, uterine, respiratory, digestive, hematological and gynecological [Navarro-Alarcon 2000]. Clinical trials showed lower incidences of total non-skin cancer, cancers of the lung, colon-rectum cancer, and prostate cancer, as well as lower overall cancer mortality rates when a daily oral supplement of selenium-enriched yeast (200 μg selenium/day) was followed [Combs 1998]. The strongest evidence for the effect of selenium in reducing cancers was shown for lung cancer (46% lower incidence) [Clark 1996], esophageal and gastric-cardiac cancers [Wei 2004] and

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especially prostate cancer (63% lower incidence) [Clark 1998, Rayman 2007]. The mechanisms of selenium-based chemoprevention are complex and incompletely understood [Combs 2001].

9.2. Selenium Nanoclusters as a Promising Anticancer Coating Material for Orthopedic Implants 9.2.1. Method of Coating Selenium Nanoclusters on Orthopedic Implants The colloidal synthesis of selenium nano-particles is known and yields stable dispersions of free nano-particles in the presence of surface stabilizing agents [Kopeikin 2003, Lin 2005, Wang 2005]. These nano-particle suspensions have also been proven to have anti-cancer properties [Wang 2005]. In a previous study, nano-rough selenium was made by chemical etching from selenium compacts [Perla 2005, Tran 2008]. This mode of selenium addition, however, can limit the mechanical properties of the implant. Moreover, this methodology does not give control over the release and the dose of selenium. Considering that selenium is toxic at high doses [Whanger 1996], stability and control over its release would be a very desirable attribute. In the meantime, currently used orthopedic materials (such as titanium, stainless steel, etc.), on the one hand, show good mechanical properties to serve as bone graft materials, but, on the other hand, do not posses any inherent mechanisms to protect cancer from reoccurring. Therefore, it is may be a good idea to coat selenium nanoparticles on the current orthopedic material substrates. To achieve this, colloidal synthesis of selenium

Figure 12. SEM images of titanium substrates: (A) Pure Ti (denoted pTi in later figures); (B) Titanium with low Se coverage level (denoted Low-nSe-Ti in later figures); (C) Titanium with medium Se coverage level (denoted Medium-nSe-Ti in later figures) and (D) Titanium with high Se coverage level (denoted High-nSe-Ti in later figures).

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Figure 13. SEM images of stainless steel (SS) substrates: (A) Pure SS (denoted pSS in later figures); (B) Substrate with low Se coverage level (denoted Low-nSe-SS in later figures); (C) Substrate with medium Se coverage level (denoted Medium-nSe-SS in later figures) and (D) Substrate with high Se coverage level (denote High-nSe-SS in later figures). Bars = 1µm.

nanoparticles [Painter 1941, Ganther 1968] was carried out in the presence of common orthopedic metals: titanium and stainless steel. Briefly, the cleaned and sterilized substrates (that is, titanium and stainless steel) were placed in a beaker with the sides to be decorated facing upward. The reduced glutathione solution was added to the beaker followed by the sodium selenite solution. Different solution concentrations were used to achieve different doses. After a gentle mixing of the solutions in the reaction beaker, 1M NaOH was introduced to bring the pH into the alkaline regime in which selenium nanoparticles were released from the solution creating selenium nanoclusters formed on the substrates. Scanning electron microscopy (SEM) images of the titanium and stainless steel before and after coating showed selenium nanoclusters formed on the surfaces of substrates as seen in Figures 12 and 13. Different coating densities were also achieved by varying the concentration of synthesis solutions.

9.2.2. Surface Characteristics of the Titanium Substrates Coated with Selenium Nanoclusters In order to determine the state of selenium on the substrates, X ray photoelectron spectroscopy (XPS) was employed and the results (Figures 14 and 15) showed that selenium present on the titanium substrates is elemental selenium. This is important since this form of selenium (elemental) is believed to have the least toxicity among selenium and selenium compounds.

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Figure 14. XPS profiles of (A) pTi; (B) Low-nSe-Ti; (C) Medium-nSe-Ti and (D) High-nSe-Ti. The heights and numbers of peaks indicate relative amounts of selenium.

Figure 15. Peak at 55.2 eV indicates elemental state of selenium on the titanium substrates.

XPS data (Figure 14) also quantitatively confirmed the increasing amounts of selenium on titanium substrates as expected from increasing selenium concentrations in the synthesis solutions.

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9.2.3. Enhanced Non-Cancerous Osteoblast (Bone-Forming Cells) and Inhibited Cancerous Osteoblast Densities on Substrates Coated with Selenium Nanoclusters Nanomaterials have been shown to increase healthy bone tissue growth [Webster 1999, Webster 2000a, Webster 2000b, Webster 2004, Khang 2007]. For example, increased noncancerous osteoblast functions (from adhesion to proliferation and calcium containing mineral deposition) were demonstrated on nanostructured ceramic (such as alumina or titania) and metal (such as titanium, titanium alloys) surfaces compared to microstructured, nano-smooth surfaces [Webster 1999, Webster 2000a, Webster 2000b Webster 2001, Webster 2004, Perla 2005]. The mechanism of this increase is believed to be the increase in initial absorption of proteins that mediate subsequent osteoblast functions (such as fibronectin and vitronectin) [Webster 2000a, Webster 2001, Khang 2007]. A similar trend was observed for the titanium substrates coated with selenium (Figure 16). In particular, non-cancerous osteoblast densities increased on the coated titanium substrates after one day of in vitro culture compared to the uncoated titanium substrates. Although more investigation is needed, the nanofeatures of selenium nanoclusters are hypothesized to increase non-cancerous osteoblast cell densities.

Figure 16. Healthy osteoblast densities after four hours ( ) and one day ( ). Increased osteoblast densities on selenium coated titanium substrates after one day of culture. Data = mean ± SEM; N=3, * p<0.05, ** p<0.1 compared to pTi. (Compared at the same time period).

Importantly, after 3 days of culture, cancerous osteoblast densities on both titanium and stainless steel substrates coated with the highest amount of selenium were significantly reduced in comparison to all other substrates of the same type (Figure 17 and 18).

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Figure 17. Cancerous osteoblast densities decreased on selenium coated stainless steel substrates compared to the uncoated substrates. Data = mean +/- SEM. N=3. * p <0.01 compared to pSS.

Figure 18. Decreased cancerous osteoblast densities on selenium coated titanium substrates compared to the uncoated substrates. Data = mean +/- SEM. N=3. * p<0.01.

Selenium has been shown to have anti-cancer properties in the form of selenium compounds such as sodium selenite, selenium-sethylselenocysteine but not in the form of elemental selenium [Ip 1998]. It was also believed that elemental selenium has the least toxic properties than other compound forms of selenium [Ip 1998]. The above results showed, for the first time, that elemental selenium inhibited cancer osteoblast growth in vitro. The

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mechanisms of cancer-inhibiting properties of elemental selenium remain not fully understood and are still under active investigation. In the first step toward studying the mechanism of cancer-inhibiting properties of the selenium nanoclusters, the titanium substrates coated with selenium nanoclusters were immersed in cell culture media without the addition of serum for three days. The media was exchanged after 2 days and the spent media collected for graphite furnace atomic absorption spectroscopy (AAS). A dose dependent release of selenium from the substrates to the cell culture medium was observed as shown in Figure 19.

Figure 19. Total selenium released into cell culture media (

) after 2 days and (

) after 3 days.

Data = mean ± SEM; N=3, * p<0.0001 compared to High-nSe-Ti. (Compared at same time period); † p<0.0001 compared to Low-nSe-Ti. (Compared at same time period).

More interestingly, approximately 5.6% of the total selenium (estimated using SEM images) was released from the titanium coated with highest amount of selenium after 2 days. The release then slowed greatly, with a negligible selenium release occurring on the third day. Importantly, the concentration of selenium released into the culture media was much less than the concentration of 3-5 ppm that was demonstrated to show chemopreventive properties [Ip 1998]. Therefore, it was hypothesized that the attached cells experienced much higher selenium concentration in the vicinity of the coated surfaces. In a summary, studies on the fabrication and applications of selenium nanoclusters coated on orthopedic implant materials have yielded some preliminary results which showed anticancer properties of elemental selenium nanoclusters. More intensive research is underway to further study this promising material for a wide range of anticancer applications.

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10. Conclusions Over the last decade, a lot of effort has been focused on finding more sensitive diagnostic techniques, more targeted delivery methods as well as more effective therapeutic agents. Some advances have been used in the clinics but most of them are still under developmental stages. The ultimate goal is to lengthen the life of cancer patients by developing the tools that can detect cancer in early stages and treat them efficiently while keeping side effects minimal. There is undoubtedly a lot more that needs to be studied and improved in order to use even just some of the above-mentioned nanoparticles in a clinical setting. Understanding the advantages and disadvantages of each type of nanoparticle, researchers are focusing on making them more targeted and less harmful to healthy tissues while still delivering efficient therapeutic agents and imaging capabilities.

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In: Nanotechnology… Editors: C.J. Dixon and O.W. Curtines, pp. 151-236

ISBN: 978-1-60692-162-3 © 2010 Nova Science Publishers, Inc.

Chapter 5

MECHANICAL CHARACTERIZATION AT NANOMETRIC SCALE OF CERAMIC SUPERCONDUCTOR COMPOSITES J.J. Roa1, X.G. Capdevila and M. Segarra Departamento de Ciencia de los Materiales e Ingeniería Metalúrgica Facultad de Química, Universidad de Barcelona, C/ Martí i Franqués, 1; 08028, Barcelona, Spain

Abstract The nanoindentation or indenter testing technique (ITT) is a functional and fast technique that can give us a lot of information about the mechanical properties of different materials at nanometric scale, from soft materials, such as copper, to brittle materials, such as ceramics. The principle of the technique is the evaluation of the response of a material to an applied load. In a composite material, if the size of the residual imprint resulting from a certain load is lower than the size of the studied phase, then is possible to determine its mechanical properties, and therefore its contribution to the global mechanical properties of the composite. Depending on the tipped indenter used, different equations should be applied to study the response of the material and calculate stress-strain curves and parameters such as hardness, Young’s modulus, toughness, yield strength and shear stress. These equations are related to the different deformation mechanisms (elastic, plastic or elastoplastic) that the material undergoes. In the case of most of the ceramic composites, when a spherical tipped nanoindenter is used, elastic deformation takes place, and Hertz equations can be used to calculate the yield stress, shear stress and the strain-stress curves. On the other hand, when a Berckovich indenter is used, plastic deformation takes place, then Oliver and Pahrr equations must be applied to evaluate the hardness, Young’s modulus and toughness. Nevertheless, in the hardness study, Indentation Size Effect (ISE) must be considered. In this work, the mechanical properties of a ceramic superconductor material have been studied. YBa2Cu3O7-δ (YBCO or Y-123) textured by Bridgman and Top Seeding Melt Growth (TSMG) techniques have been obtained and their mechanical properties studied by ITT. This material presents a phase transition from tetragonal to orthorhombic that promotes a change in 1

E-mail address: [email protected].

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J.J. Roa, X.G. Capdevila and M. Segarra its electrical properties, from insulating to superconductor, and that can be achieved by partially oxygenating the material. On the other hand, the structure of the textured material is heterogeneous, and two different phases are present: a Y-123 as a matrix and Y2BaCuO5 (Y211) spherical inclusions. Moreover, the texture process induces an anisotropic structure, thus being the ab planes the ones that transport the superconductor properties while the c axis remains insulating. The purpose of this study is the characterization of the mechanical properties, in elastic and plastic range, of orthorhombic phases of YBCO samples textured by Bridgman and TSMG technique. With the ITT technique, the oxygenation process can be followed and its kinetics established.

1. Introduction High Temperature Superconducting (HTSC) materials can be use in a wide range of applications: I) II) III)

related to energy storage as flywheels, current transport devices as cables or fault current limiters, magnetic field related applications as squids, etc.

The magnetic properties of such materials are well known but the structural or mechanical properties that will allow the construction of real devices are not still completely described. The purpose of this work is to determine this set of parameters and characterize the material behaviour under determined stress conditions. This study will help us to define the optimal dimension of the material pieces in order to resist the stress suffered during its operating life.

1.1. High-Temperature Superconductors (HTSC) Since about 1962 it has become universally recognized that there exists a whole new class of superconductors, type II, which are characterized by the fact that they exhibit a new type of reversible magnetic behaviour. This discovery has made it possible to understand many of the previously unexplained superconducting properties of a number of elements and alloys. Furthermore, it has led to the recognition of the existence of a new thermodynamic state, the mixed state, which is only shown by type II superconductors. In addition to their intrinsic scientific interest they have a technological importance: niobium-zirconium alloys and the compound Nb3Sn have been used in the construction of superconducting solenoids capable of producing steady fields of 50 or 100 kOe. In 1986 Bednoz and Muller [1 and 2], discovered a new type of superconducting materials, now known as high temperature superconductors (HTSCs), which drastically improved the superconducting transition temperature (Tc). Further, YBa2Cu3O7-δ (YBCO or Y-123) was discovered by a group of researchers at the University of Houston in 1987 [3]. The Tc of YBCO material exceeded the boiling point of liquid nitrogen (92 K vs. 77 K). As a consequence, the superconducting products were expected to be operable using liquid nitrogen, which is cheaper and easier to handle than

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Helium; this is one of the most important advantages of HTSCs [4]. Moreover, the next year new superconducting materials were discovered in the Bi (Pb)-Sr-Ca-Cu-O system [5] with a Tc of 110 K and in the Tl-Ba-Ca-Cu-O system [6] with a Tc of 125 K. In 1993 HgBa2Ca2Cu3O9, with a Tc of 134 K was discovered [7], and at present, Tc has reached 164 K, under high pressure [8]. HTSC materials present structural features of ionic crystals. They contain CuO2 planes in their crystal structure. The layers between the CuO2 planes are called the charge reservoir layers. Features of the crystal structure of HTSCs include the following: a) Crystal structure is layered and the CuO2 layer and charge reservoir layer are stacked periodically. b) The parent material is the antiferromagnetic insulator. By doping electrons or holes to the CuO2 plane from the charge reservoir layer, the CuO2 plane becomes metallic and the superconductivity appears. c) At least one CuO2 plane, where the superconductivity current flows, must be included in the unit cell. d) When the superconductivity appears, the number of oxigens in the CuO2 plane for the number of copper ions is 0.15 to 0.20, and then the effect of the antiferromagnetism is strongly exhibited. The crystal structure, space group, and lattice constants of YBCO HTSC are summarized in Table 1. Table 1. Crystal data of YBCO HTSC [9] Material YBa2Cu3O7-δ (YBCO or Y-123)

Crystal Structure Orthorhombic

a (nm)

b (nm)

c (nm)

0.38177

0.38836

1.16827

Space Group Pmmm

HTSC have unique features such as a high upper critical field, high anisotropy, and an extremely short coherence length, ξ, in addition to the Tc being very high. The physical properties of YBCO are shown in Table 2. Melt and growth processes for producing Y-123 superconductive oxide [10, 11 and 12] are considered to be an effective process to obtain high critical current densities. Many investigations 13, 14 and 15] were performed to clarify the growth mechanism of Y-123 crystal from the partial molten state where Y2BaCuO5 (Y-211) and liquid phase coexist. Recently, it was found that the peritectic reaction of the Y-123 phase formation proceeded by the solute diffusion between Y-211 particles dispersed in the liquid phase and the growing Y123 interface and 16]. The growth models proposed assuming the mass transfer limiting were suggested, and the growth rate was found to be affected by an interface undercooling as the driving force for solute diffusion, [17 and 18]. Therefore, the undercooling is a principle parameter in controlling the growth of Y-123 single crystals. The inclusion of Y-211 particles has several advantages to the growth and properties of bulk YBCO superconductors: a) Preventing the liquid flow so as to decrease the amounts of holes.

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J.J. Roa, X.G. Capdevila and M. Segarra b) Shorten space between the Y-211 particles to help growing the Y-123 crystal as well as creating more Y-123/Y-211 boundaries which are the effective flux pinning site. c) Volume fraction of Y-211 particles is approximately constant; the current density, Jc, value is an inverse measure of the mean size of Y-211 phase [19 and 20]. High critical current density values can be achieved by controlling volume fraction and particle size of the Y-211 phase. Table 2. Physical Properties of YBCO Material name

YBCO or Y-123

Critical temperature, Tc (K)

92

Upper critical field, Bc2 (at 0K) [T]

674 (//ab)

Lower critical field, Bc1 (at 0K) [T]

0.025 (//ab)

122 (//c) 0.085 (//c) -3

Carrier concentration n [cm ]

1.5·1022

Coherence length, ξGL (at 0K) [nm]

1.15 (//ab)

Penetration depth, λ (at 0K) [nm]

142 (//ab)

0.15 (//c) >700 (//c)

Recent studies of YBCO materials pointed to two major issues: creating pinning centers and eliminating weak-lines between grain boundaries. It has been reported that some defects may act as pinning centers. Pinning strength and critical current density were increased by the introduction of fine Y-211 inclusions [21 and 22]. Flux pinning may be effective by two ways: first, the defects around the Y-211/Y-123 boundary, such as dislocations or stacking faults, and second, the magnetic pinning caused by the different induction generated in Y-123 superconducting matrix and Y-211 non-superconducting phase [23]. However, the weak-link can be caused by impurity phases, micro-cracks, or high-angle misalignment of the crystals [24]. Figure 1 shows the phseudoternary phase diagram of YO1.5-BaO-CuO system at 900ºC and oxygen partial pressure of 0.21 atm.]. There are four types of quaternary compounds – YBa2Cu3O7 (Y-123), Y2BaCuO5 (Y-211), Y2Ba8Cu6O18 (Y-143), YBa6Cu3O11 (Y-163) - and five types of ternary compounds – Y2BaO4 (Y-210), Y3Ba3O9 (Y-340), Ba2CuO3 (Y-021), BaCuO2 (Y-011), and Y2Cu2O5 (Y-101). Figure 2, shows the vertical cross-sectional diagram including Y-123 and Y-211 compounds]. The Y-123 phase is produced by the following ternary peritectic reaction at 1010 ºC: L+Y-211+Y-143 → Y-123

(1)

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Figure 1. Phseudoternary phase diagram of YO1.5-BaO-CuO system at 900ºC and oxygen partial pressure of 0.21 atm (From Chemical Processing of Ceramics, Second Edition, edited by Burtrand Lee, Sridhar komarneni; Chapter 23. Synthesis and Processing of High-Temperature Superconductors, page 604).

BaO:YO 1.5 1:5

mol% YO 1.5

BaO:CuO 3:5

Figure 2. Vertical cross-sectional diagram including Y-123 and Y-211 compounds in pseudoternary phase diagram (From Chemical Processing of Ceramics, Second Edition, edited by Burtrand Lee, Sridhar komarneni; Chapter 23. Synthesis and Processing of High-Temperature Superconductors, page 604).

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The peritectic transformation of Y-211 + L → Y-123 proceeds with the solute diffusion through liquid between Y-211 particles and Y-123 interface, and the undercooling acts as a driving force of the diffusion and]. All texturing techniques make use of the peritectic reaction where Y-123 phase grows from the Ba- and Cu-rich liquid (BaCuO2 + CuO) and the solid Y-211 phase [25]. From the viewpoint of solute diffusion, the large amount of Y-211 phase can supply more yttrium solute to the growing interface. During heating, the YBCO bulk to a high temperature, CuO reacts with Y-123 to form Y-211 and liquid. CuO + Y-123 Æ Y-211 + L

(2)

The solidification temperature of the liquid is lower than that of the peritectic liquid. Thus no Y-123 grain nucleates at the sample surfaces. Another way to prevent the surface nucleation is to coat the bulk surface with CeO2 powder. CeO2 was reported to react with barium of the Y-123 phase to form BaCeO3 [26] and the reaction is given by: CeO2 + Y-123 Æ BaCeO3 + Y-211 + CuO

(3)

As by-products, CuO and Y-211 form. The CuO reacts with Y-211 to produce lower melting point liquid in the same manner as the CuO addition, and hence surface nucleation is suppressed. The peritectic growth of Y-123 phase at typically low growth rates, coupled with the relatively large amount of highly reactive and viscous liquid phase generated during peritectic decomposition, makes it difficult to support bulk samples processed by directional solidification [27]. Melt processing technique of (RE)BCO generally involves slow solidification of a Y-211 and a Ba-Cu-O liquid phase mixture typically between 30 and 40 ºC below the Tp. A crystal that grows with faceted interfaces such as Y-123 crystal needs the driving force of interface kinetics for growth; i.e., it needs the saturation. The saturation or the undercooling kinetics, σ, can be re-written as:

σ=

(C

i

−C L ,Y −123 ) C L ,Y −123

(4)

where Ci and CL,Y-123 are the Y-123 concentration in the growth interface and in the liquid, respectively. The kinetic undercooling, is the difference between the chemical potential of a crystal and its surroundings. This concept is schematically shown in Figure 3. The growth of atomically flat face takes place by a step flow mechanism, the so-called lateral growth. Some probable mechanisms are schematically shown in Figure 4. The Y-211 particle is the yttrium source for Y-123 growth, and the yttrium concentration near the particle is higher than that near the Y-123 (Figure 3). This higher concentration near the Y-211 particle causes the concentration on the growing interface to be higher by approaching close to the interface and promotes the two-dimensional nucleation (Figure 4.a). The boundary between the Y-124 crystal and the Y-211 particle on the growing interface can act as the heterogeneous two-dimensional nucleation sites (Figure 4.b), and the entrapment of Y-211

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particle may field the misfit dislocations behind it (Figure 4.c). These nucleations and/or dislocations act as the step sources. These sites may become dominant when more Y-211 particles are entrapped. Therefore, the undercooling dependence of growth rate of the sample with excess Y-211 phase becomes linear. a)

b)

Figure 3. Schematic ilustration of (a) the yttrium concentration profile between the Y-211 particle and the Y-123 interface, and (b) the phase diagram showing the relation between the concentrations and the undercoolings (From Nakamura et al., J. Mater. Res., Vol.11, No. 5, May 1996). a) Y-211

b)

Y-123

Y-211

Y-123

c)

Y-211

Y-123

Figure 4. Schematic illustration showing some possible step sources on the Y-123 interface. a) A twodimensional nucleation resulting from higher concentration with approaching Y-211 close to the interface. b) Heterogeneous nucleation at the Y-211/Y-123 interface and c) A misfit dislocation caused by the entrapment of Y-211 particle. (From Nakamura et al., J. Mater. Res., Vol.11, No. 5, May 1996).

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1.1.1. Solidification and Microstructure of YBCO Bulk Materials Fabrication of single crystals usually involves the building of a structure, adding atoms layer by layer. Techniques to produce large monodomains included slowly dragging a rotating seed out of a molten bath of feeder material (as in Crochralski process and the Bridgman technique). Previously to the thermal treatment, we must to have a perform with shape and dimensions close to the final ones. These green bodies are achieved after pressing and sintering stages. The application of pressure varies within methods. It could be in uniaxial form or with an isostatic pressure field, etc. By the other hand, the sintering step could be realised previous to texturing process, such in the Bridgman technique or included in it, such in the TSMG method. Melt processing has been shown to be a suitable technique for the fabrication of bulk YBCO high-temperature superconductors with good flux pinning properties and high critical current densities [28, 29 and 30]. Several variations of melt texturing technique under a variety of names] have been developed based on this principle with a view to improving either features of the melt process or the quality of the product bulk material. The bulk textured material can be produced using the top-seeded Melt-Growth (TSMG) technique and the Bridgman technique with the purpose to obtain single crystals or monodomains. A monodomain can be described as a monocrystal (no grain boundaries) with numerous defects like secondary phases and mosaicity. Details on the texturation process can be found elsewhere [31]. The final microstructure in both techniques shows a homogeneous distribution of Y-211 particles in the textured Y-123 matrix. It is well-known that the Y-211 inclusions play an important role for the pinning of vortices].

Top-Seeded Melt-Growth (TSMG) Technique Top-seeded melt growth (TSMG) is known as the most effective process to fabricate block-type simples used for example in energy storage application, such a superconducting flywheel system, and others. In this technique, too long isothermal step and so great processing times is regarded as one of the most significant weak points. In the case of melt texturing with a temperature gradient, processing time is shorter, but the sample shape and size are restricted. The TSMG technique has been widely used to grow YBa2Cu3OX crystals as large as several centimeters [32], but these crystals were intended for applications such as magnetic levitation. The size of single-domain is generally about several centimeters in diameter, and limited to 10 cm for high quality YBCO bulk up to now, because of the grains mis-orientation during the melt growth process [33]. The resulting cubic centimeter-size YBCO crystals are further annealed to obtain the oxygen-ordered orthorombic phase (x = 6.5). Uniquely, the TSMG process yield large, single grains of approximately the dimensions of the green body [34, 35 and 36]. The TSMG technique has become the preferred method for the fabrication of bulk (RE)BCO superconductors and is used routinely in the processing of single-grain cylindrical/square shape samples of up to 50 mm in diameter [37]. TSMG processing is classified into two types by seeding method; cold seeding and hot seeding. Cold and hot seeding are named for the moment when the seed crystal is placed on the powder compact. In handling the sample, the cold seeding method is easier than hot

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seeding, because the seeding is performed at room temperature [38]. In this study the seeding method used is cold seeding. In order to obtain large-domain-sized YBCO and to control the growth orientation, seeding effects should be taken into consideration sufficiently [39]. The seed should not only have a similar structure and lattice constant to those of as-grown YBCO bulks, but also control the orientation and ensure single-domain growth [40]. The seed crystal initiates the nucleation and growth of the Y-123 phase in the incongruent melt, which subsequently solidifies into a single grain during controlled cooling. A variety of seed have so far been applied for the melt-textured (MT) growth of YBCO bulks], which can be classified into three major categories: • • •

Non-superconductors, such as MgO. Bulk superconductors, such as RE123 MT bulks or single crystal, such as Nd-123, Sm-123, and others. RE123 thin films [41].

In principle, the seed materials should not melt during the texture process, because the maximum process temperature is below their melting points. But the seed crystals were observed to dissolve frequently when they were in contact with the Ba-Cu-O liquid that was formed as a result of the incongruent melting of Y-123 compact during a high temperature holding period [42]. As the seed dissolve during processing, it no longer acts as a seed. In the case of dissolution formed, decreasing the levitation force and trapped magnetic field by reducing the size of the shield current loop. For this reason, the growth mode of YBCO grains is significantly dependent on seed thickness [43]. To obtain the green bodies, the well-mixed powder fabricated by Alcohol polyvinyl method (PVA method) with 1 wt. % CeO2, was uniaxially pressed into pellets with a diameter of 25 mm and thickness of 10-20 mm [44]. Stoichiometric quantities of the oxides and carbonates are weighted to give 69% w/w Y-123, 30 %w/w Y-211 and 1% w/w CeO2. This proportion has been demonstrated to maximize critical current density [45]. The NdBa2Cu3O7 (NdBCO or Nd-123) single crystals were manufactured with the Bridgman technique] with a temperature gradient 1-4 ºC/m in vertical direction. The (001) Nd-123 was placed on the top of sample. The Y-123 crystal grew from the Nd-123 seed crystal epitaxially with a squared sharp interface. The growth distance was defined as the length from the edge of the seed crystal to the solid-liquid interface. Figure 5 shows the heat treatment pattern followed in this study. The optimization of the temperature profile is crucial for the production of largesingle-domain YBCO monoliths. The TSMG YBCO samples were fabricated by the following process: The YBCO samples were heated up to 1055ºC at a rate of 300 ºC/h, and held for 2 h for homogenous melting. After that, the samples were cooled to about 1010 ºC at a rate of 30 ºC/h, and further cooled to 1008 ºC at a rate of about 0.02 ºC/h. Then the samples were cooled to room temperature at a rate of 300 ºC/h. Figure 6 shows YBCO samples with tetragonal structure obtained with the thermal treatment explained over (figure a, shows a single crystal and b, shows a poly-crystal). Finally, the as-grown samples were annealed at 450ºC for 300 h (In this step occurs the transformation from tetragonal to orthorhombic phase).

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Texture process Oxigenation process

Temperature, ºC

Y-211 + L TNucleation = 1010ºC ΔT

Tf,Nucleation = 1008 ºC TOxigenation = 450ºC

Time, h Figure 5. Heat pattern of the experimental procedure in the manufacture of YBCO samples by TSMG technique.

To define the degree of undercooling, the equilibrium peritectic temperature of Y-123 phase is needed. The temperature has been reported to be 1010 ± 10ºC by many researchers].

a)

b)

1 cm

1 cm

Figure 6. Photograph of the top view of the typical seeded grown sample, a) Single crystal and b) polysingle-crystal.

During the oxygenation process, cracks in the (a,b) planes appeared as a major drawback for the c-axis elements. The occurrence and propagation of cracks was found to be directly related to the oxygenation of the material. The oxygen uptake results in a decrease of the cell unit in the c-direction. The low oxygen diffusion rate yields large oxygen gradient during a

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classical annealing treatment. The cracks are then an easy diffusion path and the shrinkage of the unit cell along the c-axis at the crack tip provides the driving force for the cracks propagation. Ceramic superconductors exhibit an extensive defect zonology ranging from point-like defects of various natures, such as chain oxygen vacancies or anisette defects in some RE substituted 123 compounds, to a variety of extended defects. In addition, 123 compounds typically display ferroelectric transition which takes place well above room temperature and results in extensive twin formation. Dislocations, as well as their dissociated configurations are confined onto a prominent glide plane (001), although dislocation lines out of this plane may be frequently found building grain boundaries or as non-assembled segments that have climbed out the glide plane. This high anisotropy of dislocation configurations is consistent with the anisotropy of the crystal structure, the prominent glide plane corresponding to the largest lattice spacing. The melt processed ceramic composites contain a dense population of fine peritectic (non-reacted) particles, which drastically affects the microstructure acting as nucleation sites for dislocations and stacking faults as well as microcrak stoppers thus enhancing the mechanical toughness of the composite [46]. For TSMG techniques, under isostatic pressing conditions, plastic deformation is only possible if there is some kind of anisotropy in the material. In melt textured Y-123 composites, two kinds of anisotropy exist: elastic anisotropy [47] in the matrix; and plastic anisotropy between the peritectic inclusions and the matrix [48]. The main microstructural features are shown in figure 7a. The two dark areas in the left side and the upper right corner are Y-211 peritectic inclusions. Several dislocations oriented along (001) appear attached to the left hand side inclusions. It can be observed that the lower dislocations (indicated by A) is decorated with stacking faults appearing at every other twin domain. Similarly, stacking faults associated to the Y-211 particle located at the upper right corner of the micrograph (indicated by B) are observed to expend selectively on one twin domain. This distinction is only relevant to the orthorhombic lattice. Figures 7b and 7c are schematic drawings of the basic mechanism associated to the observed microstructural modifications. Another way to obtain green bodies used for TSMG is High Oxygen Pressure Processing. In this case, TEM observations have revealed a strong increase in the density of stacking faults nucleated at the interface of the Y-211 particles. Typical observations are presented in Figure 8a and 8b. The main distinguish features is the development of dendritic-like morphologies of the stacking faults.

Bridgman Technique The Bridgman technique is a method single crystal ingots or boules. It is a popular method of producing certain semiconductor crystals, such as gallium arsenide, II-V Crystals (ZnSe, CdS, CdTe) and BGO, where the Crochralski process is more difficult. The method involves heating polycrystalline material in a container above its melting point and slowly cooling it from one end where a seed crystal is located. Single crystal material is progressively formed along the length of the container. The process can be carried out in a horizontal or vertical geometry.

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In this study a modified Bridgman technique have been used. In this case, the sample was introduced into the furnace at one determined thermal gradient and a external Superconducting YBCO bars were prepared by a modified Bridgman process. After heating the presintered YBCO bars well above the peritectic temperature, Tp, the semisolid bars were displaced at a constant rate of 1 or 2 mm·h-1 through a region having an axial temperature gradient of 20 K·cm-1 at Tp until the full length of the bars was cooled down to 900ºC. Details of the sample preparation are given in [49]. The thermal treatment can be observed in figure 9.

a)

b)

c)

Figure 7. a) TEM micrograph close to the (001)plane of a Y-123 sample submitted to CIP, showing the selective expansion of stacking faults in a particular twin variant. Dislocation loops are indicated by arrows. The selective dissociation of a <100> dislocation is one twin variant is indicated by A. Selective reorganization of a partial loop orientated at the interface of a Y-211 particle is indicated by B, b) Schematic drawing of the reorganization of a partial loop driven by CuO fluxes from twin variant 2 to twin variant 1 as indicated by arrows, and c) Interaction of a gliding perfect <100> dislocations with a <E1-6><031> partial loop and resulting selective dissociation in twin variant 1. (From Sandiumenge et al., Advanced Materials, 12, No.5, 2000)

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a)

b)

Y-211

Figure 8. a) Typical TEM image of dendritic-like stacking faults generated at the interface of Y-211 particles. Imaging conditions are such that the stacking faults appear dark in order to emphasize their irregular shape. b) Enlarge view where the stacking fault contrast is avoided showing the high density of partial dislocation associated to the Y-211 interface (From Puig et al., Applied Physics Letters, Vol. 75, No. 13, 1999, 1952-1954).

Temperatue, ºC

T2 Tp = 1010 ºC T1

To to

Time, h

Figure 9. Thermal treatment of bulk superconductor samples textured by Bridgman technique.

164

J.J. Roa, X.G. Capdevila and M. Segarra The two main problems found when obtaining YBCO samples with this technique are: -

The first one is related to the different liquid properties. The optimal viscosity have to be found between minimize the sample flowing and maximize atomic diffusion. It is well known that the liquid phase generally migrates to the cold zone. In this case, the bar loss the correct stochiometry and in the end of the bar have a rich Y-211 zone.

After the texture process, the YBCO bars were introduced in an oxygenation furnace at 450ºC for 240 h. Figure 10 shows a schematic drawing and optical micrograph (OM) of YBCO bars textured by Bridgman technique. By means of polarized light microscopy an initial region can be observed where polynucleation and growth competition phenomena take place. This competition region has a length which is inversely proportional to the processing rate, thus indicativy that the nucleation is promoted by an enhanced undercooling [50]. In the single domain region the c-axis of Y-123 has a tilt angle of 45º respect to the long axis of the bars.

a)

Syngle-Crystal (45º)

b)

300 μm Figure 10. a) OM and b) schematic drawing showing from bottom to top, a multinucleation region, a growth competitive region and a single domain region of a Bridgman melt-grown sample. (Figure 10.b, from Ullrich et al., Materials Science and Engineering B, 53, 1998, page 143-148).

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Figure 11 shows a TEM micrograph of a Bridgman sample. In this picture, a high quality dislocations can be observed in the matrix of this material. The dislocation substructure of Y123 is highly anisotropic, being confined to the (001) plane. Lubenets et al. reported that strong covalent and ionic bonds create high Peierls barriers, which constrain the dislocation mobility in YBCO single crystals [51]. The micrograph presented in Figure 12 also shows lower amounts of dislocations placed preferentially in the grain boundary between Y-123 and Y-211. Moreover, the trajectories of the dislocations appear to be either unaffected or changed across the twin boundaries. On the other hand, this micrograph presented a higher amount of twins inside the precipitates. These effects could be due to two different factors: the compressive strain during the cooling treatment in the texture process and the different thermal expansion coefficients between the matrix and the particles (from 20ºC to 900ºC are 1.24·10-3 K-1 and 1.70·10-3 K-1 for Y-211 and Y-123, respectively]). Finally, we observe twins and residual stress inside the Y-211 particles. Stresses associated to the Y-211/Y-123 interface can arise from two different mechanisms. The first takes into account the thermal expansion and elastic modulus mismatch between the two phases].Secondly, stress is thought to result from the incorporation of Y-211 decomposition products into the matrix. Below the peritectic temperature, there is a thermodynamically driving force for solid-state dissolution of precipitates embedded in the matrix. Under uniaxially pressing conditions, plastic deformation is only possible if there is some kind of anisotropy in the material. In melt textured Y-123 composites, two kinds of anisotropy exist: elastic anisotropy] in the matrix, and plastic anisotropy between the peritectic inclusions and the matrix.

Dislocation

200 nm Figure 11. TEM micrograph of a Bridgman sample observed of a region containing dislocations in the maximum anisotropy plane (001 plane, ab).

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R es idua l S tres s T w ins

1 μm

Figure 12. TEM micrograph of Bridgman sample of a region containing several precipitates with different sizes. Note that several twins and residual stress produced during the texture process are present into the precipitates.

Dislocation

Residual Stress

0.5 μm Figure 13. TEM micrograph of a sample textured by Bridgman technique shows a precipitate with residual stress and dislocation. Both defects are due to the compressive strain produced during the texture process of material.

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In Figure 13, inside the Y-211 particle a dislocation can be observed, as well as in the middle of the particle a microcrack and a residual stress at the corners of the particle produced during the texture process. At the grain boundary between Y-211 and Y-123 phases, a fine dislocation line surrounding the particle can be observed. If we observe the matrix we will find a great amount of dislocations.

1.1.2. Oxygenation Process Although in the standard melt-grown YBCO microcracks also appear in the vicinity of Y211 particles due to thermal expansion difference of Y-211 and Y-123 [52], these microcraks do not propagate across the Y-211 precipitates, and the length of such microcraks is limited. The Y-211 particles are under compression when the material is cooled and can therefore break the Y-123 matrix along ab-planes [53]. These are the intrinsic microcraks, which are generated during the tetragonal-orthorombic phase transformation. Large secondary phase inclusions (unreacted liquid and Y-211) can also create macrocraks due to thermal expansion mismatch. Oxygen annealing, necessary to make Y-123 samples superconducting, is reported to be responsible for further macro and microcracking]. That is why oxygen concentration gradients lead to large mechanical stress in the material. The oxygenation diffusion coefficient in the ab-plane is about 104-106 times larger than in c-direction]. Figure 14 shows a micrograph of Y-123 observed with scanning transmission electron microscopy of a YBCO samples textured by Bridgman technique. We can observe a twin produced during the oxygenation process.

Twin

0.2 μm Figure 14. TEM micrograph of YBCO sample textured by Bridgman technique of a twin present in the matrix (Y-123) in the maximum anisotropy plane (001).

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Figure 15 shows a high resolution TEM (HRTEM) micrograph viewed along the grain boundary between Y-123 and Y-211 phases. In this micrograph we can observe, along the (001) plane, the transition from a well ordered region (Y-123) to a highly defective one (Y211). The distance between crystallographic planes is 7 Ǻ. We can also observe the high anisotropy in the Y-211 phase. In the case of growing the matrix (Y-123) on Yttrium substrates, the copper oxide rich liquid reacts with the substrate to form precipitates of Y-211 at the enriched interface [54], such that subsequent nucleation of the Y-211 liquid interface [55]. The mechanism would explain the anisotropic orientation distribution of Y-211 observed in figure 14. On the other hand, for Y-123 growing directionally from the melt, the particle engulfment process is governed by the velocity of the advantaging interface, melt viscosity and particle size. For anisotropic materials, such as Y-123 and Y-211 crystals, we can expect the above process to be governed by a complex relation, between orientation and size of the particle. These considerations suggest that the incorporation of Y-211 particles into the bulk Y-123 would be favoured for particular orientations if the growth rate is parallel to ab planes.

Y211 Y123/Y211

Y123 5 nm Figure 15. HRTEM image of a Y-211/Y-123 interface viewed along the [001] direction showing a transition from ordered zone (Y-123) to a highly defective zone (Y-211). Distance between cristalographic planes in the Y-123 zone is 7 Å.

YBCO pellets textured with TSMG and/or Bridgman technique have different cracking mechanisms

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a) Mechanical Stress Macrocraks in textured pellets appear during the final cooling stage to room temperature. Three major causes for large mechanical stresses, schematically illustrated in Figure 16, are considered at this stage: •

• •

An increasing Y-211 concentration is observed with increasing resistance from the seed in the c-growth sector while it remains relatively homogeneous in the ab-growth sector. Thermal expansion coefficients (in ab- and c- direction) of an Y-123/Y-211 composite are a function of the Y-211 content. The c-axis stress due to Y-211 inhomogenities described above can reach 50-100 MPa in the center of the pellet and will tend to open ab-plane macrocracks. During cooling a thermal gradient builds up with a colder surface and a hotter bulk due to relatively low thermal conductivity of Y-123. During to cooling to room temperature it takes up oxygen mostly in ab-direction and in vicinity of the surface.

a)

b)

c)

Figure 16. Schematic of mechanical (compressive or tensile) stresses associated with a) the Y-211 distribution into a pyramid pattern, b) thermal gradient during cooling and c) oxygen gradient ( From Isfort et al, Physica C 390, 2003, 341-355).

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b) Propagation of Cracks during Oxygenation [56] Cracking due to the oxygenation uptake of the surface does not seem to damage the sample severely, as the crazing is limited to the surface. The mechanism can be described as follows: In a first stage stress is building up at the exterior surface which creates the regular crazing pattern. Oxygen will then penetrate in the cracks and oxygenate their walls. The contraction in c-direction due to the oxygen uptake will then create a stress field around the crack tip that tends to make the crack progress inside the material (See Figure 17).

σσ σ σ σ σσ σσ YBCO σσσσσ O2

σσσ σσ σ σ σ σ σσ σ σ σ Tensión Stress

Figure 17. Schematic of mechanism of propagation of cracks during oxygenation (From J.J. Roa et al., Análes de Mecánica de la Fractura, 25, 2008, Spain, ISSN-0213-3725).

1.1.3. Applications At present, many applications using melt-textured YBCO are currently discussed [57]. The first demostrators, e.g. motors, flywheels, and others, have already been built. The HTSC capability to transport electric current without any losses, together with low thermal conductivity, suggest the application of HTSC for high current transport to low temperature SC devices such as magnets. The fault current limiter (FCL) seems to be the most promising superconducting power device that will be installed in the electric power networks. The higher levitation force of the YBCO bulk means that it can be used for various applications, such as non-contacted superconducting bearing [58], flywheel [59], magnetic levitation transport system and motors [60]. The application is mainly dependent on the physical properties of the YBCO bulk, such as levitation force and others. Although YBCO compound is one of the most widely studied superconducting materials, bulk YBCO superconductors are brittle and exhibit poor mechanical properties (strength and fracture toughness) [61]. Bulk textured Y-123 has to be considered as a brittle composite material due to the presence of micro-sized Y-211 inclusions in the Y-123 matrix. The most important problem of the superconductor materials is their poor mechanical properties. However, can become as important if one takes into account the stresses appearing in practical service due to the mechanical action caused by magnetic and/or thermal cycling

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between room and liquid nitrogen temperatures [ii]. These poor and unknown mechanical properties limit the performance of melt-textured YBCO in manufacture applications. The purpose of this study is the characterization of the mechanical properties, in elastic and plastic range, of orthorhombic phases of YBCO samples textured by Bridgman and TSMG technique. With the ITT technique, the oxygenation process can be followed and its kinetics established.

1.2. Indentation Testing Technique Indentation or hardness testing has long been used for characterization and quality control of materials, but the results are not absolute and depend on the test method. In general, traditional hardness test consist of the application of a single static force and corresponding well time with a specified tip shape and tip material, resulting in a hardness impression that has dimensions on the order of nanometers or micrometers depending on the applied load. The output of these hardness testers is typically a single indentation hardness value that is a measure of the relative penetration depth of the indentation tip into the sample. Actually, there are a lot of durometers which are used to characterize the mechanical resistance of materials, such as in polymer materials, that have different spring constants and either a flat conical tip, a sharp conical tip, or a spherical tip, as specified in ASTM D 2240, Standard Test Method for Rubber Property-Durometer Hardness. The deformation of materials occurs via two distinct processes: elastic (reversible) and plastic (irreversible) deformation. Since elastic formation is a reversible process, and is governed by angstrom scale (10-10 m), interaction parameters such as the crystallographic lattice constants, elastic deformation of materials exhibit virtually no size dependence unless a large population of preexisting deffects is involved [62]. The plastic deformation response, which occurs as a result of the generation, annihilation, and motion of deffects such as dislocations, displays marked size effects when those material dimensions are in the range of microns or below. Instrumented indentation testing (ITT), also known as depth-sensing indentation, continuous-recording indentation, ultra-low-load indentation, and nanoindentation, is a relatively new form of mechanical testing that significantly expands on the capabilities of traditional hardness testing. The past two decades, ITT employs high-resolution instrumentation to continuously control and monitor the loads and displacements of an indenter. The method was introduced in 1992 for measuring hardness and elastic modulus by instrumented ITT and has widely been adopted and used in the characterization of mechanical behaviour of materials at nanometric scales [63 and 64]. The principal advantage of this technique is that the mechanical properties can be determined directly from indentation load and displacement measurements or also from load-unload curves without the need to image the hardness impression. For this reason, this method has become a primary technique for the determination of the mechanical properties of thin films and small structural features [65, 66 and 67]. ITT provides information about composite materials when the particles in the matrix have a lower size than the residual indentation imprints. This technique, also, permits the study of the mechanical properties of monolithic samples . During the past decade, ITT has introduced several important changes to the method that both improve its accuracy and extend its real of application. The changes have been

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developed both through experience in testing a large number of materials and by improvements to testing equipment and techniques. Some of this changes were [68]: new methods for calibrating indenter area functions and load frame compliance, the measurement of contact stiffness by dynamic techniques allowing continuous measurement of properties as a function of depth, and others. A nanoindenter is a type of an indentation instrument tester, which means that load and total depth of penetration are measured as function of time during loading and unloading. Depending on the details of the specific testing system, loads as small as 1 nN can be applied, and displacements of 0.1 nm (1 Å) can be measured. Mechanical properties, such as: hardness, Young’s modulus, toughness, yield strength, shear stress,…, can be obtained with the load-displacement data. The technique most frequently employed measures the hardness, and the elastic modulus or Young’s modulus. [68 and 69]. This technique, also permits evaluating the yield stress and strain-hardening curves of metals [70], characteristic parameters of damping and internal friction in polymers, such as the storage and loss modulus and the activation energy and stress exponent for creep [71 and 72]. Mechanical properties are routinately measured from submicron indentations, and with careful technique, properties have even been determined from indentating only a few nanometers deep. Many ITT testing systems are equipped with automated specimen manipulation stages. In these systems, the spatial distribution of the near-surface mechanical properties can be mapped on a point-to-point basis along the surface in a fully automated way.

1.2.1. Testing Equipment (Instrumentation) Many instrumented indentation systems can be generalized in terms of the schematic illustration shown in figure 18 [73]. Load application device

Springs

Probe Tip Displacement Sensor

Sample Load frame

Figure 18. Schematic illustration of an instrumented indentation system.

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As shown in the figure, equipment for performing instrumented indentation tests consists of three components: a) an indenter of specific geometry usually mounted to a rigid column through which the force is transmitted, b) an actuator for applying the force, and c) a sensor for measuring the indenter displacements. However, to date, most ITT development has been performed using instruments specifically designed for small-scale work. Advances in instrumentation have been driven by technologies that demand accurate mechanical properties at the micron and submicron levels, such as the microelectronic and magnetic storage industries.

1.2.2. Nanoindenter’s Tips A variety of indenters made from a variety of materials are used in ITT testing. Diamond is probably the most frequently used material because its high hardness and elastic modulus minimize its contribution to the measured displacement from the indenter. Indenters can be made of other less-stiff materials, such as sapphire, tungsten carbide, or hardened steel. The indenters can be classificated in four different groups: •

Pyramidal indenters

The most frequently sharp indenter in nanoindentation technique is the Berkovich indenter. This indenter presents a three-sided pyramid with the same depth-to-area relation as the four-sided Vickers pyramid used commonly in microhardness work. With this indenter the hardness and the Young’s modulus can be determined. •

Spherical indenters

For spherical indenters, the contact stress is initially small and produces only elastic deformation. As the spherical indenter is driven into the surface, a transition from elastic to plastic deformation occurs, which can theoretically be used to examine yielding and work hardening, and to recreate the entire uniaxial stress-strain curve from data obtained in a single test [74 and 75]. At the micron scale, the use of spherical indenters has been impeded by difficulties in obtaining high-quality spheres made from hard, rigid materials. •

Cube-Corner Indenters

A three-sided pyramid with mutually perpendicular faces arranged in geometry like the corner of a cube. The center-line-to-face angle for this indenter is 34.3º whereas for the Berkovich indenter it is 65.3º. The sharper cube corner produces much higher stress and strains in the vicinity of the contact, which is useful, for example, in producing very small, well defined cracks around hardness impressions in brittle materials; such cracks can be used to estimate the fracture toughness at small scales [76]. Also, the toughness of brittle materials can be determined with a Berkovich indenter. •

Conical Indenters

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The conical indenter is also attractive because the complications associated with the stress concentrations at the sharp edges of the indenter are absent. Curiously, very little ITT testing has been conducted with cones. The main reason is that it is difficult to manufacture conical diamonds with sharp tips, making them of little use in nanoindentation technique [77]. The most important indenters used in nanoindentation technique are the Berkovich and Spherical indenter. It allows to characterize the plastic (hardness, Young’s modulus and toughness fracture) and the elastic (yield strength, mean contact pressure, shear stress and stress-strain curves) deformation. A typical Berkovich and Spherical indenter tip are shown in figure 19.

200 μm a)

500 μm b) Figure 19. High magnification SEM scan of the most important indenters used in nanoindentation technique, a) Berkovich indenter and b) Spherical indenter.

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1.2.3. Good Experimental Practice As in any experimental work, accurate measurements can be obtained only with good experimental technique and practice. The most important factors that can be controlled in a nanoindentation test are shown below: a) Choosing an Appropriate Indenter Choosing an appropriate indenter requires consideration of a number of factors. One consideration is the strain the tip imposes on the test material. Although the indentation process produces a complex strain field beneath the indenter, it has been proven to be useful to quantify the field with a single quantity, often termed the characteristic strain, ε There are problems, however, in obtaining accurate measurements of hardness and elastic modulus with cube-corner indenters [78]. Although not entirely understood, the problems appear to have two separate origins. First, as the angle of the indenter decreases, friction in the specimenindenter interface and its influence on the contact mechanics becomes increasingly important. Second, the relation among the contact stiffness, contact area and effective elastic modulus. Corrections are required, and the magnitude of the correction factor depends on the angle of the indenter. The spherical indenter can be used when one wishes to take advantage of the continuously changing strain. In principle, one can determine the elastic modulus, yield stress, and strain-hardening behaviour of a material all in one test. b) Environmental Control To take full advantage of the fine displacement resolution available in most ITT testing system, several precautions must be taken in choosing and preparing the testing environment. Uncertainties and errors in measured displacements arise from two separate environmental sources: vibration and variation in temperature that cause thermal expansion and contraction of the sample and testing system. To minimize vibration, testing systems should be located on quiet, solid foundation (ground floors) and mounted on vibration-isolation system. Thermal stability can be provided by enclosing the testing apparatus in an insulated cabinet to thermally buffer it from its surroundings and by controlling room temperature to within ± 0.5ºC. c) Surface Preparation Surface roughness is extremely important in instrumented indentation testing because the contact areas, from which mechanical properties are deduced, and calculated from the contact depth and area function, on the presumption that the surface roughness depends on the anticipated magnitude of the measured displacements, and the tolerance for uncertainty in the contact area. The greatest problems are encountered when the characteristic wavelength of the roughness is comparable to the contact diameter.

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d) Testing Procedure To avoid interference, successive indentations should be separated by at least 20 to 30 times the maximum depth when using a Berkovich or Vickers indenter. For other geometries, the rule is 7 to 10 times the maximum contact radius. The importance of frequently testing a standard material can not be over emphasized. e) Detecting the Surface One very important part of any good IIT testing procedure is accurate identification of the location of the specimen’s surface. This is especially important for any small contacts, in other words, when the applied load is very low, for which small errors in surface location can produce relatively large errors in penetration depth. For hard and stiff materials, such as hardened metals and ceramics, the load and/or contact stiffness, both of which increase upon contact, are often used. However, for soft, compliant materials, like polymers and biological tissues, the rate of increase in load and contact stiffness is often too small to allow for accurate surface identification. In these situations, a better method is sometimes offered by dynamic stiffness measurement [79and 80].

1.2.4. Experimental Techniques a) Hardness and Elastic Modulus Measurements The analysis of force-displacement or load-unload curves produced by instrumented indentation system is often based on work by Doerner and Nix [81] and Oliver and Pharr. The two mechanical properties measured most frequently by ITT methods are hardness and elastic modulus or Young’s modulus with a Berckovih indenter. For materials that do not experience pile-up, which includes most ceramics, hard materials, and soft metals that work harden, these mechanical properties can be determined generally within ± 10 %, sometimes better. The method was developed to measure the hardness and elastic modulus of a material from indentation load-displacement data obtained during one cycle of loading and unloading. Although it was originally indented for applications with sharp indenter, like Berckovich. A schematic representation of a typical data set obtained with a Berkovich indenter can be observed in figure 20, where P designates the load, and h the displacement relative to the initial undeformed surface. The load-displacement curve shows the elastic/plastic behaviour of each sample. From the difference between total indentation depth at maximum indented load (ht) and depth of residual impression upon loading (hf), the elastic recovery can be calculated [82]. In figure 20, there are four important quantities that must be measured: a) the maximum load, Pmax, b) the maximum displacement, hmax, c) the elastic unloading stiffness, S = dP/dh, defined as the slope of the upper portion of the unloading curve during the initial stages of unloading. The parameter S has the

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dimensions of force per unit distance, and is known as the elastic contact stiffness, or more simply, the contact stiffness, and d) the final depth, hf. The characteristics of figure 20 and such experimental results are summarized as follows: -

-

The initial load-displacement (P-h) response is elastic and can be described by continuum level contact mechanics. The first departure from this elastic response occurs when the local maximum shear stress level sustained by the indented material is on the order of the theoretical shear strength of material. Subsequent to this initial plastic event, a series of similar discontinuities in the P-h response occurs. Although the fundamental mechanisms responsible for the experimentally observed discrete deformation processes under this nanoscale contact are still debated in the literature.

Figure 20. Schematic illustration of indentation load-displacement data showing important measured parameters (From J. J. Roa et al., Nanotechnology, 18, 2007, 385701/1-385701/6).

The loading response of the material will show, therefore, whether if the indentation probe is blunt o sharp, producing an elastic initial response or an elasto-plastic response. Once the loading is sufficiently high, both indenters will produce similar elasto-plastic response in the material. It has to be taken into account that all real sharp indenters have a tip curvature which translates into a blunt indentation at the initial contact depths. The analysis of the P-h curves depends if it is a loading curve (which can be elastic or elasto-plastic) or a unloading curve, where the material is usually deformed. The accuracy of the mechanical properties in plastic deformation range (hardness and Young’s modulus) depends on how well these three parameters can be experimentally measured. The last parameter is the permanent depth of penetration after the indenter is fully unload; in other words, with this value the plastic work necessary to deform the material of the study can be calculated.

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The exact procedure used to measure H and E is based on the unloading process shown schematically in figure 21, in which it is assumed that the behaviour of the Berckovich indenter can be modelled by a conical indenter.

Figure 21. Schematic illustration of the unloading process showing parameters characterizing the contact geometry (From J. J. Roa et al., Nanotechnology, 18, 2007, 385701/1-385701/6).

The fundamental relations from which hardness and Effective Young’s modulus or reduced Young’s modulus are determined as follows:

H=

Pmax A

(5)

where Pmax is the maximum applied load and A is the projected contact area at that load. The hardness valued is a measure of the load-bearing capacity of the contact computed by dividing the applied load by the projected area of contact under load. This should not be confused with the more traditional definition of hardness, the load divided by the projected area of contact of the residual hardness impressions. These two different definitions of hardness yield similar values when the plastic deformation process dominates and a fully plastic permanent impression is formed. However, they give very different values when contact is predominantly elastic, because for purely elastic contact, the residual contact is vanishingly small, giving an infinite hardness based on the traditional definition. The Effective Young’s modulus can be determined as a function of S, A and a constant which depends on the geometry of the indenter.

E eff = f (S , A, β )

(6)

π S ⋅ 2⋅β A

(7)

Equation 6, can be re-written as:

E eff =

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Equation 7 is found in elastic contact theory and holds for any indenter that can be described as a body revolution of a smooth function Because this equation was derived for an axysimetric indenter, it formally applied only to circular contacts, for which the indenter geometry parameter is β = 1. However, it has been shown that the equation works equally well when the geometry is not axysimetmetric, provided that different values of β are used [83, 84 and 85]. For indenters with square cross sections like the Vickers pyramid, β = 1.012, for triangular cross sections like the Berkovich and the cube-corner indenters, β = 1.034 [iv]. This factor plays a very important role when accurate property measurements are desired. This constant affects not only the elastic modulus calculated from the contact stiffness by means of equation 7, but the hardness as well because procedures for determining the indenter area function are also based on equation 7, and area functions can be mistaken if the wrong value of β is used. An effective modulus, Eeff, is used in equation 7 to account for the fact that elastic displacements occur in both the indenter and the sample. The elastic modulus of the studied material, E, is obtained from Eeff, using the next equation:

1 1 − ν 2 1 − ν i2 = + E eff E Ei

(8)

where ν is the Poisson’s ratio for the test material, and E is the Young’s modulus. The subindex i denote the values of the indenter. For diamond indenter, the elastic constants Ei =1141 GPa and νi = 0.07 are often used. While it may seem counterintuitive that one must know the Poisson’s ratio of the studied material in order to calculate its Young’s modulus using the equation 8, even a rough estimate, say ν = 0.25 ± 0.1, produces only about a 5% uncertainty in the calculated value of E for most materials. The most important thing is to determine the contact stiffness and the contact area as well as possible. From equation 5 and 7, it is clear that, in order to calculate the hardness and elastic modulus of the studied material from indentation load-displacement curves, an accurate measurement of the S and the A must be performed. One of the main distinctions between IIT and conventional hardness testing is the way to obtain the contact area. Rather than by imaging, the area is established from an analysis of the indentation load-displacement data. The analysis used to determine the hardness, H, and the elastic modulus, E, is essentially an extension of the method proposed by Doener and Nix that accounts for the fact that unloading curves are distinctly curved in a manner that cannot be accounted for by that flat punch approximation. Doener and Nix consider that the contact area remains constant as the indenter is withdraw, and the resulting unloading curve is linear. The Oliver and Pahrr method begins by fitting the unloading portion of the load-displacement curve to the powerlaw relation:

P = B ⋅ A ⋅ (h − h f

)

m

(9)

where Β and m are power law fitting constants , and hf is the final displacement after complete unloading, also determined from the curve fit. The S is established by analytically

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differentiating equation 9; and evaluating the results at the maximum depth of penetration, h = hmax, that is:

S=

dP dh

(10) h = h max

Differentiating equation 9, the S can be obtained with the following equation:

S = B ⋅ m ⋅ (hmax − h f

)

m −1

(11)

It is thus prudent, when the S has been calculated with equation 11, to fit only the upper portion of the unloading curve; moreover, the value of S determined from this fit should be checked by comparing the curve fit to the data. Fitting the upper 25 to 50% of the data is usually sufficient. Now, the contact depth, hc, has been calculated, which for elastic contact is less than the total depth of penetration (hmax) as illustrated in the figure 21. The basic assumption is that the contact periphery sinks in a manner that can be described by models of indentation of a flat elastic half-space by rigid punches of simply geometry [86]. This assumption limits the applicability of the method because it does not account for the pile-up of material at the contact periphery that occurs on some elastic-plastic materials. Assuming, however, that pileup is negligible, the elastic models show that the amount of sink-in, hs, is given by the next equation:

hs = ε ⋅

Pmax S

(12)

where ε is a constant that depends on the geometry of the indenter. Typical values are: 0.72 for a conical punch indenter, 0.75 for a parabolic of revolution and 1 for flat punch. From the geometry of figure 21, the depth along contact is made between the indenter and the specimen, hc, can be calculated by:

hc = hmax − hs

(13)

Using equation 12 and 13, the contact between the indenter and the specimen is:

hc = hmax − ε ⋅

Pmax S

(14)

The projected contact area is calculated by evaluating an empirically determined indenter area as a function of A = f (d) at the contact depth hc; that can be re-writen as:

A = f (hc )

(15)

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The area function, also known as the shape function or tip function must be carefully calibrated by independent measurements, so that deviation from non-ideal indenter geometry are taken into account. These deviations can be quite severe near the tip of the Berkovich indenter, where some rounding inevitably occurs during the grinding process. b) Spherical Indentation (Determination of the Stress-Strain Curves) In the past 20 years, instrumented nanoindentation experiments have emerged as a powerful tool in understanding the mechanical behaviour of solids in general, and single crystals and thin films in particular. However, since the majority of the work has been carried out using Berkovich indenters, the emphasis has been put on extracting the hardness and the Young’s modulus of the material . Berkovich indenter is quite sharp and results in plastic deformation almost instantly; consequently, much of the information about the purely elastic region and, as important, the elastic to plastic transition is lost, a fact that has long been appreciate. Given that the conversion of load-displacement curves to indentation stress-strain curves is almost as old as the technique of using indentations to probe the mechanical properties of solids, it is somewhat surprising that this conversion is not much more common than it is. This comment notwithstanding, there have been a number of papers in which spherical nanoindenters have been used [87]. Roughly a decade ago, Field and Swain suggested a method to extract indentation stress-strain curves from load-displacement curves . Before acquiring the Continuous Stiffness Measurement (CSM) option, Field and Swain method was used to convert load-displacement results obtained on Ti3SiC2 [88], and single crystals of mica [89] and graphite [90], loaded parallel to the c-axis to indentation stressstrain curve. Typically, a nanoindentation test results in the load, P, and displacement into surface, ht, data. Additionally, the CSM attachment provides the harmonic contact stiffness, S, values over the entire range of loading. The elastic contact between two bodies was first described by Hertz in 1882, and the equations describing this response are usually known as Hertzian equations. Hertzian indentation has been applied to ceramic systems with heterogeneous microstructures, where an intermediate form of damage is observed. The vast majority of spherical nanoindentation analysis is based on the Hertz equation in the elastic region

P=

3 1 3 ⋅ E eff ⋅ R 2 ⋅ he 2 4

(16)

where R is the radius of the indenter, he is the elastic distance into the surface (see figure 22), and Eeff is the effective modulus given by equation 8. For a rigid spherical indenter, Sneddon showed that the elastic displacements of a plane surface above and below the contact circle are equal, and given by:

he = ht =

a3 R

(17)

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where a is the contact radius during the indentation, see figure 22. Combining equation 16 and 17, next equation can be obtained:

po =

3 ⎛a⎞ ⋅ Eeff ⋅ ⎜ ⎟ 4 ⋅π ⎝R⎠

(18)

The left side of equation 18 represents the indentation stress or mean contact pressure, also referred to as the Meyer hardness [91]. The expression in parentheses or a/R on the right side represents the indentation strain. Whith the spherical indenter two different mechanisms can be studied: the elastic and elasto-plastic mechanism.

Figure 22. Schematic representation of spherical indentation (From S. Basu et al., J. Mater. Res., 2006, Vol. 21, No. 10, page 2628-2637).

a) Elastic Regime Both the Oliver and Pharr and Field and Swain methods use the slopes of the initial portions of the unloading curves, dP/dh, to calculate he. Differentiating equation 16 with respect to h can yield: 3 1 dP = 2 ⋅ E eff ⋅ R 2 ⋅ he 2 dh

(19)

which when substituted this equation in equation 16, results in

P= Therefore,

2 dP ⋅ ⋅ he 3 dh

(20)

Mechanical Characterization at Nanometric Scale…

he =

3 dh ⋅P⋅ 2 dP

183

(21)

Since dP/dh is nothing but the stiffness, S*, of the system composed by the specimen and the load frame, the stiffness of the material itself can be calculated from next equation:

1 1 1 = * − S S Sf

(22)

where Sf is the load-frame stiffness. This value is obtained from the manufacturer of the instrument. b) Elasto-Plastic Regime Again following the Oliver and Pharr and Field and Swain methods, it can be assumed that the contact depth, hc, can be defined as the distance from the circle of contact to the maximum penetration depth, see figure 22, to be given by the next equation:

he 2

(23)

3 P ⋅ 4 S

(24)

hc ≈ ht − Combining equations 22 and 23 yields:

hc = ht −

Equation 24, can be modified and re-written as follows:

hc = ht −

3 P ⋅ +δ 4 S

(25)

where δ is an adjustable parameter of the order of a few nm needed to obtain the correct elastic moduli. Once hc is known, a can be calculated with the next equation:

a = 2 ⋅ R ⋅ hc − hc2 ≈ 2 ⋅ R ⋅ hc

(26)

Equation 26 is only valid for hc <<< a, and when the indenter tip is perfectly spherical. In the purely elastic regime, hc = ht/2 = he/2 and equations 17 and 23 become identical. Note that for the most part of the plastic regime, since ht >> he/2 it follows that hc ≈ ht (equation 23)

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c) Fracture Toughness Measurement Ceramics are generally brittle and prone to generation of cracks when indented. Toughness estimation by microfracture Vickers indentation is a well-known and broadly employed technique in ceramic materials. This technique consist of the application of a Vickers or Berkovich indenter at a given load to the material sufficiently high to nucleate cracks at the corners of the imprint, and further measure the crack lengths produced at the corners of the imprint, c, in order to evaluate the fracture toughness of the studied material. The fracture toughness of brittle bulk materials, an important measure of the resistance of these materials to fracture and crack propagation, can be evaluated through conventional microindentation using a microindenter. However, because of the bluntness of the microindenter tip, large forces are necessary to produce the cracks needed for analysis. Hardness and modulus of thin films or particles in a matrix can not be measured with microindentation without tedious work in correcting the substrate effect or the contribution of the matrix on the results. So, the fracture toughness of thin film or ceramic composite, such as YBCO, cannot be easily determined with the microindentation method. Fracture toughness at small scales can be measured by ultra-low load indentation using techniques similar to those developed for microindentation testing [92]. Nanoindentation can be used to evaluate the fracture toughness of material and interfaces in a similar manner to that conventionally used in large scale testing. During loading, tensile stresses are induced in the specimen material as the radius of the plastic zone increases. Upon unloading, the additional stresses arise as the elastically strained material outside the plastic zone attempts to resume its original shape but is prevented from doing so by the permanent deformation associated with the plastic zone. There exists a large body of literature on the subject of indentation cracking with Vickers and other sharp indenters, such as Berkovich indenter. Recently, several methods for fracture toughness estimation without visualization of the crack length have been proposed, for example the works of Field et al. [93], where they relate the crack length produced to the pop-in behaviour during loading or the work of Dahami et al. [94] where they related the crack length with the total load penetration depth for fused silica. Generally, there are three types of crack. A scheme of it can be observed in figure 23: a) Radial crack are vertical half penny type cracks that occur on the surface of the specimen outside the plastic zone and at the corners of the residual impression at the indentation site. These radial cracks are formed by a hoop stress and extend downward into the speciment, but are usually quite shallow. See Figure 23a. b) Lateral cracks are horitzontal cracks that occurs beneath the surface and are symetric with the load axis. They are produced by a tensile stress and often extend to the surface, resulting in a surface ring that may lead to chipping of the surface of the specimen. See Figure 23b. c) Median cracks are vertical circular penny cracks that form beneath the surface along the axis of symmetry and have a direction aligned with the corners of the residual impression. Depending on the loading conditions, median cracks may extend upward and join with surface radial cracks, thus forming two half-penny cracks that intersect the surface. See Figure 23c.

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a)

b)

c) Figure 23. Scheme of the different types of cracks, a) radial crack, b) lateral cracks and c) median cracks.

In adition, other fracture events, such as delamination or radial cracking at the interface, may be activated in case that the indentation field comprises both materials. As in the case of radial cracking, the delamination event will only be detected in an instrumented indentation test for certain materials. Fracture mechanics treatment of these types of cracks seek to provide a measure of fracture toughness based on the length of the radial surface cracks. Attention is usually given to the length of the radial cracks as measured from the corner of the indentation and then radially outward along the specimen surface as shown in figure 24.

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l c a

Figure 24. Crack parameters for Berkovich indenter. Crack length c is measured from the center of contact to end of crack at the specimen surface.

Conventional indentation toughness methods were initially developed for monolithic bulk materials tested by microindentation when well-developed radial cracks form (for example Marshall and Lawn [95], Anstis et al. [96]). The toughness KIC is related to the applied load P, H, E and the cracks dimension, c. The fracture toughness can be obtained with the next equation: 1

K IC

⎛E⎞ 2 P = χ ⋅⎜ ⎟ ⋅ 3 ⎝H⎠ c 2

(27)

where E and H are the Young’s modulus and the Hardness of the material; these values are obtained with nanoindentation technique using the Oliver and Pahrr approach. For Berkovich and Vickers indenters, χ = 0.016. The values obtained by this method will depend on the residual stress in the coating since equation 27 is strictly only valid in the absence of internal stresses. At higher loads using a Berkovich indenter, two crack systems are observed [97]. Initially, radial cracks are observed along the edge of the indenter. These are followed by picture-frame cracks at the edge of the impression once sufficient bending has occurred. For Vickers and Berkovich indenters, cracking threshold in most ceramic materials are about 250 mN or more [98], and since the indentation produced at these loads are relatively large, the cracking thresholds place severe restrictions on the spatial resolution which can potentially be achieved. It should be noted that the cracking threshold depends on the condition of the indenter tip, generally being higher for tips that have been blunted by wear. At a given load, the cube-corner and Berkovich diamonds should, to a first approximation, penetrate the material to produce approximately equal projected contact areas, such as the hardness measured with the two indenters should be about the same.

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Given that the nucleation and propagation of indentation cracks are promoted by large stress and strains, one would then qualitatively expect a reduction in the threshold for the sharper indenter.

1.2.5. The Effective Indenter Shape The effective indenter shape is outlined in figure 25. The basic principles are derived from observation gleaned from finite element simulations of indentation of elastic-plastic materials by rigid conical indenter with a half included angle of 70.3º. During the initial loading of the indenter (figure 25a), both elastic and plastic deformation processes occur, and the indenter conforms perfectly to the shape of the hardness impression. However, during unloading, figure 25b, elastic recovery causes the hardness impression to change its shape. A key observation is that the unload shape is not perfectly conical, but exhibits a subtle convex curvature that has been exaggerated in figure 25b, the contact area increases gradually and continuously until full load is again achieved, a process which must be the reverse of what happens during unloading because both processes are elastic. It is this continuous change in contact area that produces the nonlinear unloading curves. Furthermore, the relevant elastic contact problem is not that of conical indenter on a flat surface, but a conical indenter pressed into a surface that has been distorted by the formation of the hardness impression.

1.2.6. Errors Due to Pile-up and Sinking-in In an indentation into an elastic material, the surface of the specimen is typically drawn inwards and downwards underneath the indenter, and sinking-in occurs. When the contact involves plastic deformation, the material may either sink in, or pile up around the indenter. a)

b)

Load

Unload

P

c)

Reload P

Z = U(r)

elastic Elastic/ Plastic

elastic Z

P P

PP Z = U(r)

r

Effectiv indenter shape Effective indenter

shape

r r

Figure 25. Concepts used to understand and define the effective indenter shape, a) loaded, b) unloaded and c) reloaded. (From W. C. Oliver et al., J. Mater. Res., Vol. 19, No. 1, 2004, pages 3-20).

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In the fully plastic regime, the behaviour is seen to be dependent on the ratio E/σys, where σys is the yield strength of the material, and the strain-hardening properties of the material. The mechanical nature of a typical specimen can be described by a conventional stressstrain relationship that includes a strain-hardening exponent:

σ = E ⋅ ε when ε ≤

σ ys

σ = k ⋅ ε x when ε ≥

(28)

E

σ ys

(29)

E

where k is equal to:

⎡ E ⎤ k = σ ys ⋅ ⎢ ⎥ ⎢⎣σ ys ⎥⎦

x

(30)

Such a complication arises from the pile-up or sink-in of the material around the indenter, which is primarily affected by the plastic properties of the material [99]. In a low-strainhardening alloy, plastically displaced material tends to flow up to (and pile-up against) the faces of the indenter due to the incompressibility of plastic deformation. The result is a barrelshaped impression due to pile-up around the sharp-indenter. In high strain hardening materials, the plasticity deformed region is pushed out from the indenter with the imprint sinking below the initial surface level. The result is a pin-cushion like impression around the sharp indenter, as shown in figure 26. Apparent contact diameter

Apparent contact diameter

Pile-up

True contact diameter

True contact diameter

Sink-in

True contact perimeter Indenter Side Surface

a)

b)

Figure 26. Schematic illustration of a) pile-up and b) sink-in around a sharp indenter.

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As a consequence of pile-up or sink-in, large differences may arise between the true contact area and the apparent contact area which is usually observed after indentation. One significant problem with the determination of the contact area with equations from 12 to 15, is that this method does not account for pile-up of material around the contact impression. When pile-up occurs, the contact area is greater than that predicted by these equations, and both the hardness estimated from equation 5 and the effective modulus from equation 7 are overestimated, sometimes by as much as 50% [100]. Pile-up is large only when the relation between hf/hmax is close to 1 and the degree of work hardening is small. It should also be noted that when this relation is lower than 0.7, very little pile-up can be observed no matter what the work-hardening behaviour of the material. In this case, the contact area given by this method match very well with the true contact area obtained. Furthermore, when this relation is higher than 0.7, the accuracy of the method depends on the amount of work-hardening in the material. For Berkovich indenters, indentations with a large amount of pile-up can be identified by the distinct bowing out at the edges of the contact impression.If pile-up is large, accurate measurements of H and E cannot be obtained using the contact area deduced from the loaddisplacement curve. In this case, the most useful method to correct these problems is known as Cheng and Cheng method [101 and 102]. These equations are function of the total work of indentation (Wtot), and the work recovered during unloading (Wu). The method that they proposed to account for pile-up is based on the work of indentation, which measure from the areas under indentation loading and unloading curves. The relation proposed by Cheng and Cheng can be observed below:

Wtot − Wu H ≅ 1− 5⋅ Wtot E eff

(31)

Combining equation 5 and 7 and considering β as 1. Another equation involving H and Eeff can be obtained:

H 4 Pmax ⋅ 2 = 2 π S E eff

(32)

where Wtot, Wu, Pmax, and S are all measurable from load-displacement curves. Spherical indentation differs from conical or pyramidal indentation in that there is no elastic singularity at the tip of the indenter to produce large stresses. Stress-strain curves can be approximated from indentation data using the classical approach of Tabor . Field and Swain [103 and 104] have applied Tabor’s approach to instrumented indentation and have developed a method that uses the indentation load-displacement data to approximate the stress-strain curve and the work-hardening exponent The approach requires the deformation to be fully plastic. Pile-up geometry can change considerably during the course of spherical indentation and it is therefore not possible to predict the pile-up based on the mechanical properties of the material alone, even when fully plastic deformation is achieved.

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1.2.7. Indentation Size Effect, ISE The indentation size effect (ISE) is extensively studied in literature and the mechanisms driving it remain irresolute. Examples of suggested mechanisms include inadequate measurement capabilities of extremely small indents, presence of oxides or chemical contamination on the surface, indenter-specimen friction, and increased dominance of edge effects with shallow indents [105]. In conclusion, ISE produces a greater hardness at a smaller applied load [106]. Depth-sensing indentation with loads higher than 1 N was proposed to avoid the inaccuracy due to the ISE effect [107]. The Indentation Size Effect (ISE) can be classified in two different types: -

-

Intrinsic: From interference of the size (or thickness) of the sample, the plastic region of the sample, L, or the autonomous plastic regions of the structure (the grains), D, with some of the characteristic microstructural lengths. Extrinsic: Associated to deformation gradients. Size effects in the resistance to plastic flow appear in metals when the dimensions of the specimen or of the zone subjected to plastic deformation are in the range of μm. The heterogeneity of the field of plastic deformation implies differences in the dislocation fluxes across a crystalline volume element (internal storage of “geometrically necessary dislocations, GND”). The analysis of this phenomenon is important because: i) the crystal deformation and fracture processes take place at this length scale and ii) implications on the development of micro-electro-mechanical systems (MEMS) and in the micro-electronics industry.

Recently, it is has become possible to perform indentation test at dimensions of tens to hundreds of nanometers using nano- and microindentation methods. At these small indentations depths classic plasticity theory predicts constants hardness using a geometrically self-similar indenter on a homogenous material. Nevertheless a strong size dependent indentation hardness result is well known. ISE is characterized by an increasing hardness as the indentation depth is reduced to the order of microns or submicrons. This phenomenon was interpreted by Nix and Gao (1998) [108], De Guzman et al. (1993) [109], Poole et al. (1996) [110], McElhaney et al. (1998) [111] and Fleck et al. (1994) [112] who proposed a strain gradient theory to explain the presence of the ISE. Nix and Gao showed that the ISE for crystalline materials can be explained using the concept of geometrically necessary dislocations, which leads to a strain gradient plasticity law. Swadener et al. [113] extended this model for the case of spherical tip indenters. For loads up to and exceeding that to initiate pop-in at the beginning of the elastic/plastic part of the loading curve all indents can be described by an approximately spherical contact. They were analyzed using the approach of Swader et al. , which proposes that spherical indenters show a dependence of hardness on the indenter radius rather than on the depth of the penetration. The load dependence of the hardness is referred to as the ISE. The ISE can be seen remarkably under extremely low loads such as nanoindentation test [114]. Up to low various factors such as strain-hardening and friction effects have been reported to explain the phenomenon of the ISE [115].

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This theory assumes that the stress flow of metals depends on the density of statistically stored and geometrically necessary dislocation. The density of the statistically stored dislocation varies with the effective strain whereas the density of the geometrically necessary dislocations depends on the strain gradient. In crystalline solids (such as ceramics), dislocations are responsible for sustained plastic deformation and dislocations impedes the motion of new dislocations. Dislocations are generated during plastic deformation (when the applied load is higher than yield strength of the studied material). The dislocations are then moved, and stored. Storage of dislocations abets strain hardening. It is postulated that dislocations become stored because they either accumulate by randomly trapping each other or they are required for compatible deformation of various parts of the material. When they randomly trap each other, they are often known as the statistically stored dislocation , whereas when they are required for compatibility purposes, they are often called geometrically necessary dislocation and they are related to the gradient of plastic shear strain in a material ISE have been observed since the early days of indentation testing [116]. In many cases, this effect is largely due to the rising uncertainties involved in making and measuring small indentations. There has been voluminous literature devoted to study the origin of the ISE. Consequently, several empirical or semi-empirical equations widely applied in ceramic materials can be applied to solve this problem, such as Meyer’s law [117], the Hays-Kendall approach [118], the elastic recovery model , the energy-balance approach [119], the proportional specimen resistance model [120], and others, that have been proposed for describing the variation of the indentation size with the applied test load. a) Meyer’s Law The most widely used empirical equation for describing the ISE is the Meyer’s law, which correlates the test load and the resultant indentation size using a simple power law ,

Pmax = A ⋅ hcn

(33)

where A and n are constants that can be derived directly from the curve fitting of the experimental data. n is known as Meyer’s index and it is usually considered as a measure of the ISE. To obtain the A and n values, the values obtained from the nanoindentation data are plotted in an ln Pmax – ln hc scale. Each set of the data shows an excellent linear relationship, implying that the traditional Meyer’s law is suitable for describing the nanoindentation data. Trough linear regresion analyses, the best fit values of the A and n were obtained. b) Hays-Kendall Approach When examining the ISE in the Knoop hardness testing of a number of metals, Hays and Kendall advanced that there exists a minimum level of the applied test load, W, named the test-speciment resistance, below which permanent deformation due to indentation does not initiate, but only elastic deformation occurs.

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An effective indentation load has been introduced, Peff = Pmax – W, and proposed the following relationship,

Peff = A1 ⋅ hc2 = Pmax − W

(34)

where W and A1 are constants independent of the test load for a given material. Equation 34 predicts that a plot of Pmax versus hc2 would yield a straight line. This equation provides a satisfactory description of the nanoindentation data for the studied material. c) Elastic Recovery Model or Elastic/Plastic Deformation Model In microhardness test, the indentation size is measured after the indenter is removed from the specimen surface. Note that the elastic recovery would occur in the vicinity of the remaining indentation impression after the indenter is removed so that the indentation size would shorten to a certain degree [121]. Considering this effect, Tarkanian et al. suggested that the measured indentation size should be corrected with a revised term in order to obtain the true hardness. The true hardness can be calculated as follows:

Ho = k ⋅

P (d + d o )2

(35)

where do is the correction in the indentation size d due to the elastic recovery and k is a constant dependent on the indenter geometry. Furthermore, several authors have pointed out that, for calculating the hardness from the recorded indentation test such as the nanoindentation, similar correction in indentation size should be considered because of the elastic recovery associated with the new bands of plastic deformation [122] and /or the blunting of the indenter tip [123]. Thus it is necessary to check if equation 35 is suitable for describing the nanoindentation data obtained for ceramics. To analyze the nanoindentation data, equation 35 may be re-written in the next form: 1

Pmax2 = χ

1

2

⋅ hc + χ

1

2

⋅ ho

(36)

where ho is the correction in hc and χ = Ho/k is a constant related to the true hardness. Equation 36 allows to determine ho and χ from the plots of Pmax1/2 against hc. There are two ways to calculate the true hardness based on the elastic recovery model: Directly use equation 35, or Obtain χ. d) Proportional Specimen Resistance Model or PSR Model Proportional specimen resistance (PSR) model was proposed by Li and Bradt . This model can be considered as a modified Hays-Kendal approach. In this model, the test-

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193

specimen resistance to permanent deformation is assumed not to be a constant, but this factor increases linearly with the indentation size. This phenomena is governed by the next equation:

W = a1 ⋅ hc

(37)

To a first approximation, the equation 37 can be considered to be similar to the elastic resistance of a spring with the opposite sign to the applied test load. The effective indentation load and the indentation dimension can be related as follows:

Peff = Pmax − W = Pmax − a1 ⋅ hc = a 2 ⋅ hc2

(38)

where a1 and a2 are constants for a given material. According to the analysis of Li and Bradt , the parameters a1 and a2 can be related to the elastic and the plastic properties of the test material. a2 was suggested to be a measure of the socalled true hardness, Ho. For the nanoindentation test with a Berkovich tip indenter, Ho can be calculated directly from a2 with:

Ho =

Peff 24.5 ⋅ h

2 c

=

Pmax − a1 ⋅ hc a = 2 2 24.5 ⋅ hc 24.5

(39)

Equation 38 can be rearranged as;

Pmax = a1 + a 2 ⋅ hc hc

(40)

which enables to determine both a1 and a2 from the plot of Pmax/hc against hc. An alternative explanation for the physical meaning of equation 40 was proposed by Flöhlich et al. [124] based on energy-balanced analysis. According to the energy balance consideration, the parameters a1 and a2 in equation 40 are related to the energies dissipated for creating a new surface of a unit area and for producing the permanent deformation of a unit volume, respectively. Also, a2 is a measure of the true hardness. e) The Modified PSR Model Examining the load-dependence of the microhardness of some ceramics measured in a wide load range, Gong et al. [125], found that the resultant P/d-d (where d is a half-length of the Vickers indentation) curves exhibit significant nonlinearity and argued that the PSR model mentioned above may only be used to represent the experimental data measured in a narrower range of applied loads. Gong et al. suggested that the PSR model should be modified as follows:

Pmax = a o + a1 ⋅ hc + a 2 ⋅ hc2

(41)

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where a0 is a constant related to the surface residual stress associated with the surface machining and polishing and a1 and a2 are the same parameters as those in equation 40. An equation with the same form as 41 has been also deduced based on a modified energy-balanced analysis The modified PSR model provides two different ways to obtain the true-hardness:

Ho 1 =

Pmax − ao − a1 ⋅ hc 24.5 ⋅ hc2

(42)

a2 24.5

(43)

Ho

2

=

2. Mechanical Properties 2.1. State of the Art of Mechanical Properties of YBCO Samples The mechanical properties of YBCO samples have been studied during the last years. The most important properties studied have been the hardness (at micrometric and nanometric scale), the Young’s modulus and the toughness fracture. Some authors studied the mechanical properties at room or at cryogenic temperatures also known as work temperature. The techniques used to perform these studies have been: microindentation or nanoindentation, to obtain the hardness of the YBCO samples; the bending, X-ray diffraction and others, to obtain the Young’s modulus; and microhardness to obtain the toughness fracture. Reported values of Young’s modulus, hardness and fracture toughness of YBCO composite obtained using different experimental techniques are summarized in tables 3, 4 and 5, respectively. Table 3. Literature values of Young’s modulus for YBCO with different techniques Author Joo. et al. Lucas et al. [126] Soifer et al. [127] Güçlü et al. [128] Ledbetter et al. [129] Sheahen et al.[130] Goyal et al. [131] Soifer et al. [132] Goyal et al. [133] Goyal et al. Reddy et al. [134]

Material YBCO YBCO + 5% vol. Ag YBCO + 10% vol. Ag YBCO + 15% vol. Ag Y-123 YBCO film YBCO polycrystalline, 50K YBCO polycrystalline, 160K YBCO polycrystalline, 180K YBCO polycrystalline, 293K YBCO polycrystalline Syntherized with Ag Syngle crystal Thin film Y-211 Texturized Texturized

Young’s modulus (GPa) 110 103 103 97 154.30 ± 16.34 210 47.20 29.63 28.47 9.39 90.8-101.8 75-120 220 ± 20 210 213 182 95.89

Method Pulse echo technique Indentation Nanoindentation Vickers indentation Ultrasonic Ultrasonic Nanoindentation with AFM Nanoindentation Nanoindentation Ultrasonic

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195

Table 4. Literature values of Hardness for YBCO with different techniques Author Lucas et al. Li et al. [135] Soifer et al. Yoshino et al. [136]

Güçlü et al. [137] Lim et al. [138] Cook et al. [ Goyal et al. Goyal et al. Soifer et al. Goyal et al.

Material Y-123 YBCO, MTG-1100ºC, 5 min YBCO, MTG-1100ºC, 10 min YBCO, MTG-1100ºC, 15 min YBCO, Solid state reaction YBCO film YBCO 40K YBCO 293 K YBCO polycrystalline, 50K YBCO polycrystalline, 160K YBCO polycrystalline, 180K YBCO polycrystalline, 293K YBCO single crystal YBCO Textures Textures Thin film Y-211

Hardness (GPa) 10.28 ± 1.67 5.4 5.0 5.1 4.7 8.5 18 ± 2.5 5.2 ± 0.5 3.58 1.03 0.95 0.53 7.81 ± 0.23 8.7 6.7 10.8 8.5 14.0

Method Indentation Vickers indentation Nanoindentation Vickers indentation

Vickers indentation Nanoindentation Vickers indenter Vickers indenter Nanoindentation Indentation with AFM Nanoindentation

Table 5. Literature values of fracture toughness for YBCO with different techniques Fracture Toughness (MPa·m1/2) YBCO 1.60 YBCO + 5% vol. Ag 2.10 Joo. et al. [139] YBCO + 10% vol. Ag 2.50 YBCO + 15% vol. Ag 2.80 1.53 Lenblond-Harnois et al. YBCO [140] YBCO + 5% wt Ag 1.88 YBCO, MTG-1100ºC, 5 min 1.9 YBCO, MTG-1100ºC, 10 min 1.7 Li et al. YBCO, MTG-1100ºC, 15 min 1.7 YBCO, Solid state reaction 1.3 YBCO 40K 0.4 Yoshino et al. YBCO 293 K 1.3 Joo et al. [141] YBCO 1.6 YBCO + 5 vol% Ag 2.2 YBCO + 10 vol% Ag 2.6 Leenders et al [142] YBCO + 30 mol% Y-211 1.01 YBCO + 60 mol% Y-211 1.44 Cook et al. [143] YBCO 1.1 Sheahen et al. YBCO 0.8-1.0 Sheahen et al. YBCO textured 1.6 Fujitomo et al. [144] YBCO textured 0.99-1.20 Sheahen et al. YBCO with 5, 15 and 25 % of Ag 1.6 textured Sheahen et al. YBCO with 20% Ag 3.8 Fujitomo et al. YBCO with Ag 1.6-2.1 Author

Material

Method Single-edge-notch beam Vickers indentation method Vickers indentation

Vickers indentation Single-edge-notch beam Vickers indenter Vickers indenter Bending Bending Vickers indenter Bending Bending Vickers indenter

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From last tables it can be observed the high scattered differences between the different studies carried out last century. Now, we have performed a study of the mechanical properties of YBCO samples textured by Bridgman and TSMG techniques by Nanoindentation. Figure 27a to 27c, a representation of Young’s modulus, Hardness and fracture toughness respect the technique can be observed. 250 220

210 213 200

210

182 154,3

E (GPa)

150

96,3

100

110

95,98

103 103

97

47,2

50

29,6 28,4 9,3

0

Indentación Ultrasonic

Nanoindentación AFM

Vickers Pulse echo technique

a) 20 18

18 16

14

H (GPa)

14 12

10,8

10,28

10

8,7

8 6

8,5 6,7

5,4 5,1 4,7

4

8,5

7,81

5,2 3,58 1,03 0,95 0,53

2 0

Indentación

Vickers

Nanoindentation

b) Figure 27. Continued on next page.

AFM

Mechanical Characterization at Nanometric Scale…

197

4

3,8

3,5

1/2

KIC (MPa·m )

3

2,8

2,6

2,5

2,5

2,2

2,1

1,88 1,9

2 1,6

1,6

1,5

1,53

1,85

1,7 1,7 1,3

1,3

1,01

1

1,6 1,6

1,44 1,1 1,05

0,9

0,4

0,5 0

Single edge notch beam

Vickers

Bending

c) Figure 27. Representation of the bibliographic mechanical property versus the different technique for a) Young’s modulus, b) Hardness and c) Fracture toughness.

2.2. Mechanical Properties of YBCO Samples Textured by Bridgman and TSMG Technique by Nanoindentation 2.2.1.Plastic Deformation Experimental Conditions The nanoindentation technique was performed by a Nano Indenter® XP Systems (Systems Corporation) equipped with Test Works 4 Professional level software. Nanoindentation imprints were observed with optical microscope, with an AFM NanosCope III-A atomic force microscope, and with a field emission Hitachi H-4100 scanning electron microscope. The experiments were performed on the (001) plane at room temperature for the different indenters used (Berkovich and Spherical tip indenters). First of all, a plastic study was carried out with a Berkovich indenter. In this case, the hardness, Young’s modulus and fracture toughness were studied at different applied loads: 5, 10, 30 and 100 mN. The loading/unloading time was selected to be constant for all indentations: 15 s. Table 6 shows fixed test parameters to perform measurements of nanoindentation when a Berkovich indenter was used. The small nanoindentations were made by a three-sided pyramid Berkovich diamond indenter, see figure 19.a. The displacement or also known as penetration depth, was continuously monitored and load-time history of indentation recorded. Each hardness, Young’s modulus, fracture toughness and other values listed in this section, are an average of 40 measurements performed on two different samples in order to achieve statistical

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significance. The values obtained by the Oliver and Pharr equation (equations 5, 7 and 8) have been corroborated by FE-SEM. After that, we studied the Young’s modulus evolution (from 0 to 700 mN of applied load) with continuous stiffness measurement, CSM, and with a spherical tip indenter. CSM supplies information about the first steps of the indentation and allows to know the evolution of the Young’s modulus with the penetration depth. The radii of this tips was 25 μm. . Experimental Curves The load/unload curves can give a qualitative information of the hardness of each phase of the study when the applied load is lower than 10 mN. For these loads, the mechanical properties of each phase can be isolated. Table 6. Test inputs of nanoindentation Name Allowable drift rate Load rate multiple for unload rate Maximum load Number of times to load Peak hold time Per cent to unload Time to load

Value 0.05 1 5, 10, 30 and 100 5 30 90 15

Units nm·s-1 mN s % s

Figure 28, shows the load/unload curves for YBCO samples textured by Bridgman technique when the applied load was 5 mN. 5,0 4,5 4,0 3,5

P (mN)

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

Y-123 Y-211 Y-123/Y-211

1,0 0,5 0,0 0

50

100

h (nm)

150

a) Figure 28. Continued on next page.

200

Mechanical Characterization at Nanometric Scale…

199

5,0 4,5 4,0

P (mN)

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

Y-123 Y-211 Y-123/Y-211

1,0 0,5 0,0 0

50

100

h (nm )

150

200

b) Figure 28. Load/Unload curves for YBCO samples when the applied load was 5 mN textured by a) Bridgman technique and b) TSMG technique.

Figure 28a, shows the qualitative manner to predict the hardness of each phase. The hardness of one material is a function of two different parameters, the maximum applied load and the contact area. The last parameter is a function of contact penetration and this one of the maximum penetration. For this reason, when one phase presents a high penetration depth, this one has a low hardness value; one material has a high hardness value when the penetration depth is low. From figure 28a, the distribution of hardness for each phase is: HY-123 < HY-123/Y211 < HY-211. From figure 28b, the same effect that in the figure 28b can be observed. In this case, the difference between the different phases is less than in the figure 28a. The YBCO samples textured by TSMG technique present the same relation: HY-123 < HY-123/Y-211 < HY-211. This technique presents a little difference in the hardness value at ultra low load such as 5 mN, because after the oxygenation process, the orthorhombic phase presents a high porosity in the ab-plane. Figure 29, shows the load-unload curve for YBCO samples textured by Bridgman and TSMG technique recorded by CSM when the applied load was 700 mN. Figure 29, shows that for the two different techniques of study the first steps of nanoindentation present the same tendency. The scattered occurs mainly in the unload curve, for this reason the Young’s moduli will be different. Characterization Imprints All imprints realise in the samples of study have been observed with Optical Microscopy and FE-SEM. In order to know the correct value of the hardness, Young’s modulus and fracture toughness of Y-123, Y-211 and Y-123/Y-211 in the case that the applied load permit isolate each mechanical property.

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J.J. Roa, X.G. Capdevila and M. Segarra 700 600

P (mN)

500 400 300 200 100

Bridgman technique TSMG technique

0 0

100

200

300

400

500

600

700

800

900

1000

1100

h (nm) Figure 29. P-h curve for YBCO samples textured by Bridgman and TSMG technique recorded by CSM when the applied load was 700 mN.

Figure 30 shows indentation imprints performed by applying 30 and 100 mN on the ab plane of the monodomain, the Y-123 matrix and Y-211 inclusions.

50 μm

a) Figure 30. Continued on next page.

Mechanical Characterization at Nanometric Scale…

30 mN

100 mN

201

50 μm

b) Figure 30. Optical microscope micrographs of nanohardness impressions developed on the surface of a sample of YBCO orthorhombic phase (ab plane) at room temperature, a) for Bridgman textured sample (From Roa et al. Nanotechnology, 18, 2007, page 38571-1 to 38571/1-38571/6) and b) for TSMG textured sample, when the applied load was 30 and 100.

In this figure, the particles or inclusions can be observed homogeneously distributed in the textured samples, so that they can be easily identified but not indented separately. It is important to highlight that the size of Y-211 inclusions (from 1 to 5 μm, approximately) is smaller than the nanoindentation imprints performed at these loads, so that we can only measure the mechanical properties of the composite (Y-123+Y-211) when the applied load was higher than 10 mN. In Figure 30b, it can be observed that the residual indentations are highly affected by the superficial porosity. Every residual imprint shows a crack in its corner. Thus indicating, give us that YBCO is a brittle material. When the applied loads are higher than 10 mN (such as 30 and 100 mN), only the mechanical properties of composite can be determined and we cannot isolate the contribution of each phase. Figure 31 shows Y-123, Y-211 and YBCO composite imprints for samples textured by Bridgman technique. Figure 31b, shows propagation of the cracks following the corners of the indentation, for Y-211 precipitate. Nanohardness of Y-211 is twice as high as the Y-123 matrix (see table 7); for this reason the Y-211 phase presents a brittle fracture. The analysis of SEM images can be used to determine the profile impressions with nanometer resolution and to provide information about the shape change on unloading. As can be seen in this figure, the indentation exhibit triangular geometry.

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1.2 μm a)

1.2 μm b)

1.2 μm c) Figure 31. Micrograph of nanoindentation imprints obtained by FE-SEM when the applied load was 10 mN. a) matrix Y-123; b) precipitate, Y-211 and c) Y-23/Y-211 composite (From Roa et al. Nanotechnology, 18, 2007, page 38571/1-38571/6).

Figure 32 shows Y-123, Y-211 and YBCO composite imprints at 10 mN for samples textured by TSMG technique.

Mechanical Characterization at Nanometric Scale…

203

1.50 μm a)

1.50 μm

b)

1.50 μm

c) Figure 32. Micrograph of nanoindentation imprints obtained by FE-SEM when the applied load was 10 mN. a) matrix Y-123; b) precipitated, Y-211 and c) Y-23/Y-211 composite.

In figure 32b, a crack at the corners of the imprint and another fracture mechanism known as chipping, can be observed. When the applied load is higher than 10 mN, the mechanical properties of each phase cannot be isolated because the size of the residual imprint is lower than the size of Y-211. Figure 33, shows the nanoindentation imprints obtained by FE-SEM when the applied load was 30 and 100 mN for samples textured by Bridgman technique.

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2.0 μ m a)

μm m 3.0 μ 3.0 b) Figure 33. Micrograph of nanoindentation imprints obtained by FE-SEM of samples textured by Bridgman technique at applied load of, a) 30 mN, and b) 100 mN.

30 μm Figure 34. Micrograph of nanoindentation imprints obtained by FE-SEM of samples textured by TSMG technique.

Figure 33a, shows cracks at the corners of the imprints; in this case the fracture toughness can be calculated. Inside de imprints, a radial cracks can also be observed. The fracture mechanism in this case will be further discussed.

Mechanical Characterization at Nanometric Scale…

205

Figure 33b, also shows cracks, radial ones inside the imprint and length ones at two of the three corners of the imprint. The missing crack at one corner has been stopped by a Y-211 particle. In this case, the fracture toughness cannot be calculated. In the case of the samples textured by TSMG technique, when the applied load was 100 mN, we can observe a lot of porosity and sink-in besides the residual imprints (Figure 34). Figure 35, shows residual nanoindentation imprints at applied loads of 30 and 100 mN. When the load applied was 30 mN (figure 35a), cracks at the corners of the imprints cannot be observed. No cracks can be observed at the corners of imprints, but only radial cracks.

2.00 μm

a)

5.00 μm

b) Figure 35. Nanoindentation imprints obtained by FE-SEM of samples textured by TSMG technique, a) 30 mN of applied load and b) 100 mN of applied load.

Note that when the applied load was 100mN, all imprints show a high porosity beside the imprint for both texturing technique, see figure 34 and 35b.

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Hardness Table 7 shows the calculated nanohardness values of the orthorhombic phase of Y-123, Y-211 and Y-123/Y-211 composite for the different monodomains samples textured by Bridgman technique, when the applied loads were 5, 10, 30 and 100 mN. Table 7. Nanohardness of orthorhombic phases of YBCO textured by Bridgman (From Roa et al. Nanotechnology, 18, 2007, page 38571-1 to 38571/1-38571/-6) and TSMG technique for applied loads of 5, 10, 30 and 100 mN Applied Load (mN)

Phase

5

Y 123 Y 211 Y123/Y211composite Y 123 Y 211 Y123/Y211composite

Hardness; H (GPa) Bridgman TSMG 11.0 ± 0.5 11.41 ± 0.51 20.0 ± 1.0 15.0 ± 1.05 15.2 ± 0.3 14.86 ± 1.05 9.8 ± 0.4 10.59 ± 0.45 18.1 ± 0.5 15.35 ± 0.58 14.4 ± 0.7 14.92 ± 0.98

Y 123 Y 211 Y123/Y211composite

11.4 ± 0.4 17.1 ± 0.5 15.3 ± 0.3

11.12 ± 0.96 17.14 ± 0.25 14.58 ± 0.78

Y 123 Y 211 Y123/Y211composite

11.0 ± 0.3 16.7 ± 0.6 14.9 ± 0.2

9.59 ± 0.76 16.89 ± 0.86 14.86 ± 0.95

11.2 ± 0.3

8.71 ± 0.95

Sample 2

11.0 ± 0.3

8.26 ± 0.79

Sample 1

8.8 ± 0.2

7.96 ± 0.72

9.1 ± 0.1

7.88 ± 0.89

Monodomain

Sample 1

Sample 2

Sample 1 10 Sample 2

Sample 1 30

100 Sample 2

Y123/Y211 composite

Y123/Y211 composite

Nanohardness and Young’s modulus of each phase (Y-123 matrix and Y-211 inclusions) can be determined when the applied load was ultra-low (less 10 mN), each phase can be indented and the respective mechanical properties can be isolated. For higher loads, the mechanical properties of the interaction of both phases can be obtained the YBCO composite or Y-123/Y-211.

Mechanical Characterization at Nanometric Scale…

207

Values of Y-211 phase were higher than for Y-123 in both cases of study, for the samples. Nanohardness of Y-211 was about twice as high as for Y-123. This fact could be due to different reasons: I). II). III).

Ionic bond of Y-211 is stronger than the Y-123 bond, High anisotropy of dislocation confined onto a (001) plane [145], and/or the melt processed ceramic composites contain a dense population of fine peritectic inclussion embebbed in a matrix, which drastically affects the microstructure acting as nucleation sites for dislocation

When the applied load is 100 mN, overall nanohardness is very similar for both samples (for Bridgman technique 8.90 GPa, and TSMG technique 7.92 GPa) but is lower than the nanohardness of the Y-123 phase. Therefore, these results are in agreement with a previous work reported by Lim and Chaudhri which studied the hardness of YBCO single crystal by Nanoindentation technique and the hardness value obtained was 7.81 GPa. This value is in agreement with the result obtained for samples textured by TSMG technique. On the other hand, the hardness value obtained for samples textured by Bridgman technique is in agreement with a previous work reported by Soifer et al. and Cook et al.. When the applied load is higher than 10 mN we are working within the microindentation range and we can observe microcraks at the corners of imprints (see figure 33a), thus causing the reduction of the hardness value (H100mN < H30mN). This effect is known as indentation size effect, ISE. Verdyan et al [146] reported a nanohardness for orthorhombic YBCO thin film can be around 8.5 GPa, when the applied load varies between 0.1 and 9mN, which are similar values to those found in the present study for the orthorombic phases. Values of the Young’s modulus are also comparable to our results for the orthorhombic composite, studied at loads of 30 and 100 mN. If we compare our data for Hardness with data reported of YBCO (see table 7), we conclude that the broad distribution of hardness observed in table 4 can be attributed both to the different measuring techniques and to different quality of the studied YBCO samples (grain structure, porosity, texturing process, etc). Young’s Modulus Table 8 shows the calculated nanohardness and Young’s modulus values of the orthorhombic of Y-123, Y-211 and Y-123/Y-211 composite for the different monodomains samples studied when the applied loads were 5, 10, 30 and 100 mN textured by TSMG technique. When the applied load is 100 mN, overall Young’s modulus are very similar for both samples and techniques (for Bridgman technique 173 GPa and TSMG technique 129 GPa) but is lower than the Young’s modulus of the Y-123 phase. Therefore, the results obtained by Bridgman technique are in agreement with a previous work reported by Goyal et al. which studied the Young’s modulus of texturized YBCO by Nanoindentation technique and the Young’s modulus value obtained was 182 GPa. The Young’s modulus value obtained for samples textured by TSMG technique is in agreement with a previous work reported by Joo et al. which value had been obtained by pulse echo technique. Goyal et al. studied the Young’s

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modulus of Y-211 phase present in YBCO samples by Nanoindentation; this value was 213 GPa. This value obtained by Goya et al.] is in agreement with a value obtained by Nanoindentaion for YBCO samples textured by TSMG technique when the applied load was 10 mN and the Y-211mechanical properties have been obtained. Table 8. Young’s modulus of orthorhombic phases of YBCO textured by Bridgman (From Roa et al. Nanotechnology, 18, 2007, page 38571-1 to 38571/1-38571/-6) and TSMG technique for applied loads of 5, 10, 30 and 100 mN Applied Load (mN)

Phase

5

Y 123 Y 211 Y123/Y211composite Y 123 Y 211 Y123/Y211composite

Young’s modulus; E (GPa) Bridgman TSMG 176.44 ± 15.31 193 ± 7.60 224.35 ± 20.45 199 ± 10.20 207.33 ± 21.65 204 ± 7.12 200 ± 6.53 175.89 ± 15.39 189 ± 8.56 223.54 ± 16.23 198 ± 7.02 208.26 ± 19.52

Y 123 Y 211 Y123/Y211composite

185 ± 3.16 209 ± 4.29 201 ± 6.92

173.62 ± 18.20 207.27 ± 10.58 189.61 ± 17.87

Y 123 Y 211 Y123/Y211composite

192 ± 4.21 203 ± 3.56 206 ± 5.49

174.59 ± 15.35 206.58 ± 9.59 190.12 ± 15.2

179 ± 5.49

140.09 ± 14.21

Sample 2

181 ± 4.35

139.85 ± 13.25

Sample 1

171 ± 2.51

128.88 ± 5.54

175 ± 3.87

130.37 ± 5.69

Monodomain

Sample 1

Sample 2

Sample 1 10 Sample 2

Sample 1 30

100 Sample 2

Y123/Y211 composite

Y123/Y211 composite

Several measurements of Young’s modulus for Y-123 have resulted in values scattered within the range E=40-200 GPa. Most probably, this scatter is caused by residual porosity and bad contacts between the grains [147]. Sakai et al [148] reported a value of 370 GPa for 5 mm3 cubic specimens cut from a single-crystal Y-123 prepared by TSMG using a single domain of Sm-123 as a seed. The authors attribute the difference of Young’s modulus to the 40% excess of Y-211 phase present in the material. Obtained results of Young’s modulus are in agreement with Johansen et al. and Alford et al [149] when the applied loads were 30 and 100 mN.

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If we compare our data for Young’s modulus with data reported for YBCO (see table 8) we also conclude that the broad distribution of Young’s moduli observed in table 3 can be attributed both to the different measuring techniques and to the different quality of the studied YBCO samples (grain structure, texture, and others). The tendency of the Young’s modulus versus penetration depth obtained with a spherical tip indenter and CSM recorded for each one of these texturing techniques is presented in the next figures. 160 140 120

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Figure 36a, shows that the Young’s modulus increases until a constant value, which is around 300 nm of penetration depth. When the penetration depth is lower than 300nm, the spherical contact area could be smaller than the size of Y-211 particles. So, at first stages, we could measure each phase or interface separately. This can explain the highly scattered values that we can see in figure 36a. When the residual imprints are higher than the superficial defects, the Young’s modulus value of the samples textured by Bridgman technique present only a constant value with the penetration depth. Figure 36b, shows two different tendencies of the Young’s modulus: until 300 nm a high dependence of the Young’s modulus with the penetration depth can be seen. On the other hand when the penetration depth is higher than 300 nm, this trend decreases. In this case, a constant value of Young’s modulus cannot be achieved. The difference sample response between Bridgman and TSMG monodomain could be related to the fragilization of the material due to the high porosity density present in the TSMG samples. Even for penetrations depth lower than 300 nm, we cannot isolate the porosity effect over the Young’s modulus value. This is reflected in figure 36b by the high homogeneity on the measurements. For higher applied loads, in the Bridgman samples with lower porosity the Young’s modulus tends over the Y-123/Y-211 value. In TSMG samples, the inner porosity not allows this kind of stabilization. Table 9, shows the Young’s modulus values for the YBCO samples obtained with a spherical tip indenter at 300 nm of penetration depth using CSM. Table 9. Young’s modulus of YBCO samples obtained by spherical indentation at 300 nm of penetration depth using CSM Material YBCO

Technique Bridgman TSMG

Young’s modulus (GPa) 120 ± 20 108 ± 8

The Young’s modulus of YBCO samples textured by Bridgman technique obtained with spherical tip indenter are similar to the values obtained with Berkovich indenter, see table 8. In Figure 36 and Table 9, the values present a difference tendency before and after than 300 nm of penetration depth. When the sample was textured by TSMG technique at 300 nm of penetration depth, the different curves of Young’s modulus versus penetration depth present a lower scattered value. In this case, when the penetration depth is lower than 300 nm the sample present a high homogeneity between Y-123 and Y-211 particles. For penetration depth higher than 300 nm, the Young’s modulus does not present a constant value. In the other hand, for samples textured by Bridgman technique when the penetration depth is lower than 300 nm, the material presented a high interaction with the surface’s defects; for this reason the Young’s modulus present high scattered values. For penetrations depths higher than 300 nm, the material presents a constant value of Young’s modulus. Indentation Size Effect The indentation size effect is generally attributed to train gradient plasticity that generates geometrically necessary dislocations; for ceramic compounds where the plasticity is limited allow loads, the elastic recovery can be significant. When the indentation size is smaller, the density of geometrically necessary dislocations decreases and, as a result, hardness becomes

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higher at small loads [150]. In fact, from the present calculations of the indentation height before and after removal of the load, it is found that 10% of total work done during the indentation at loads of 5, 10, 30 and 100 mN, is due to elastic deformation. While most of the methodologies of analysis of instrumented indentation test assume that materials are homogeneous and continuous this is not the case for real materials, as the continuous broke at a given length (size effect). However, nanohardness can be strongly affected by the presence of defects and impurities that can cause almost no change in dislocation movement, which would affect the hardness. This phenomenon is known as ISE. Instrumented indentation is, a priori, a good technique for local evaluation of the residual stresses, and, for that purpose, several methods have been developed, specially for metals with relatively low yield stress and that do not work harden appreciably when indented with sharp indenters. However, these tests work well when the measured residual stress is in the order of magnitude of the order of the hardness of the material. In YBCO materials, this is often the case because of the high hardness of these materials. In this material, the differences obtained in the load-unload curve from a stressed to an unstressed material are too small to qualitatively extract a value of residual stress. The load-displacement curves obtained for various loads and techniques are shown in figure 37. For each material, all the data points of ten measurements are given in this plot when the applied load was 100 mN. The same effect also can be observed at different studied applied loads, such as: 5, 10 and 30 mN. For YBCO samples textured by Bridgman and TSMG techniques, slight scatters exist among the different load-displacement curves. These scatters may be attributed to the intrinsic microstructural inhomogeneity of these two different techniques to produce single crystals. 100 90 80

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Figure 37. a and b, shows the load-unload curves for YBCO samples textured by Bridgman and TSMG techniques at 100 mN of applied load, respectively. At this load, the hardness and Young’s modulus of YBCO sample has been studied because the residual indentation imprint is higher than the size of Y-211 particles. If the hardness of YBCO sample is the same for all indentations performed on the ab-plane, then we only have to look at one load-unload curves. Nevertheless, in this figure we can observe a high scattered loadunload curves; this effect may be probably due to the ISE. When the applied load is increased, their hardness value is decreasing due to the same factor. Some authors have studied the ISE effect and proposed some equations to solve these problems, such as: Meyer’s lawn, HaysKendall approach, elastic recovery model, proportional specimen resistance model or PSR model and the modified PSR model (see equations 33 to 43). Figure 38 shows the experimentally determined nanoindentation hardness, H, as a function of the peak load, Pmax, for the tested materials. Until 30 mN of applied load, the high tendency is due to the existence of two different phases embedded in the samples of study. When the applied load is 10 mN lower, the mechanical properties of different phases can be studied (Y-123, Y-211, Y-123/Y-211). In figure 38, a high dependency at low applied loads can be observed. This effect is due to the difference between the different phases present in YBCO samples. The ISE can be used for isotropic materials. In fact, YBCO samples present high anisotropy properties in the ab-plane. The different models proposed to study the ISE for ceramics, present some problems to apply in YBCO samples, such as:

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These approaches do not calculate or evaluate the area factor, and the contact area can be strongly affected at low applied loads. This effect produces and overestimated hardness value. The ISE is only affected by the dislocation movement. YBCO samples and a lot of ceramic compounds, present another mechanism such as radial cracks, fracture mechanisms, etc. If the material presents this mechanism, the equations proposed for this study cannot be applied.

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In conclusion, YBCO samples textured by Bridgman and TSMG techniques may present ISE. These phenomena’s cannot be solved because the different equations for ceramics compounds do not a present the different fracture mechanism. If someone has to evaluate ISE for ceramic compounds, the best thing to do is perform a high amount of nanoindentation imprints, and apply the Weibull approach in order to have an average of measurements. Investigations have confirmed that hardness numbers calculated with the Oliver and Pharr method show load dependency as the manner of ISE depicted in figure 38. In order to describe the ISE behaviour, several relationships between applied indentation test load, P, and penetration depth, h, have been presented in the literature [151]. In the same figure it, can be observed that at small loads (less than 10 mN), the Bridgman and TSMG samples have the same Hardness. When the applied load increase, the difference between the two different methods of texture increase. It is due to the high porosity present in TSMG sample, see figure 39.

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2 μm Figure 39. Optical micrograph of YBCO samples textured by TSMG technique.

Table 10. Relation hf/hmax for each phase, load and technique Phase

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D. Pile-up and Sink-in Problems Another common problem in the determination of the hardness by Nanoindetation technique, is the pile-up and sink-in, for more detail see figure 26. The most common effect in YBCO samples is the sink-in effect. The pile-up or sink-in can be obtained from the rate between the final and the maximum penetration depth. If this relation is higher than 1.0, YBCO may present a pile-up effect. On the other hand, if the relation is lower than 1.0, the sample could presented sink-in effect. This study has been performed in both samples textured by these two different methods. Next table summarizes the relation hf/hmax for each phase, load and technique. Every value presented in the table is an average of eighty indentations performed in two different single crystals.

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When the applied load is lower than 10 mN, the relationship can be isolated for each phase. At 5 and 10 mN of applied load, the Y-211 and Y-123/Y-211 phases do not present pile-up effect. On the other hand, Y-123 can present sink-in effect. In this case, the area of the residual imprint can be overestimated and produce a reduction of the hardness of these phases; this effect strongly affected the samples textured by Bridgman technique. Whereas it cannot be observed on samples textured by TSMG technique. In conclusion, the pile-up/sinkin can be related to the hardness of each phase. When the applied load is higher than 10 mN, we cannot isolate this relation for each phase present in the studied material; in this case, only the existence of the sink-in in the Y-123/Y-211 can be determined. When a residual imprint present pile-up or sink-in effect, principally for samples textured by Bridgman technique, two different things can be performed: -

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Observe all the residual imprints with scanning electron microscopy (SEM) and calculate the contact area. Compare the contact area obtained with SEM and the area supplied by the Nanoindenter. If the areas have the same value, the pile-up effect will not exist, but a fracture mechanism takes place. In the case that the two areas have a high scattered value, the Cheng and Cheng approach must be used, and the hardness corrected.

In this study the Cheng and Cheng equations have not be used to correct the contact area, because the size of the imprints obtained by SEM are very similar to the size obtained by the Indenter.

5.00 μm Figure 40. Sink-in effect besides a residual imprint due to the uniaxial compression focused under the Berkovich indenter.

This effect has been widely observed in the residual imprints when the YBCO was textured by TSMG technique and the applied load was higher than 10 mN. This will be discussed widely under next heading related to fracture toughness.

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Fracture Toughness Ceramics are generally brittle, and prone to generation of cracks when indented. Toughness estimation by micro fracture Vickers indentation is a well-known and broadly employed technique in ceramic materials. This technique consists of the application of a Vickers (or Berkovich) indenter at a given load, onto the material, sufficiently high to nucleate cracks at the corners of the imprint, and further measure the crack length produced at the corner of the imprint, c, in order to evaluate the fracture toughness, KIC. In case that very small volumes of material are going to be evaluated obtuse indenters are required (like Berkovich), which do not produce cracks large enough for a correct estimation of toughness fracture. YBCO samples present different mechanisms of fracture, such as crack at the corners of the imprints, chipping, radial cracks inside the imprint, sink-in, and others. When the sample was textured by Bridgman technique method the following observations have been made: -

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Y-211: at applied loads lower than 10 mN, have lateral cracks producing chipping; this effect can be observed in figure 31b. Y-123: when the applied load is lower than 10 mN any fracture mechanism can be observed (see figure 31a). Y-123/Y-211: at 10 mN of applied load, a propagation of crack is only observed at one of the corners; in this case, the fracture toughness cannot be calculated (see figure 31 c). When the applied load was 30 mN, the toughness of YBCO samples textured by Bridgman technique can be obtained. Figure 33a, shows a crack at the corner of the nanoindentation imprints performed at 30 mN. With the crack length and the Palmqvist equation (equation 27), the toughness can be calculated for an applied load of 30 mN. The obtained value is 2.85 ± 0.11 MPa·m1/2. When the applied load was 100 mN, porosity besides the imprint can be observed due to the localized deformation during the nanoindentation test, and also radial cracks inside the imprint (figure 33b). The radial cracks are vertically halfpenny type cracks that occur on the surface of the specimen outside the plastic deformation zone, and at the corners of the residual impression at the indentation site. These radial cracks are formed by a hoop stress and extended downward into the indentation; an example of the radial cracks can be observed in figure 33b.

When the sample was textured by TSMG technique method the following can be observed: -

When the applied load was 10 mN, for Y-211 particles, the same fracture mechanism has been observed. In this case, chipping is the most important mechanism but also exists another mechanism that reduce the toughness of the studied sample; this factor is the porosity near the imprints. For Y-123 and Y-123/Y-211 at the same load, any fracture mechanism has been observed. At this applied load, the fracture toughness has not been calculated (see figure 32).

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When the applied load was 30 mN, any crack propagation have been observed; only one fracture mechanism could be detected, in this case; in the residual imprint internal cracks can be observed, perpendicular to the surface (see figure 32a). Figure 32a, shows that the Y-211 are harder than Y-123, it is due to Y-211 particles, which after the unloading are not deformed. When the applied load was 100 mN, any crack propagation in the corners of the imprint can be observed. In this case a sink-in and porosity near the indentation imprint can be observed, see figure 35b. The sink-in had been predicted with the relation between the final and the maximum penetration depth; this result is in accordance with the value obtained and presented in table 15. Nest figure, shows cracks inside the nanoindentation imprint, and a sin-in near the imprint.

In conclusion, YBCO samples textured by Bridgman and TSMG technique are brittle materials. The Y-123 phase is a soft phase because absorb all the plastic deformation during the nanoindentation imprint. On the other hand, Y-211 is a hard material and it is broken during the indentation. When the indenter was performed at the interface, any fracture mechanisms have been detected. When the applied load was higher than 10 mN, the fracture toughness can be calculated for the samples textured by Bridgman technique. Nevertheless, for TSMG technique, internal perpendicular cracks have been detected. At maximum applied load, 100 mN, a sink-in effect is produced and porosity can be created due to it.

5.00 μm Figure 41. Residual imprint with different fracture mechanism when the applied load was 100 mN and the material textured by TSMG technique.

The fracture toughness obtained for YBCO samples are in concordance with the values published in the bibliography, concretely with Joo et al. for samples with 15% w/w of Ag which value is 2.80 MPa·m1/2, or 10% w/w Ag which present a value of 2.60 MPa·m1/2. Both values have been obtained with single-edge-notch beam. Sheahen et al. present a fracture toughness value of 3.80 MPa·m1/2 for a sample with 10% w/w Ag.

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2.2.2. Elastic Deformation Experimental Conditions The nanoindenter (XP System, MTS) used in this work was equipped with a CSM attachment. The harmonic displacement for the CSM was 2 nm with a frequency of 45 Hz. The test was carried out on two different samples for each technique of texture studied. Figure 42 shows indentation imprints performed when the maximum load was 700 mN on the ab-plane of the monodomain with a spherical tip nanoindenter. Also in this figure also, it can be observed that the Y-211 particles are homogeneously distributed in the textured sample, so that they can be easily identified but not indented separately. It is due to the size of the Y-211. Stiffness versus Displacement into Surfaces Before plotting the indentation stress-strain curves it is crucial to determine the effective elastic stiffness of the YBCO material examined. Plot of S versus displacement into surface, allows to determine if the contact point indenter-sample is good or not. In that case a straight line can be observed, whereas a non linear relation is found when the contact is not good enough.

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Figure 42. OM micrographs of residual impressions performed on the surface of YBCO orthorhombic phase (ab-plane) at room temperature, when the applied load was 700 mN and the sample textured by Bridgman technique.

Figure 43, shows the tendence of stiffness with the penetration depth. This figure gives us information about which stress-strain curves will be suitable or present high scattered values. When S versus h is a straight line, the contact point between indenter and sample is right and therefore the stress-strain curve will be suitable. This implies that S is not affected by pop-ins or plastic deformation. A bad contact point can give us an overestimated strain value, and therefore inaccurate yield strength is obtained. In this case, MTS software must be used in order to correct experimental values. Next figure shows the experimental curves and the same corrected with MTS software, for both samples studied.

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d) Figure 44. Contact point for samples textured by, a) Bridgman technique with bad contact point, b) TSMG technique with bad contact point, c) Bridgman technique with good contact point, and d) TSMG technique with good contact point.

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Determination of the Elasto-Plastic Transition In most materials, and according to the Tresca criterion, when the value of τmax reaches a certain limit (τmax = σys/2, where σys is the yield strength of the material), the material will start to flow and develops a permanent deformation. This criterion has been also used to explain the nucleation of dislocations during pop-in effects [152], yielding the material to make the transition from an elastic to an elasto-plasic response [153]. In this case, mean contact pressure will converge on the hardness of the material. 25

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Figure 46. Load/depth-of-penetration results for YBCO monodomain ab-planes, with a 25 μm spherical indenter, for Bridgman technique at 200 nm of penetration depth.

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However, it has to be taken into account that, if yield strength is not overcome, the slope of the unloading curve will the same as for the loading curve, and the same equation, 16 will hold, see figure 45.The typical load-displacement curves obtained when the YBCO ab-plane is indented with a 25 μm spherical indenter, figure 46, are characterized by the elasto-plastic transition or commonly known as pop-in, indicated in the figure as a change in the slope of the loading curve. If we observe the surface of the imprints with AFM, residual imprint is found, for penetrations depths lower than 150 nm which correlates with the change in the slope of the loading curve (pop-in effect). For higher penetration depth, different residual imprints can be observed by AFM (Figure 47). During the first steps of nanoindentation (elastic regime), the load curve can be adjusted a Hertzian contact (P = C·h3/2); when the applied load is higher than the elasto-plastic transition, a residual imprint can be observed, and the load curves can be adjusted as P = C·h2. In both cases, C is the slope of the load curve in each regime.

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2,2 μm b) Figure 47. AFM image at different penetration depths for YBCO monodomains, textured by Bridgman technique, a) 200 nm and b) 800nm.

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Stress-Strain Curves In the case of spherical indentation, by plotting po (mean contact pressure) against a/R (contact radium/indenter radium) it is possible to determine the point at which there is the transition from elastic to elasto-plastic regime, shown in Figure 46 as a change in the slope for the studied samples. Elastic (Hertzian theory)

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b) Figure 48. Indentation stress/strain curves for YBCO textured by a) Bridgman, and b) TSMG techniques.

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J.J. Roa, X.G. Capdevila and M. Segarra These values are around 3.5 GPa and 3.0 GPa for Bridgman and TSMG, respectively. To obtain the yield strength, next equation has been used:

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σ ys = ⋅ (1 − 2υ ) ⋅ po

(44)

where, σys is the yield strength, υ is the Poisson ratio of the material and po is the mean contact pressure at the point of elastic to elasto-plastic transition (obtained from data in Figure 48). Values obtained for Bridgman and TSMG samples are 700 MPa and 600 MPa respectively. The difference between both techniques can be attributed to the different amount of defects, as well as correlated to data for hardness, Young’s moduli and fracture toughness previously reported.

2.2.3. Kinetic Study by Nanoindentation Technique of TSMG Samples along the CAxis At 450ºc An important phenomenon in bulk superconductors obtained by TSMG, is the formation of cracks due to the inherent brittleness of the Y-123 phase matrix. These cracks form during the texturing of the superconducting monolite and play an important role in the limitation of current flow. As-grown Y-123/Y-211 bulk superconductors prepared by the TSMG process must be further oxygenated in order to transform from tetragonal to orthorhombic phase. Oxygenation at temperatures around 400-450ºC could take between one to two weeks for a 2 cm diameters and 4 mm height sample. Oxygen diffuses through cracks previously formed during cooling in the texturing process. Two types of macrocraks are formed: I). II).

a/b-macrocraks parallel to the a/b-plane and c-macrocraks parallel to the a/c-plane.

While the a/b-macrocraks extend over almost the whole sample, the c-macrocrak length is limited by the a/b-macrocraks spacing [154]. c-axis macrocrack network represents a mesoscopic defect in bulk superconductors obtained by TSMG, and limits the local current density in the sample. Reddy and Rajasekharan [155] studied the inner structure of partially oxygenated Y-123 bulk melt-textured material. During the oxygenation process microcraks parallel to the a/bplanes (a/b-microcracks) appear in the melt processed YBCO (dark lines in Figure 49). Y-211 phase particles (dark), and twins can also be seen. Experimental Conditions Samples used to perform this study were obtained by the TSMG technique. A small NdBCO melt textured grain was used to seed the melt growth process; it was placed at the centre on top of the basal pellet surface before heating and the melt-growth process was applied. From textured sample, a 20 mm x 20mm x 4mm monodomain was obtained that was further cut into four pieces of 10 mm x 10 mm x 4 mm. Afterwards, these four pieces were introduced in a horizontal furnace for an oxygen annealing (99.999% purity) at 450ºC for

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different oxygenation times, 24 h, 48 h, 96 h and 187 h. After oxygenation each piece has been cut following next schema.

a,b c

10µm Figure 49. The microcracks parallel to the a/b-planes (a/b-microcracks) in the melt processed YBCO (dark lines). Y-211 phase particles (dark) and twins can also be seen. (From Diko et al., Superconductor Sci. Technol, 16, 2003, pages 90-93).

Figure 50. Preparation samples scheme.

The applied load used to determine the mechanical properties was 10 mN. At this load, the mechanical properties of Y-123 can be isolated, and the hardness evolution with the

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oxygenation time can be obtained, as oxygen concentration gradients lead to large mechanical stress in the material. The oxygenation diffusion coefficient in the ab-plane is about 104-106 times larger than in c-direction [156]. The structure of all samples was observed by optical microscopy in polarized light after polishing and etching in 1 w/w % HCl in ethyl alcohol. Nanoindentation imprints were observed by FE-SEM. Oxygenation Defects and Macro-microckracking in Melt-Textured YBCO Bulks First, the c-axis section of the oxygenated samples was investigated in order to locate regions of tetragonal phase which can be clearly distinguishable, under polarized light as darker or brighter regions depending on the orientation of the Y-123 crystal axes to the vector of polarized light and/or the mutual position of the analyses and polarizer [157]. The macrocracking and the porosity were the dominant structural changes in the samples after oxygenation. Observation at optical microscope under polarized light revealed a high density of porosity formed during the oxygenation process. Figure 51, shows the optical micrographs of c-axis YBCO samples textured by TSMG technique at different oxygenation times.

Figure 51. Optical micrographs of the c-axis at different oxygenation times, a) 24 h, b) 48 h, c) 96 h and d) 187 h. O, orthorhombic phase and T, tetragonal phase.

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Different defects produced after the oxygenation process can be seen in Figure 52.

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c) Figure 52. Defects observed by FE-SEM after the oxygenation process, a) formation of microcracks around the porosity, b) generation of macrocracks and c) micrograph of the c-axis after oxygenation process.

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If we observe in detail a porosity, we will see a microcrack (Figure 52a). Another problem during this process is the presence of macrocracks,.that have been generated during the cooling in the texture process due to the different expansion coefficients between Y-123 and Y-211, and that growth during the oxygenation process, see figure 52b. Both defects produce an embrittlement of the material of study. Macrocracks can easily form in regions with low Y-211 concentration and 158], because the fracture toughness increases with the concentration of Y-211 particles [159 and 160]. Moreover, Y-211 phase has a lower thermal expansion coefficient than Y-123 phase in the cdirection [161], and Y-211 is therefore under tension during cooling down, and a/bmacrocracking in the Y-211 is enhanced. Length and the spacing of a/b- and c-macrocracks, increase continuously with increasing the oxygenation temperature. At higher oxygenation temperature the a/b- and especially the cmacrocracks spacing is significantly increasing. In contrast, at lower oxygenation temperature (such as 450ºC) the c-macrocracks are very fine and their density is very high. Tensile stress in the c- and a/b- direction are induced by shrinkage of the lattice parameters of the Y-123 phase with oxygen content. These stresses are then responsible for a/b- and c-crack formation in the oxygenated layer. Oxygen may then flow along the c-cracks to the a/c-surfaces, oxygenate the c-crack’s surface and c-crack’s tip and further c-crack propagation. It can be observed in figure 52c, that in the Y-123 material, the network of pores, dispersed in the Y-123/Y-211 skeleton, also supports the orthorhombic transformation. This means that the region around the pores is oxygenated by fast diffusion in the a/b-direction. All these microstructural items are important during the process of oxygenation of YBCO bulk materials. Determination of the Kinetics of Oxygenation by Nanoindentation The hardness of YBCO bulk material has been studied from the oxygenation surface oxygenation until the bottom of the sample along the c-axis. For this study the thickness of the samples used were 4 mm. Figure 55, shows the different nanoindentation imprints along the c-axis of the samples studied.

Figure 53. Nanoindentation imprints along the c-axis in YBCO samples textured by TSMG technique.

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Figure 53, shows a matrix of indentations being the distance between each are 40 μm. The representation of hardness versus the position from the upper surface of the sample, can give us information about the oxygenation process.

Tetragonal

11.0

24 h 48 h 96 h 187 h

Hardness (GPa)

10.5 10.0 9.5 9.0 8.5

Orthormhombic

8.0 7.5 7.0 0

500

1000

1500

2000

2500

3000

3500

4000

Height from the upper surface along the c-axis (μm) Figure 54. Hardness versus height for YBCO samples textured by TSMG technique.

Figure 54, shows the tetragonal-orthorhommbic transition in YBCO samples textured by TSMG technique. In this figure, it can be observed that the transition is function of the oxygenation time. When the sample is oxygenated during 187h, has a little portion of tetragonal phase. The high disperssion presented in this figure is due to superficial defects produced during the oxygenation process. At the end of oxygenation, all samples present a tetragonal structure; in this case, the hardness value is higher than for orthorhombic structure, which has a macro- and microcraks that produce a reduction of the hardness of the YBCO samples. When the oxygenation time was 187h, the transformation has taken place in the first 3000 μm, remaining the rest of the sample with a structure not purely tetragonal. In this case, an interface between tetragonal and orthorhombic phase could be formed. For samples with a height of 4 mm, the oxygenation time needed to complete the transformation is greater than 187 h. Prediction of the Oxygenation Time in YBCO Bulk Materials To predict the oxygenation form a bigger sample, we represent the distance from the surface that has undergone the transition versus the time needed to achieve it.

J.J. Roa, X.G. Capdevila and M. Segarra

Distance from surface along the c-axis (mm)

230

4000 3500 3000 2500 2000 1500 1000 500 0 0

50

100 150 200 250 300 350 400 time (h)

Figure 55. Distance from the surface along the c-axis transformed into orthorhombic phase versus time of oxygenation for YBCO samples textured by TSMG technique.

As can be seen in this figure the time necessary to complete oxygenate a sample with a height of 4 mm is around 380h. Moreover, figure 55 shows two different tendends, a quick oxygenation process through macrocracks formed during cooling in the texture process, and a slower oxygenation process that take place when microcracks are formed, induced by oxygen.

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[142] Leenders, A.; Ullrich, M.; Freyhardt, H. C.; IEEE Transaction on applied superconductivity 1999, Vol. 9, Num. 2, pages 2074-2077. [143] Cook, R. F.; Dinger, T. R.; Clarke, D. R.; Appl. Phys. Lett. 1987, Vol. 51, Num. 6, page. 454456. [144] Fujitomo, H.; Murakami, M.; Koshizuka, N.; Physica C 1992, Vol. 203, pages 103-110. [145] Sandiumenge, F.; Puig, T.; Rabier, J.; Plain, J.; Obradors, X.; Adv. Mater. 2000, Vol. 12, Num. 5, pages 375-381. [146] Verdyan, A.; Soifer, Y. M.; Azoulay, J.; Rabkin, E.; Kazakevich, M.; IEEE Trans. Appl. Supercond. 2000, Vol. 15, Num. 2, pages 3585-3588. [147] Johansen, T. H.; Superconductor Sci. Technol. 2001, Vol. 13, R121. [148] Sakai, N; Murakama, F.; Diko, P.; Takebashi, S.; Yoo, S. I.; Murakami, M.; Adv. Supercond. 1997, X 645. [149] Alford, N. M.; Birchall, J. D.; Clegg, W. J.; Harmer, M. A.; Kendall, K.; Jones, D. H.; J. Mater. Sci. 1988, Vol. 23, Num. 3, pages 761-768. [150] Mukhopadhyay, N. K.; Eatherly, G. C.; Embury, J. D.; Mater. Sci. Eng. A. 2001, Vol. A35. [151] Sangwall, K.; Surowska, B.; Blaziak, P.; Mater. Chem. Phys. 2003, Vol. 80, Num.2, pages 428-437. [152] Lorenz, D.; Zeckzer, A.; Hilpert, U.; Grau, P.; Johansen, H.; Leipner, H. S.; Phys. Rev. B. 2003, Vol. 67, Num. 17, pages 172101/1-172101/4. [153] Gaillard, Y.; Tromas, C.; Woirgard, J.; Phil. Mag. Let. 2003, Vol. 83, Num. 9, pages 553-561. [154] Diko, P.; krabbes, G.; Physica C 2003, Vol. 399, Issue 3-4, pages 151-157. [155] Reddy, E. S.; Rajasekharan, T.; Physica C 1997, Vol. 279, Issue 1-2, pages 56-62. [156] Isfort, D.; Chaud, X.; Tournier, R.; Kapelski, G.; Physica C 2003, Vol. 390, Issue 4, pages 341-355. [157] Diko, P.; Gawalek, W.; Habisreuther, T.; Klupsch, Th.; Görnert, P.; J. Microsc. 1996, Vol. 184, page 46. [158] Diko, P. ; krabbes, G. ; Supercond. Sci. Technol. 2003, Vol. 16, page 90. [159] Okudera, T. ; Murakami, A. ; Katagiri, K. ; Kasaba, K.; Shoji, Y.; Noto, K.; Sakai, N.; Murakami, M. ; Phys. C 2003, Vol. 392, pages 628-633. [160] Raynes, A. S.; Freiman, S. W.; Gayle, F. W.; Kaiser, D. L.; J. Appl. Phys. 1991, Vol. 70, Issue 10, pages 5254 - 5257. [161] Diko, P.; Supercond. Sci. Technol. 1998, Vol. 11, num. 1, pages 68-72.

In: Nanotechnology… Editors: C.J. Dixon and O.W. Curtines, pp. 237-274

ISBN: 978-1-60692-162-3 © 2010 Nova Science Publishers, Inc.

Chapter 6

ZNO NANOWIRE ARRAYS: TEMPLATE-FREE ASSEMBLY GROWTH AND THEIR PHYSICAL PROPERTIES Bingqiang Cao* and Weiping Cai Key Laboratory of Material Physics, Institute of Solid State Physics, Chinese Academy of Sciences, Hefei, 230031, Anhui, China

Abstract In this chapter, growth and physical properties of ZnO nanowire arrays were reviewed. It begins with some general remarks on semiconductor nanowires and basic properties of ZnO. In the second part, different kinds of growth methods that have been applied to grow ZnO nanowires are summarized. Vapor phase methods usually based on VSL or VS mechanism, depending on the presence or absence of a metal catalyst, were discussed in general. Typical solution methods for growth of ZnO nanowires were discussed separately as there is no common growth mechanism that can be applied to describe them. A new template-free strategy based on self-assembly process to grow ZnO nanowire into arrays were emphasized and discussed in detail. The obtained sample qualities were characterized with scanning electron microscopy, transmission electron microscopy, X-ray diffraction and energydispersive X-ray spectrum. The third part deals with the physical properties of ZnO nanowire arrays. Raman spectrum, including resonant Raman spectrum, was applied to test the crystal quality and phonon interaction of ZnO nanowires. Temperature-dependent photoluminescence spectra were measured to probe the intrinsic exciton and defect-related emission process of ZnO nanowires. Field emission properties of such ZnO nanowire arrays were also studied in view of the possible application for flat-panel displays. Some brief conclusions were summarized at the end.

*

E-mail address: [email protected]

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1. Introduction 1.1. Semiconductor Nanowires Nanowires, including nanobelts and nanorods, represent an important and broad class of one-dimensional (1D) nanostructures at the forefront of nanoscience and nanotechnology. Nanowires are typically single-crystalline, highly anisotropic, semiconducting, insulating, and/or metallic nanostructures that result from rapid growth along one direction [1]. They usually have big aspect ratio (length/diameter) and uniform cross section. In contrast to carbon nanotubes, which are another class of 1D nanostructure that have received considerable attention, nanowires have some unique merits [2]. For example, the electronic properties of nanowires can be more precisely controlled by doping, while those of carbon nanotubes are largely determined by the clarity of the graphene layer. It is generally accepted that 1D nanowires provide a good system to investigate the dependence of electrical and thermal transport or mechanical properties on dimensionality and size. Nanowires are also expected to play an important role as both interconnects and functional unit for optoelectrical nanodevices as they represent one of best defined and controlled classes of nanoscale building blocks. Nanowires can be synthesized in single-crystal form by controlling the experimental parameters during growth [3]. Moreover, growths of nanowires with precisely controlled chemical composition, diameter, length, doping, growth direction, and surfaces properties are also possible till now [4]. However, many experimentally achieved growths of nanowires have not got a satisfactory elucidation of the underling mechanism. A well understanding of their growth mechanisms can help to design the nanowire growth more rationally and predictably. Generally speaking, the nanowires are synthesized by promoting the fast crystallization of solid structures along one direction. There are two main reasons for such preferential growth. One is the intrinsic crystal properties of the grown samples and another is the optimized growth conditions. Vapor phase strategy is probably the most widely explored methods for growths of one-dimensional nanowires, including nanorods and nanobelts, especially for semiconductor nanowires. Numerous techniques have been developed to prepare precursors into the gas phase for the next nanowire growth. Among all vapor-phase methods, the so-called VLS (vapor-liquid-solid) or VS (vapor-solid) mechanism seem to be the most successful in growing nanowires with single crystalline in large quality [5]. In comparison with vapor-phase methods, a variety of solution-phase methods have also been developed to grow nanowires. But, among them, there is no a general principle that could be applied to explain the different experimentally employed chemical processes. Many different mechanisms have been adopted in solution-phase methods, such as, the use of template with one-dimensional channels to confine the growth of nanowire [6], the introduction of a liquid/solid interface to introduce the symmetry of the seed [7], the use of selected capping reagents to direct the crystal growth by preferentially absorption on special facets of a new nuclear [8], and the self-assembly of zero-dimensional nanoparticles into 1D nanowires [9]. Although growth of nanowires has gained much intension and great success, it must be acknowledged that such nanowires need to be organized into three dimensions or arrays for the construction of functional nanodevices. The assembly of nanowires into designed hierarchical structures represents a key step for creating arrays of nanodevices. Using a

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patterned catalyst, nanowires can be directly grown into two-dimensional arrays on a solid substrate by applying VLS-based vapor phase methods [10] without complicated lithography. But it still faces many of the traditional constraints of planar growth and device fabrication. For solution methods, electric-field has been used to attract and align nanowires dispersed in solutions due to their highly anisotropic structure and large polarization [11]. Other techniques, such as fluid-flow-directed pattern [12] and Langmuir-Blodgett method [13] have also been developed to organize as-grown nanowires into hierarchical structures. Alternatively, if the nanowires are synthesized under optimized conditions, their growth can be organized into arrays by a self-assembly process. In the self-assembly process, atoms, molecules, particles and other building blocks organize themselves into functional nanostructures driven by energetics of the system [14]. Self-assembly has become a very effective and promising approach to synthesize and assembly nanowires. Nanowires have been widely used to construct many kinds of nanodevices based on their intrinsic properties and nanoscale dimensionality [15]. Here we limited our discussion on semiconductor nanowires. A prototypical kind of such devices is the nanowire field effect transistor (FET). It has been illustrated that silicon, germanium and gallium nitride nanowire FETs show good performance comparable to the best reported for planar devices made from the same material [16]. Due to the large surface ratio, FETs fabricated with single nanowires also show ultra sensitive sensor properties for detecting a wide range of gases, chemicals and biomedical species in both commercial and research applications [17]. Nanowires also represent very promising building blocks for active optoelectronic nanodevices, such as lightemitting-devices (LED) [18] and lasers [19]. Moreover, complex logic gates and even computational circuits have been illustrated and used as basic digital calculators [20]. Recently, nanogenerators that produce continuous direct-current by harvesting mechanical energy from arrays of ZnO nanowire have also been successfully demonstrated [21].

1.2. Basic Properties of ZnO Nanowire Arrays Zinc oxide (ZnO), a typical II-VI compound semiconductor with a direct band-gap of 3.37eV and a relatively high exciton binding energy (60 meV) at room-temperature, displays excellent electrical, optical, catalysis and novel piezoelectric properties [22]. Table 1 gives a summary of the basic physical properties of the ZnO. The most important characteristics of ZnO that needs to be pointed out firstly is that it has a big exciton binding energy of 60 meV, lager than GaN (21 meV), ZnSe (20 meV), which is also lager than the thermal energy (25 meV) at room temperature. This means that exciton emission can be obtained from ZnO crystals more efficiently in the ultraviolet (UV) region. Therefore, ZnO is expected to be used as a material for blue/UV optoelectronics, including LED and even laser diodes in addition to GaN-based materials [23]. ZnO is also theoretically predicted to show magnetic properties when it is doped with magnetic Co, Mn, Fe, V, etc. elements, which could be used for spinelectronic devices [24]. Moreover, ZnO is also a transparent, highly conducting oxide (TCO), when doped with Al, Ga, In, etc., and can be used as a cheaper alternative to indium doped tin oxide [25]. Besides these, ZnO is also a radiation-hard material for electronic devices in space area and a biocompatible material for biomedical applications without coating.

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Bingqiang Cao and Weiping Cai Table 1. Physical Properties of ZnO Properties Crystal Structure Lattice Parameters (300 K) Density (300 K) Rigidity (Mohs) Static dielectric coefficient (/oC) Melting point Thermal conductivity Linear expansion coefficient (/oC) Refractive index Electron affinity Band gap (300 K) Intrinsic carrier concentration Exciton binding energy Effective mass Electron mobility (300 K) Hole mobility (300 K)

Parameters Wurtzite a＝0.32495 nm c＝0.52069 nm 5.606 g/cm3 5 8.656 1975 oC 0.6; 1-1.2 W/m·K a-axis: 6.5*10-6 /oC c-axis: 3.0*10-6 / oC 2.008, 2.029 3.0 eV 3.37 eV (direct) ~106 /cm3 60 meV Electron: 0.24 m0 Hole: 0.59 m0 20~200 cm2V·S 5-50 cm2V·S

It is well known that ZnO is a polar crystal and has positively Zn2+-terminated (0001) and negatively O2-terminated (000-1) polar surfaces, which induce a net dipole moment along the c-axis. Thus the surface energy of the polar {0001} plane is higher than those of non-polar {01-10} and {2-1-10} planes [26]. So preferential growth along the c-axis (<0001>-direction) is energetically favorable. So, due to this structural polarity, ZnO has diverse nanostructures. Many kinds of ZnO nanostructures, such as, one-dimensional (1D) nanowire, nanorod, nanoring, nanohelix, and two-dimensional (2D) nanoplates with different crystal morphologies and physical properties, have been synthesized by various methods [27]. Although many ZnO nanostructures have been reported, it is must be acknowledged that these need to be integrated or to be arrayed in three dimensions to find their practical applications in nanodevices. Among them, due to their structural uniformity and applications in nanoscale devices, ZnO nanowire arrays represent an important step towards realizing nanoscale devices. Till now, ZnO nanowire array based nanodevices, such as, room-temperature UV lasers [28], solar cells [29], gas sensors [30] and nanogenerators [31], have been illustrated. The main problem now that prevents the ZnO films and nanostructures from being used for practical devices is the lack of reproducible p-type conductivity. Despite numerous the reports of p-type ZnO using various growth methods and different dopants (N, P, As, Sb, Li, etc.), a reliable and reproducible high quality p-type conductivity is still lacking. In the meanwhile the advantages of ZnO are being explored and exploited by alternative approaches: such as heteroepitaxy in which n-type ZnO films are formed on other p-type materials, utilizing ZnO as an active layer. Since the growth of nanowires is a strain nearly

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fully-relaxed process and very high structural quality can be obtained on common substrates, ZnO nanowires have a definite potential for achieving stable p-type doping. Primary results of phosphorous-doped ZnO nanowires grown by chemical vapor deposition and pulsed laser deposition methods have shown very promising p-type conductivity [32-34]. More discussions about p-type conductivity of ZnO have been reviewed by several papers [35-37] and will not be included in this chapter. This chapter reviews the recent research activities of ZnO nanowire arrays, which emphasizes the self-assembly growth and the physical properties. The first part is an introduction on semiconductor nanowires and basic properties of ZnO. The second part discusses different methods and concepts that have been applied to grow ZnO nanowire arrays. The third part discusses the physical properties of ZnO nanowire arrays, such as Raman scattering, photoluminescence, and field emissions, where the influence of morphology and defect are emphasized. The final part summarizes this chapter with some brief conclusions.

2. Growths of ZnO Nanowire Arrays 2.1. Growth of ZnO Nanowire Arrays by Vapor Phase Methods

Figure 1. Schematic illustration of the ZnO nanowire growth through the vapor-liquid-solid (VLS) mechanism. (I) metal particles as catalyst; (II) alloying; (III) nucleation; (IV) axial growth.

Many methods have been developed to grow ZnO nanowire arrays by different groups. It can be generally divided into two groups, vapor phase method and solution phase method. Vapor phase method includes physical vapor deposition (PVD), chemical vapor deposition (CVD), metal-organic CVD, molecular beam epitaxy (MBE), metal-organic vapor phase epitaxy (MOVPE), pulsed laser deposition (PLD), etc. The general growth mechanism for ZnO nanowires applied in vapor phase methods is the so-called Vapor-Liquid-Solid (VLS) and/or Vapor-Solid (VS) processes. The VLS mechanism, first found in the early 1960s [38], has achieved great success in recent years for the growth of different semiconductor nanowires, including ZnO nanowire arrays. In this mechanism, the liquid metallic small droplets act as catalysts to adsorb and dissolve reactant species and guide the growth of nanowires. So the catalyst can be used to control the growth of nanowires in terms of diameters and sites on substrates. Gold nanoparticles are the most frequently used catalysts. It

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usually requires a high growth temperature (~ 900 °C) so that zinc vapor can be dissolved into a gold catalyst droplet to form alloy droplet. After saturation, zinc precipitates out from the droplet and is oxidized as the ZnO nanowire grows with preferred crystallographic direction. Another intrinsic feature of the VLS growth method is that at the tips (or bottoms) of the nanowires there are always impurity particles that are undesirable for device fabrication.

Figure 2. (a) Gold nanoparticles arrays prepared by nanosphere lithography (NSL) method; (b) ZnO nanowire arrays grown with pulsed laser deposition method using gold nanoparticles arrays shown in (a) as catalyst on a- plane of sapphire substrate. [After ref. 39, Rahm, et al. Appl. Phys. A, 88, 31 (2007)].

Sometimes, it is also possible to grow ZnO nanowire arrays by VS process without using any metal catalyst [40]. From the viewpoint of growth kinetics, the VS mechanism emphasizes that surface diffusion of reactant species and preferential incorporation at highsurface-energy sites feed and maintain the continuous growth of nanowires. Figure 3 shows the scanning electron microscopy (SEM) image of ZnO nanowires grown with pulsed laser deposition method on ZnO film coated sapphire substrate without using any catalyst. There are clear hexagonal pyramids at the bottom of the ZnO nanowires. It indicates that the zinc and oxygen vapor are absorbed here and diffuse to the tip of the pyramid for the continuous

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wire growth. It is a typical VS growth process. In this method, usually no catalysts are needed and, accordingly, it is not easy to control the diameter and density of the grown nanowires. Table 1 gives some typical methods that use VLS or VS mechanism to grow ZnO nanowire arrays.

Figure 3. ZnO nanowires grown with pulsed laser deposition method on ZnO film coated sapphire substrate without using any catalyst. (a) tiled-view; (b) top-view.

The obtained ZnO nanowires with vapor phase methods are usually of high structural quality. But, besides high preparation temperature and energy-consuming experiment facilities, these vapor phase methods usually face other limitations in terms of sample uniformity, low product yield and substrate choice. For example, some expensive and/or insulting substrates, such as, sapphire, gallium nitride (GaN), and titanium nitride (TiN), are necessary for the epitaxial growth.

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Bingqiang Cao and Weiping Cai Table 2. Vapor phase growth methods of ZnO nanowire arrays

Method

Catalyst

Substrate

PVD

germanium

silicon

CVD

gold

GaN

MOCVD

gold

silicon

MBE

silver

PLD PVD

gold no

silicon oxide silicon nitride a, m, c-sapphire Silicon (100)

MOVPE

no

c-sapphire

CVD

no

TiN/Silicon (100)

PECVD

no

Silicon (111)

PLD

no

Silicon (100)

Source ZnO, carbon ZnAs2 Diethylzinc oxygen Zinc O3/O2 ZnO Zinc, Oxygen Diethylzinc oxygen Zn; Oxygen Diethylzinc oxygen ZnO

Temperature

Reference

900 oC

41

600

42

800 oC

43

600 oC

44

850 oC 500 cC

45 46

400 ~500 oC

47

550 oC

48

700 oC

49

750 oC

50

2.2. Growth of ZnO Nanowire Arrays by Solution Phase Methods Table 3. Solution phase growth methods of ZnO nanowires Growth Method Sol deposition Hydrothermal

Source Zn(NO3)2 ; ZnCl2 ; HMT Zn(NO3)2; HMT

Sonochemical method

HMT; Zinc

Epitaxial Chemical Deposition

Zn(NO3)2; NaOH

Electrophoretic Deposition

Zn(OAc)2 NaOH

Template-assisted Electrochemical Deposition Template-free Electrochemical Deposition

ZnSO4

Zn(NO3)2

Substrate no Glass, Silicon Silicon; Glass; polycarbonate ZnO film/ Glass Anodic Alumina membrane Anodic Alumina membrane Silicon

Resulting Morphology

Temperature

Ref.

Nanorod

78 oC; 100 oC

51

Nanorod array

95 oC

52

Nanorod array

Room tem.

53

Nanorod array

Room tem.

54

Nanowire

50 oC

55

Zinc/ZnO nanowires

30 oC

56

nanowires

70 oC

57

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In comparison with the vapor phase methods mentioned above, another kind of techniques that have been widely used to grow ZnO nanostructures could be generally named as solution method, including hydrothermal method, sol-gel method, chemical bath deposition, hydrolysis method, and electrochemical deposition method (ECD), which are also attractive due to their low cost and high yield. Many groups have adapted different solution strategies to grow different ZnO nanostructures. Among them, nanowire array is also an important kind nanostructure that has been widely studied. Table 3 summarizes some typical solution methods that have been reported to grow ZnO nanowire arrays. For solution methods, it is more flexible and powerful to control the resulting morphologies and properties of the final samples by changing their growth conditions. In this part, we will review and compare some typical solution methods that have been well developed to synthesis ZnO nanowire arrays.

2.2.1. Hydrolysis Growth The first aqueous solution growth of rod-like ZnO crystals was reported in 1990 by Verges et al. [51], where zinc nitride and zinc chloride solutions were hydrolyzed in the presence of hexamethyltetramine (HMT), respectively. Now it is the most successful route that has been widely adopted to synthesis ZnO films and nanostructures with a solution consisting of Zn(NO3)2 and HMT or dimethylamineborane (DMAB). From the viewpoint of chemical reactions, hydroxide ions are formed by the decomposition of HMT or DMAB, and they react with the zinc ions to form ZnO after dehydration. The hydrolysis and condensation reactions of zinc salts can result in 1D ZnO crystals under a wide variety of conditions. In general, low-temperature (< 300 °C) nanowires (or nanorods) growth is possible in slightly basic conditions as divalent metal ions do not readily hydrolyze in acidic media. It has been suggested [58] that these additives (e.g. HMT, DMAB) function, in part, by decomposing during the reaction and increasing the pH to a basic value. The crystal morphology can be controlled by various species in the solution, which act as promoters or inhibitors for nucleation and growth. These species can include the zinc counterion, additives such as amines, and acids and bases. It has been demonstrated that [59], by changing the DMAB concentration, the ZnO morphology evolves from 1D nanorods to two-dimensional nanoplates. ZnO nanorods grow fast along c-axis direction due to the high surface energy of the polar {0001}- plane when the concentration of OH- ions produced by hydrolysis of DMAB is low. While, when the OH- concentration is increased, more OH- ions preferably adsorb on the (0001) plane of ZnO and the growth of the ZnO nanocrystallite along c-axis is partially suppressed. But they can still grow sideways along <2-1-10> directions. So with the OH- concentration increased, the average aspect ratio (high/width) of ZnO nanorods is decreased. Finally, two-dimensional ZnO nanoplates are formed. In this case, the OH- ions act as both precursor and inhibitor for the crystal growth. Although pH may play an important role in this case, it is shown that the species in solution can have a strong effect on the resulting morphology. In general, the obtained ZnO nanowires deposit randomly on the bottom of the reaction vessel or on the immersed substrate, as shown in Figure 4.

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Figure 4. ZnO nanorods and nanoplates grown through hydrolysis method with Zn(NO3)2 (0.5 M for nanorods and 0.1 for nanoplates) and DMAB (0.05 M) as precursor solutions. [After ref. 59, Cao, et al. J. Phys. Chem. C, 112, 680 (2008)].

Later, Vayssieres et al. [60] developed this method and succeeded in growing ZnO nanowire arrays through monitoring of the nucleation, growth, and aging process by means of chemical and electrostatic control of the interfacial free energy. In most cases, homogeneous nucleation of solid phases requires a higher activation energy barrier, and therefore, heteronucleation is easier and energetically more favorable. Indeed, the interfacial energy between crystals and substrates is usually smaller than the interfacial energy between crystals and solutions. Consequently, nucleation takes place at a lower saturation ratio onto a substrate than in homogeneous solution. Epitaxial crystal growth occurs from substrate-generated nuclei along the preferential direction of crystallization, and if the concentration of precursors is high, a condensed phase of single-crystalline rods perpendicular to the substrate is obtained. According to this strategy, ZnO nanorod arrays were obtained on different substrates, such as glass and silicon wafer, with precursor composed of equimolar (0.1 M) zinc nitride and HMT aqueous solution. Moreover, by decreasing the concentration of the precursors while keeping the same 1:1 ratio, the diameter of the obtained ZnO nanorods also decrease likewise by about an order of magnitude, due to the critical diffusion of the monomers and the subsequent limited growth. Microrods with diameter of 1-2 μm were grown at a concentration of 0.1 M zinc nitride and DMAB. While nanorods of 100-200 nm in diameter were grown at 0.01 M. Greene et al. [61] also reported the growth of ZnO nanowire arrays on ZnO nanocrystals coated silicon substrates. They firstly spin-coated a 150-nm-thick ZnO crystal seeds on silicon substrate, which is composed of ZnO nanocrystals with an average size of 5-10 nm. Then ZnO nanorod arrays were grown on such substrate with precursor solution composed of NaOH (0.3 M) and zinc acetate dihydrate (0.01 M). It is also possible to control the aspect ratio of the nanorod by changing the precursor concentration.

2.2.2. Template-Assisted Electrochemical Deposition Recent advances have made electrodeposition an attractive approach for preparation of nanomaterials [62]. When a metallic salt is dissolved into water or other solution it dissociates to form positively charged ions. The solution that contains these charged ions is called as an electrolyte. By passing a sufficient amount of electric current through this electrolyte, one can

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reduce the metal ions to form solid metals or compounds on electrodes. Especially, if the metals are oxidized again when they are deposited, oxide compound will be acquired instead of pure metals. This process is referred in literature to electrochemical deposition (ECD). In common with other relatively low temperature solution synthesis techniques, ECD of semiconductors typically results in a small crystal size. This is clearly advantageous when the goal is formation of nanostructured materials [63]. There are many different ways of electrodepositing semiconductors and different tactics to control crystal size and nanostructure of the resulting samples. Electrodeposition is usually carried out in a three electrode electrochemical cell that contains a working electrode, a reference electrode, and a counter electrode. ECD has some advantages over physical vapor phase deposition methods for growth of nanostructures [64]. (a) Electrochemical reaction usually takes place close to the electrode within the electric double layer, which has a very high potential gradient of 105 Vcm-1. Under these conditions, the reactions often lead to products which cannot be obtained in an ordinary chemical synthesis. (b) The samples are deposited on the electrode, which is also the substrate for the sample growth. So the sample morphologies can be influenced by the nature of the substrate, especially for the interface between the substrate and electrolyte. (c) Growth kinetics can be controlled by the deposition current passed through the cell, while growth thermodynamics control can be excised by the applied potential. (d) The sample composition can be tuned by varying the composition of the electrolyte. (e) ECD is fast and usually performed at low temperature (< 100 oC). (f) The experimental instruments are not complex and inexpensive therefore well-suited for large-scale production. Growth of ZnO polycrystalline films by ECD method was firstly illustrated by Lincot et al. [65] and Izaki et al. [66], respectively, in 1996. As for the growth mechanism, the one-step ECD reactions using Zn(ON3)2 as electrolyte have been generally proposed and widely −

−

adopted [67]. The reduction of the nitrate ( NO3 to NO2 ) in mild acid solution of Zn2+, which results in increase of the pH value near the electrode surface, is crucial. With the increasing concentration of OH-, Zn(OH)2 will form and deposit on the cathode electrode. The deposited Zn(OH)2 will subsequently be decomposed and form ZnO on the substrate at the temperature of 70oC. Another electrochemical route that has also been used to grow ZnO is the oxidation of the zinc ions by introducing oxygen gas when the electrodeposition is performed [68]. An historic overview of work carried out in this field has been summarized elsewhere [69]. Here we limit the present discussion on the ECD growth of ZnO nanowires. Although rod-like films were firstly reported by epitaxial ECD growth of ZnO on GaN [70] and single crystal gold substrates [71], growth of ZnO nanowire arrays by the templateassisted method is the most direct route, where the nanochannels of the template, such as anodic aluminum oxide (AAO) membrane, were filled by electrochemical deposition. In 2002, Li et al. [56] firstly developed a two-step template-assisted ECD method to grow ZnO nanowires. They prepared zinc nanowires by filling the AAO channels with ECD from zinc electrolyte. Then the zinc nanowires were annealed in air. As a result, ZnO nanowire arrays were formed by oxidation of the zinc metal nanowires. The acquired ZnO nanowires were typically polycrystalline and only the typical defect-related green emission were observed from the ZnO nanowires.

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Figure 5. PL spectra of the as-prepared ZnO nanowires (a) and annealed samples (c-d) grown with template-assisted ECD method. (Inset) TEM images of such ZnO nanowires embedded in AAO template. [After ref. 56, Li, et al. Appl. Phys. Lett. 76, 2011 (2002)].

Later, Zheng et al. [72] adopted the one-step electrochemical method to grow ZnO nanowires into the AAO template. The prepared ZnO nanowires are also of low crystal and optical qualities. Another question about the AAO template-assisted growth method is that when the nanowire/AAO composite is dissolved to remove the template, the ZnO nanowire arrays may be destroyed. More important, if the nanowires are embedded in the channels of the AAO template, the nanowires could not be integrated with other circuitry as the aluminum oxide is an insulator. Wong et al. [73] reported that ZnO nanorods can be grown on polycrystalline zinc foil by electrodeposition in an aqueous zinc chloride/calcium chloride solution at 80 oC. It was illustrated that the growth of ZnO nanorods on zinc foil by solution electrochemistry was strongly influenced by electrolyte solution concentration as well as by the substrate electrode. Intrinsic ultraviolet (UV) emission of ZnO was observed from the annealed ZnO nanorods. More recently, Pauporté et al. [74] grew arrays of vertically oriented epitaxial nanorods of ZnO on GaN at low temperature (80 oC) from an electrolytic aqueous bath using electrochemical reactions. The preparation is a template-free and does not use any foreign species. At room-temperature, UV simulated emissions were also measured at 382 nm with an excitation threshold at 4.4 MW/cm2. Due to the big lattice mismatch between the ZnO and silicon substrate (electrode), it is difficult to epitaxially grow ZnO nanowire arrays directly on silicon substrate, which is extremely important for their device applications. Usually, only polycrystalline ZnO films or nanoparticles can be grown on silicon substrate by ECD methods [75]. So preparation of ZnO 1D nanostructure arrays on widely used silicon substrate by a soft and template-free method at a low cost is very appealing for their potential applications in nanodevices. The next parts will discuses the progress in this area.

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2.2.3. Template-Free Electrochemical Deposition a. Substrate-Induced Growth of ZnO Nanowire Arrays [76] Experiment A 50nm-thick gold layer was firstly deposited on Si(100) substrate with low resistivity (<10 Ω·cm). Galvanostatic electrochemical deposition was then employed on the gold coated Si substrates (Au/Si) at a current density of 0.9 mA/cm2. Zinc sheets (99.99% purity) acted as the anode electrode. The electrolyte solution was Zn(NO3)2 aqueous solution (0.05 M). The pH value of the solution was about 6. The deposition temperature was fixed at 70 oC by a temperature-controllable water bath and the deposition time was 2 hours.

Figure 6. (a) SEM image (tilted view) of well-aligned ZnO nanorod arrays grown on Au/Si substrates; (b) XRD spectrum of ZnO nanorods array on Au/Si substrate and (inset) θ -rocking-curves of ZnO(0002) peak. [After ref. 76, Cao, et al. Nanotechnology, 16, 2567 (2005)].

Characterizations Fig.6 (a) is the general morphology of the as-synthesized samples on Au/Si substrate. A layer of well-aligned nanorod arrays with high density has grown on the substrate surface erectly. Typically, the nanorods electrodeposited in 0.05 M Zn(NO3)2

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electrolyte under a current of 0.9 mA for two hours show a mean length of two micrometers with tip of tens of nanometers in diameter. From the XRD pattern of the as-prepared sample shown in Fig.6 (a), only four peaks are observed, which can be assigned to: ZnO(002)/(004), Au(111) and Si(400). The absence of any other peaks confirms the well alignment of these nanorods with c-axis orientation. Meanwhile, the XRD θ -rocking curves of the ZnO (0002)peak shows a full width at half maximum (FWHM) value of 7.5o, which is comparable to that of the ZnO nanowire arrays grown on silicon substrates (3o~10o) or on ZnO films (~6.5o) by vapor phase methods [77-78]. It indicates the alignment of the nanowires is normal to the surface of the Au/Si substrates.

Figure 7. (a) TEM image of a single ZnO nanorod; (b) corresponding HRTEM image and its selected area electron diffraction pattern (inset), showing clearly the (0001) growth direction. [After ref. 76, Cao, et al. Nanotechnology, 16, 2567 (2005)].

TEM images (Fig.7) indicate that the ZnO nanorod is of single-crystal structure without visible defects. It is also clearly shown that the (0001) crystal planes are perpendicular to the axis of the nanorod and <0001> is the preferred growth direction. The corresponding selected area electron diffraction pattern (SAED) pattern further proves the single-crystal nature of the nanorods and their <0001>-growth direction. The EDS result of a single ZnO nanorod, illustrated in Fig.8, indicates that the atomic ratio of zinc to oxygen is about 1.05, which is near the stoichiometric composition of pure ZnO. The carbon and copper signal are from carbon-coated copper TEM grid. Growth Mechanism This template-free ECD growth strategy of ZnO nanowire arrays makes use of the surface properties of the substrate. Working electrodes have an important influence on the initial ZnO nucleation process in the ECD method. It is shown in the SEM image (Fig 9a) of the Au/Si electrode that the gold film is composed of smooth and featureless grains (nanoparticle film) with about 100 nm in size. XRD measurement (Fig.9b) has proved that such gold nanoparticle film has a strong preferential orientation of Au(111) planes parallel to the Si substrate. Obviously, during the initial period of the ECD process, ZnO will nucleate on the flat gold nanoparticles on the cathode. Since the interface energy is directly related to the lattice mismatch of interfaces, to lower the interface energy, there will exist an orientation relationship between the ZnO nuclei and Au film to reduce the lattice

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mismatch, which can be deduced to be (1x1)ZnO(0001)[11-20]//(1x1)Au(111)[-110]. The lattice mismatch between ZnO(0001) and gold(111) planes is 12.7%. So the ZnO nuclei on the substrate should be {0001}-oriented. In addition, as mentioned above, fast crystal growth along the <0001>-direction for ZnO is energetically favorable due to its higher surface energy of the polar {0001}-plane. For thermodynamic stability of growth habit, the growth rate along <0001> direction is faster than other directions, which leads to the formation of ZnO nanowires perpendicular to the substrate. So the ECD growth of the nanowire arrays can be described in two steps: (I) formation of {0001}-oriented ZnO nuclei on the {111}-oriented Au film due to the smaller lattice mismatch between them, (II) and then the oriented ZnO nuclei serve as seeds and grow preferentially along c-axis due to the high surface free energy of the {0001} plane. Thus, the growth of ZnO nanorod arrays on Au/Si substrate through this template-free ECD method is a typical self-assembly process.

Figure 8. EDX spectrum of a single ZnO nanorod on a TEM copper grid and its elements analysis (inset) [After ref. 76, Cao, et al. Nanotechnology, 16, 2567 (2005)].

Figure 9. (a) SEM image of the surface of Au(111)/Si substrate; (b) XRD spectrum of the Au/Si substrate and (inset) θ -rocking-curves of Au(111) peak. [After ref. 76, Cao, et al. Nanotechnology, 16, 2567 (2005)].

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b. Seed Layer-Induced Growth of ZnO Nanowire Arrays [79] Experiment A 50nm-thick layer of ZnO film (seed layer) was deposited on silicon (100) substrates (resistivity < 10 Ω•cm), by radio frequency sputtering methods. Then, galvanostatic cathodic deposition was employed on the ZnO film coated silicon substrates at a current range from 0.8 to 1.25 mA. The deposition area on cathode is all the same in these experiments. Zinc sheets (99.99% purity) act as the anode electrode and the electrolyte solution was zinc nitrate aqueous solution (0.05 M). The pH value of the solution was about 6. The deposition temperature was fixed at 70oC with a water bath and the deposition time was 3 hours.

Figure 10. XRD spectra comparison of the samples electrodeposited under different currents: (a) 0.8 mA, (b) 1.0 mA, and (c) 1.25 mA. [After ref. 79, Cao, et al. J. Phys. Chem. C, 111, 2470 (2007)].

Characterizations All XRD spectra (Fig.10) of the samples electrodeposited under different ECD currents show strong (002) peaks in addition to very weak (101) peaks corresponding to the ZnO wurtzite structure. It indicates that these ZnO samples are all of highly c-axis orientation. Macroscopically, all the whole cathodic substrates were covered with a layer of white and homogeneous ZnO film or particle film in squared centimeter order. But their microstructures show a notable evolution with the increasing deposition currents. Fig. 11a shows the SEM image of ZnO electrodeposited under current of 0.8 mA. The film shows flat and compact surface morphology and seems to be of many hexagonal nanosheets from the top view. They are the (001) planes of wurtzite ZnO crystal. So such film has a corientation, consisting with its XRD result in Fig. 10a. When the ECD current increases to 1.0 mA, large-area well-aligned ZnO nanowire arrays in high density are acquired, as shown in Fig. 11b. These ZnO nanowires are all straight, smooth and relatively vertical to the substrate with uniform diameters of about 100 nm. Their lengths could be easily controlled by the growth time. The XRD spectrum (Fig. 10b) also indicates that all the ZnO nanowires grow along the c-axis. When the current further increases to 1.25 mA, the electrodeposited film

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seems to be composed of white ZnO particles watched with the naked eyes. But SEM images of Fig. 12(a, b) show that such ZnO particles are somewhat like connective microcalabashes, and their shells are made up of densely arrayed ZnO nanowires with sharp ends. These hierarchical ZnO nanostructured arrays are growing nearly vertical to the substrate. One broken microcalabash shown in Fig. 12c indicates that there are ZnO nanoparticles filled inside. The TEM images and SAED pattern shown in Fig. 12d and e prove that such singlecrystal ZnO nanowire also grows along the c-axis.

Figure 11. SEM images of ZnO samples electrodeposited under current of 0.8 mA (a) and 1.0 mA (b). [After ref. 79, Cao, et al. J. Phys. Chem. C, 111, 2470 (2007)].

Growth Mechanism To elucidate the growth mechanism of different ZnO nanostructures in this electrochemical route, the properties of the ZnO seed-layer were firstly checked. It is shown in the AFM image of Fig. 13a that the film is composed of smooth and featureless grains about 20 nm in size. The XRD measurement shown in Fig. 13b indicates that the ZnO seed-layer has a strong c-axis orientation. Since preferential growth along the c-axis (<0001>direction) is energetically favorable for ZnO, the seed-layer will have an important influence on the initial ECD nucleation process. The ZnO seed-layer on the cathode has a c-orientation, where the ZnO(001) planes are parallel to substrate. So, when the thermodynamically favored (0001)-oriented ZnO nuclei deposit on the seed-layer, they will continue fast growth along caxis easily due to the exact lattice match. As a result, (0001)-oriented ZnO films or nanowire arrays are formed, as shown in Fig. 11a and b. The reason for these two different morphologies is that different deposition current will cause different growth velocity ratio along c-axis to a-axis [80]. Obviously, under higher deposition current, the crystal growth velocity along c-axis will be much faster than those of other directions and, therefore, the growth ratio along c-axis to a-axis will be larger. So deposition at low current results in compact (0001)-oriented ZnO film and, when the deposition current is higher, the samples are well-aligned nanowires. If the current is further increased, the crystal growth velocity is faster with a result that the lengths of nanowires are not quite uniform, and some nanowires are obviously longer. When the newly formed ZnO clusters grow on the protruded parts by secondary nucleation, it will lead to the accumulation of ZnO nuclei and then nanoparticles,

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which serves as the cores of microcalabashes, as shown in Fig. 12c. Such ZnO nanoparticles continually grow densely fast along the c-axis and then lead to the formation of prickly ZnO microspheres (Fig.12 a and b), which are composed of nanorods. If such process is repeated, microcalabashes are finally formed by this self-assembly process.

Figure 12. SEM images of ZnO microcalabashes (a, 30o tiled-view) and a single ZnO microcalabash composed of densely-arrayed ZnO nanowires electrodeposited under current of 1.25 mA (b), and a split microcalabash (c). TEM image of a single ZnO nanowire and inset is the corresponding selected-area electron diffraction pattern (d), and its lattice fringe image (e). [After ref. 79, Cao, et al. J. Phys. Chem. C, 111, 2470 (2007)].

Figure 13. AFM image (a) and XRD spectrum (b) of the ZnO seed layers on silicon substrate with an obvious (001) orientation and mean roughness about 3nm. [After ref. 79, Cao, et al. J. Phys. Chem. C, 111, 2470 (2007)].

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3 Physical Properties of ZnO Nanowires 3.1. Raman Scattering Property 3.1.1. Vibration Properties of ZnO Due to its sensitivity to the crystallization, structural disorder and defects in nanostructures, Raman scattering has been widely applied to investigate the crystal quality and vibration properties of the ZnO nanostructures. Wurtzite ZnO crystal has two formula unites per primitive cell and belongs to C 6 v space group. The optical phonons of ZnO with 4

wurtzite structure at the Г-point of the Brillouin zone belong to the following irreducible representation [81] Г= 1A1 + 2B1 + 1E1 + 2E2, where, the branches with E1 and E2 symmetry are two-fold degenerated. Both A1 and E1 modes are polar, and split into transverse optical (TO) and longitudinal optical (LO) phonons with different frequencies, due to the macroscopic electric fields associated with the LO phonons. The short-range inter-atomic forces cause anisotropy, and, therefore, A1 and E1 modes possess different frequencies. As the electrostatic forces dominate the anisotropy in the short-range forces in ZnO, the TO-LO splitting is larger than the A1-E1 splitting. For the lattice vibrations with A1 and E1 symmetry, the atoms move parallel and perpendicular to the c-axis, respectively, as indicated in Fig. 14. Both A1 and E1 modes are Raman and infrared active. The two non-polar E2 modes, E2(high) and E2(low), are Raman active only. The B1modes are both infrared and Raman inactive (silent modes). For crystals with wurtzite crystal structure, pure longitudinal or pure transverse phonons of well-defined symmetry can be observed only if the phonon propagation is along or perpendicular to the c-axis. Group theory in combination with polarization and propagation considerations can identify the symmetry of the Raman active optical modes by applying different scattering configurations. Table 4 compares the theoretical optical phonons of wurtzite ZnO with the experimental data.

Figure 14. Scheme of optical phonon modes of wurtzite ZnO at the Г-point of the Brillouin zone. Black dots: zinc atoms; black circle: oxygen atoms. Vertical direction: c-axis.

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Table 4. Comparison of the experimental data of the optical phonon modes of ZnO with theory predictions (after ref. [82]) Modes A1(TO) E1(TO) A1(LO) E1(LO) E2(Low) E2(high)

Raman(exp.)

Infrared(exp.)

Theory

378 380 407, 409 410, 413 574, 576 579 583, 587 588, 593 98, 99 101, 102

409.1, 408.2 412 574.5 577.1 588.3 592.1 --

126 98

437, 438 444

--

335 433

380

382 386 316 407 548 628

3.1.2. Raman Spectrum of ZnO Nanowire Arrays

Figure 15. Raman spectrum of the ZnO nanowire arrays grown on gold coated silicon substrate. [After ref. 76, Cao, et al. Nanotechnology, 16, 2567 (2005)].

Fig. 15 shows the Raman spectrum of the ZnO nanowire arrays grown on Au/Si substrate by ECD method. The excitation light is Ar+ laser of 514.5 nm. In the Raman experiment, the incident laser light was parallel to the ZnO nanorods and the Raman signal was recorded in the backscattering geometry. According to the Raman selection rules [83], LO and E2 modes are allowed while the TO modes of A1(379 cm-1) and E1(407 cm-1) are forbidden. From Fig.

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15, the strong and characteristic E2(high) mode at 437cm-1 is clearly observed, in addition to a weak second order 2E2 mode at 332 cm-1. The wide peak at around 573cm-1 is too weak to identify its physical origin but it is in the Raman shift region of ZnO LO modes. The peak at 520 cm-1 is the Raman signal from the silicon substrate because of the penetration of the excitation laser light through the nanowire arrays. The peak position of the E2 mode was red shifted about 3 cm-1 compared with that of bulk ZnO [84]. This redshift is usually attributed to the optical phonon confinement effect in nanostructures that can cause uncertainty in the phonon wave vectors, a downshift of the Raman peaks. [85]

Figure 16. (a) Resonant Raman scattering spectrum of the ZnO nanowire arrays excited with He-Cd laser of 325 nm; (b) 1LO and 2LO resonant scattering modes and their Lorenz fitting results. [After ref. 76, Cao, et al. Nanotechnology, 16, 2567 (2005)].

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As we know, ZnO has a direct band gap of 3.37eV and large polaron coupling coefficient. Therefore, if we use the exciting photon energy in the proximity of the electronic interband transition energy, the polar LO modes will be dominant since the resonance enhancement of LO phonons is much stronger than that of the TO modes and other nonpolar phonons, such as, E2. That is, multiphonon resonant Raman scattering (MRRS) process will happen [86]. Fig.16a is the corresponding Raman spectrum of the ZnO nanorod array, excited with 325 nm He-Cd laser. Three major peaks, centered at 571, 1142, and 1723cm-1 respectively, are observed, which superimpose upon the PL background, in addition to the almost imperceptible E2 and 2E2 modes. These sharp lines are the multiples of the frequency of 571cm-1 and can be attributed to the Raman 1LO mode and its overtones. Due to the anisotropic short-range force in the uniaxial ZnO lattice, the 1LO mode usually consists of the contributions from both A1(LO) and E1(LO) (9 cm-1in frequency interval between them). However, Lorentzian curve fitting has revealed that 1LO and 2LO peaks in Fig.16a consist mostly of the A1(LO) mode and the contribution from E1(LO) is neglectable, as illustrated in Fig.16b. Since the E1(LO) mode is ascribed to the defects in ZnO, such as, oxygen vacancies or zinc interstitials [87], absence of E1(LO) mode indicates the good crystal quality of the nanorods. Moreover, the linewidths of the 1LO and 2LO modes are 21 and 35cm-1, respectively, which are much broader than the results predicted by the formula:

Δλ−1 ( jLO ) = 9 jcm −1 , where Δλ－1is the FWHM of the jth order LO mode [88], and the resonant Raman peaks also show obvious asymmetry. These facts are further evidence of the optical phonon confinement effect of nanostructures.

3.2. Photoluminesce of ZnO Nanowires 3.2.1. General Remarks Compared with its competitors, such as GaN and ZnSe, ZnO has higher exciton bound energy of 60 meV. The stable excitons could lead to more efficient UV light-emitting even laser actions. This is the most important physical properties of ZnO as optoelectronic materials, which could be developed for light-emitting devices (LED) and laser diodes. Besides its novel UV emission properties, visible luminescence in range of 400-700 nm has also been widely studied because of its potential applications in new low-voltage field emissive display [89]. So the optical properties from the UV to the visible region are all the basal physical properties of ZnO and, therefore, extremely important for its applications. Although luminescence properties have been the subject of ZnO studies for several decades, the emission centers and mechanisms are still controversial due to their complexity and sensitivity to the surface states, mirror changes of electronic configurations and preparation methods [90-98]. For instance, many controversial viewpoints have been proposed on the origin of the ordinary green emission around 520 nm [93-96]. As for the violet (around 410 nm) and red emission (around 600 nm) bands, there are only few discussions without more proofs [97]. So the full understandings of the different luminescence mechanisms are still a challenging and significant subject in the ZnO optical studies. In this part, intrinsic exciton UV emission and defect-related visible emissions of ZnO nanowires are discussed with emphasis on their luminescence processes.

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3.2.2. Temperature-Dependent Photoluminescence Behavior of ZnO Nanowires Fig. 17 is the PL spectra of ZnO nanowire arrays grown on silicon substrate with gold buffer layer by ECD method, as previously described. These PL spectra were measured over a temperature range from 10 to 300 K under excitation with 267 nm Ti: sapphire laser. There are two distinct peaks situated at green (2.2 eV) and UV (3.3 eV) regions, respectively, and one inconspicuous band superimposed between them for all spectra. With the temperature increasing, the UV peak shows obvious redshift. But the green peak blue shifts and its relative

Figure 17. PL spectra of the as-prepared ZnO nanowires. (a), PL spectra at different temperatures. All spectra are normalized by the UV peak intensity and shifted vertically for clarity. (b), Typical Gaussianfitting analysis of the PL spectrum at 50 K. Three emission bands are named Green, Violet, and UV emission bands, respectively. [After ref. 98, Cao, et al. Appl. Phys. Lett, 88, 161101 (2006)].

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intensity to UV peak decreases. Further analyses have revealed that all spectra can be well fitted by three Gaussian peaks. The inconspicuous band between the green and the UV peaks is mainly contributed from a violet (3.1 eV) emission. Fig. 17b gives a typical Gaussian analysis result for the full PL spectrum at 50 K, where a violet emission band is clearly illustrated. The Gaussian-fitted violet emission bands from 10 K to 300 K are shown in Fig.18. The peak position shows a different style from the other two peaks. As the temperature increases from 10 K to 70 K, the violet peak red shifts by 35 meV, which is much larger than the bandgap shrinkage value (4 meV) over this temperature range predicted by Varshni formula (more discussion later). When the temperature increases from 70 K to 150 K, however, the peak blue shifts by about 160 meV without considering the bandgap shrinkage of 11 meV in this temperature range. Further increasing to room-temperature (RT) induces the peak to redshift again by 31 meV that is similar to the bandgap shrinkage in this temperature range. In a word, the violet emission shows an abnormal S-shaped (or red-bluered) shift behavior with increasing temperature. Fig. 19 compares the temperature-dependent behavior of these three emission peaks observed from the ZnO nanowires.

Figure 18. Gaussian-fitting violet emission bands at different temperatures, the curves from top to bottom correspond to the temperatures shown in the left, respectively. [After ref. 98, Cao, et al. Appl. Phys. Lett, 88, 161101 (2006)].

3.2.3. Physical Origin of Different ZnO Emission Bands For the UV emission of ZnO crystal, it is usually attributed to the near bandgap recombination of electrons and holes in form of excitons. Generally, due to the temperatureinduced lattice dilatation and electron-lattice interaction, the near bandgap emission peak energy follows the Varshni formula [99]:

E g (T ) = E g (0) − α ⋅ T 2 /(T + β )

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where Eg(t) is the band gap at an absolute temperature (T), and α, β are the material-related Varshni thermal coefficients. The solid line in the lower frame of Fig. 19 denotes the fitting results for the UV emission from 10 K to 300 K. The obtained fitting parameters of E0, α and β are 3.379 eV, 7.5x10-4 eV/K and 1050 K, respectively, which are comparable to the values reported for ZnO films [100]. So the well-fitting of redshift with increasing temperature for the UV emission proves its exciton origin. The visible (green and violet) bands show anomalous temperature-dependent behaviors, which differ remarkably from the monotonous bandgap shrinkage with increasing temperature. Therefore, they must be associated with the defect-related levels in the ZnO nanorods. To investigate their physical origins, X-ray photoelectron spectrum (XPS) and electron paramagnetic resonance (EPR) were applied to detect the defect states existing in the ZnO nanowires. In the full XPS spectrum of the ZnO nanowires (Fig. 20a, with sensing depth of 2 nm), there is no any impurity detected besides the Zn, O, adsorbed carbon and gold (from the Au/Si substrate) signals. Fig. 20b presents the high-resolution XPS peak of O1s centered at 531.25 eV with a large FWHM of 2.1 eV, which can be well fitted with three Lorentzian peaks. The three binding energy components, centered at 531.89 eV, 531.13 eV and 530.40 eV, are usually attributed to the absorbed oxygen on the surface of the ZnO nanorods, O2ions in the oxygen-deficient regions within the matrix of ZnO, and O2- ions on the wurtzite structure of the hexagonal Zn2 + ion lattice, respectively [101]. It is also to be mentioned that the intensity of the component at 531.13eV is higher than the other two components, indicating the relatively higher concentration of oxygen-deficient states in the nanorod surface layer.

Figure 19. The temperature-dependent peak position for the three typical emission bands. Dash-line: aid to eyes; Solid line: fitting result by Varshni formula. [After ref. 98, Cao, et al. Appl. Phys. Lett, 88, 161101 (2006)].

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Figure 20. (a) Full XPS spectrum of the ZnO nanowire array; (b) High-resolution spectrum of the O1s peak and its Lorentzian fitting results.

Fig. 21 is the EPR spectrum of the as-prepared ZnO nanowires, which shows a clear gfactor of 1.96. Although this EPR signal was formerly attributed to the singly ionized oxygen +

vacancy Vo

[93], now it has been unambiguously proved that this signal is typically

associated with shallow donors [102-106] and its position appears to be independent on the shallow donor identity, such as, intrinsic Zni [102-103] or doped copper [105], nitrogen donors [106].

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Figure 21. EPR spectrum of the as-prepared ZnO nanorod array showing a clear g-factor of 1.95.

As the XPS data proves that there is no any other impurity detected, therefore, the EPR signal of g~1.96 in the ZnO nanowires corresponds to the Zni shallow donors. It has a 0.2 eV energy gap below the conduction band [107]. Moreover, the O1s peak analysis in the highresolution XPS spectrum shown in Fig. 20b proves the relatively higher concentration of oxygen-deficient states in the nanowire surface layer. The oxygen vacancy is a deep level situating at 0.9 eV higher than the valance band [108]. Similarly, temperature-dependent S-shaped shift behavior has also been observed in nitride semiconductors, which are known to be closely related to the carrier localization caused by alloy composition or inhomogeneities [109-111]. Accordingly, based on these two existing defect levels in the ZnO nanorods, it is proposed that the violet emission (3.03 eV at 10 K) comes from the electron-hole recombination between the electrons localized at the Znishallow donor levels and holes in the valence band as shown in Fig.22(a, b). Its emission energy is well consistent with the theoretical value between these two levels. As the temperature increases from 10 K to 70 K, a redshift appears because the excitons gain thermal energy sufficiently to overcome small energy barriers and are trapped in adjacent lower levels of the localized states from where the recombination takes place (T2 process in Fig.22 (b)). Obviously, this redshift also includes the contribution from the temperature-induced bandgap shrinkage of 4 meV. The following blueshift of 160 meV from 70 K to 150 K is caused by the thermal population of electrons localized at higher energy states and transitions from these higher levels to the valance band (T3 process in Fig.22 (b)). In this temperature range the influence of bandgap shrinkage is totally counteracted by the localization effect. At temperatures over 150 K, the shift of the emission band is mainly determined by the behavior of the bandgap shrinkage, and the shift energy is consistent with the prediction value by the Varshni relation. Therefore, the S-shaped shift of the violet emission is the result of competition between the temperature-induced localization effect and bandgap shrinkage.

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Figure 22. (a) Schematic illustration of the ZnO band structure and the proposed UV, violet, and green emission processes, NR: nonradiative transition; (b): model of the localized electrons at Zni defect levels and their transition process with increasing temperature. (c): model of the green emission with increasing temperature. [After ref. 98, Cao, et al. Appl. Phys. Lett, 88, 161101 (2006)].

As for green emission, theoretical studies have identified the electron-hole radiative recombination at the Vo center as the source of the green luminescence [108]. However, the transition process is still not clear yet. Since it blueshift monotonously and its intensity decreased with increasing temperature, it is proposed that the green emission at near room temperature is composed of two transitions: from the conduction band to Vo levels and from Zni levels to Vo levels as shown in Fig.22 (a, c). In fact, the existences of these two types of transitions have been confirmed by optically detected EPR (ODERP) experiments, respectively [104,112]. With temperature decreasing, more electrons in conduction band firstly relax to lower Zni levels nonradiatively, and then recombine with holes at Vo levels, which results in the green emission shift to lower energy with narrower FWHM, changing from 600 meV at 50 K to 440 meV at 300 K. Accordingly, the green and violet emission intensities, compared with that of UV emission, are increased as more electrons at Zni levels recombine with holes at oxygen vacancy and in valance band.

3.3. Field Emission of ZnO Nanowire Arrays 3.3.1. Field Emission Theory Field emission source are widely used in flat-panel displays. The principle of field emission is based on the application of a high electric field to extract electrons from a metal or semiconductor surface [113]. A high electric field near the emitter can be sufficient to

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lower the surface barrier, so that the barrier height is sufficiently reduced to increase emission substantially. When the electric field is around several voltages per nanometer, the width of the barrier is of order 1 nm, and electrons can escape from the emitter surface by tunnelling, as schematically shown in Fig. 23. This process is field emission (FE). The electrical field of FE plays a similar role to the temperature in thermionic emission, where electrons are emitted from a heated filament. Compared with thermionic emission, field emission has the advantages of less thermal shift, low energy spread, and low operating voltage. FE can be observed even at room temperature, so it is also called as cold FE. Traditional field emission devices require complicated micro-fabrication processes, e.g. electron beam lithography [114]. An ideal field emitter should be highly conductive, very sharp at the tip end, robust, and easy to fabricate. ZnO has a high melting point, low emission barrier, and high saturation velocity. Furthermore, as a widely studied transparent conductive oxide, ZnO can be made both highly conductive and optically transparent from the visible to near-UV range through proper doping. These properties make ZnO nanowires a promising candidate for field emission displays [115].

Figure 23. Scheme of the electron emission process from a metal or semiconductor emitter under an electrical field (Fowler-Nordheim process).

Fowler-Nordheim (F-N) tunneling [116] has been studied extensively in Metal-OxideSemiconductor structures, where it has been shown to be the dominant current mechanism of FE, especially for thick oxides. The basic idea is that quantum mechanical tunneling from the adjacent conductor into the insulator limits the current through the structure. Once the carriers have tunneled into the insulator they are free to move within the valence or conduction band of the insulator. According to F-N theory, the relationship between current density (J) and applied electric field (E) can be expressed as follows:

⎛ − BΦ 3 / 2 ⎞ ⎛ β 2E2 ⎞ ⎟⎟ , ⎟⎟ exp⎜⎜ J = A⎜⎜ ⎝ βE ⎠ ⎝ Φ ⎠

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where A=1.54x10-10 (AV-2eV), B=6.83x109 (Vm-1eV-3/2), and Φ is the work function of the emitter. β is the field enhancement factor defined as: E local = βE = βV / d , where Elocal is the local electric field nearby the emitter tip. β can be expressed ideally as

β =h/r, where h is

the height of the emitter (for example, nanowire) and r is the radius of curvature of the tip apex. It can be easily deduced that the physical geometry of one-dimensional semiconductor nanowire provides a field enhancement. Therefore, electrons can be extracted more easily from nanowires and they are expected to be good emitter. According to F-N equation, the plot of Ln( J / E ) Vs E 2

−1

(F-N plot) should be a straight line and is a good check on the field

emission mechanism. Experimentally, FE properties are measured in a vacuum chamber at room temperature under a two-parallel-plate configuration, as schematically shown in Fig. 24. The distance between the emitter and the current detector need to be precisely controlled. Besides this, the emission area also needs to be calculated accurately.

Figure 24. Illustration of the FE measurement from a single nanowire emitter.

3.3.2. Field Emission of ZnO Nanowire Arrays Fig. 25a shows the current density versus applied voltage (J-V) curves for the ZnO nanowire array shown in Fig. 6 measured at various distances (d). The turn-on field, defined as the applied field to draw an emission current of 10 μA/cm2, is 22.1, 19.3 and 18.9 V/μm, corresponding to the d values of 100, 150 and 200 μm, respectively. They are lower than the turn-on value (18 V/μm for 0.01μA/cm2) of ZnO nanoneedles grown on gallium-doped ZnO film by vapor phase method [117] and comparable to that of the silicon nanocone arrays prepared by ion sputtering [118]. The current density as high as 1 mA/cm2 corresponds to the applied voltage of 4.3 kV (Fig. 25a) when d=150 μm, which is comparable to those of other wide-band-gap semiconductor, such as GaN nanowire [119] or AlN nanocone [120] arrays grown on silicon substrates.

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Figure 25. (a) Field emission curves of the ZnO nanowire array measured at the different distances between the anode and cathode; (b) F-N plots corresponding to (a)

Corresponding F-N plots transformed from data of Fig. 25a for different d values are shown in Fig. 25b. All the F-N plots show nearly straight-lines with slightly different slopes, indicating the field emission process from the ZnO nanowire array is an F-N quantum tunnelling process. The observed kinks on the F-N plots, which are located in the field region of turn-on voltage, can attributed to the current-induced changes in the structures of the field emitters [121] or absorbates-induced emission saturation [122]. The slopes obtained from the F-N plots can be used to estimate the β values, as shown in Fig. 25b. They are comparable to the value of multi-wall carbon nanotubes [123] (830 at d=125μm). Such well-aligned ZnO nanowire array, with low preparation cost and wafer scale, could promise the industrial application of flat display in a near future.

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3.3.3. Field Emission Comparisons of ZnO Film and Nanowire Arrays

Figure 26. Field emission comparison of ZnO film and nanowire arrays measured at the same conditions: (a) J-E curves and (b) corresponding F-N plots.

The successful synthesis of ZnO films and nanostructures with different morphologies by ECD method paves the way to study its FE property contrastively. Fig. 26 shows the emission current density versus applied electrical field (J-E) curves for ZnO film, nanowire array and hierarchical nanowire array (H-array) corresponding to Fig. 11~12. All the ZnO film, ordered nanowire array, and H-array show steady FE properties at various measuring distances (d) between the electrodes. For simple comparison, Fig. 26 (a) shows the J-E curves measured at the same conditions (d=200μm). The turn-on fields, defined as the applied field to draw an emission current of 10 μA/cm2, are 16.9, 15.5, and 9.5 V/μm, for ZnO film, dense nanowire array, and H-array, respectively. From the data in Fig. 26(a), according to the F-N theory, corresponding F-N plots of

Ln( J / E 2 ) Vs E −1 for different ZnO are shown in Fig. 26(b). All the F-N plots are nearly

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straight-lines with different slopes, indicating the field emission process from the ZnO film and nanowire arrays is consist with F-N mechanism. There is a kink on the F-N plot of ZnO film, which is located in the region of turn-on field, maybe caused by the absorbates-induced emission saturation of the film surface. On the other hand, the obvious nonlinearity is also observed from the F-N plots, indicating that local electric field enhancement factor ( β ) is not exclusively determined by the emitter radius. The β values obtained from the slopes of the FN plots have been shown in Fig. 26(b). In the normal electron emission region, the FE properties of ZnO nanowire arrays, such as turn-on electric field and β factor, are both better than that of ZnO film due to their high aspect ratio and small curvature radius, which are all efficient for FE as emitters. But the β value of the ordered ZnO nanowire array with high density is much smaller than that of the hierarchical ZnO nanowire array, which is caused by the local electric field screening effect around the tips of nanowires [124]. This result indicates that, for practical field emission display application, the planar density of semiconductor nanowire array needs to be optimized to get efficient emission current.

4. Conclusion In this chapter, a review on growth and physical properties of ZnO nanowire arrays are presented. Among various growth techniques for ZnO nanowires, all kinds of vapor phase methods based on VSL or VS mechanism, depending on the presence or absence of a metal catalyst, are all widely adopted probably due to the simplicity of the growth. In general, vapor phase method is proved to be a high-temperature synthesis process with low productivity. On the contrary, many solution methods, which are typical low temperature synthesis process with high productivity, have been proved well suitable for growth of ZnO nanowires. Among them, template-assisted growth is the simplest in principle but may have difficulties in producing high quality single-crystal ZnO nanowire arrays. To overcome this point, a new template-free strategy is developed to grow ZnO nanowire arrays. The growth mechanism can be attributed to formation of {0001}-oriented ZnO nuclei on the designed substrates and then faster growth along <0001> direction in terms of minimum interface and surface energy, which is a template-free self-assembly process. As for physical properties of ZnO nanowire arrays, Raman, photoluminescence and field emission are studied with the samples gown by this template-free growth strategy. Raman studies have confirmed that the ZnO nanowires are of good crystal quality. Resonant Raman scattering spectrum were also observed and applied to study the phonon interaction of ZnO nanowires. Detailed photoluminescence spectra of ZnO over the temperatures from 10 to 300 K show three typical emission bands. They are ultraviolet, violet, and green emission, respectively. With increasing temperature, these bands show different temperature dependences: a normal redshift for the ultraviolet emission, S-shaped shift for the violet emission and blue shift for the green one. The origins of these three bands and their temperature dependent shifts are explained based on defect levels (zinc interstitial and oxygen vacancy levels) and electron localization effect at the defect levels in addition to bandgap shrinkage. ZnO nanowire arrays with optimized planar density also exhibit good field emission properties due to their physical geometry. Such wafer-scale ZnO nanowire arrays

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with good field emission performance and low preparation cost could promise the industrial application of flat display in a near future.

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In: Nanotechnology… Editors: C.J. Dixon and O.W. Curtines, pp.275-292

ISBN: 978-1-60692-162-3 © 2010 Nova Science Publishers, Inc.

Chapter 7

SPATIALLY RESOLVED CONTROL OF ELECTRICAL RESISTIVITY IN ORGANIC MATERIALS — DEVELOPMENT OF A NEW FABRICATION METHOD OF JUNCTION STRUCTURES Toshio Naito Division of Chemistry, Graduate School of Science, Hokkaido University, Hokkaido, Japan

Abstract Organic materials (OMs) are diverse and are interesting in terms of application in electronic devices. In particular, organic charge transfer salts (OCTSs), which are typically composed of positive and negative ionic (and often radical) organic molecules, attract continued attention. Without carrier doping, they generally have high conductivity, magnetism, and well-defined unique nanostructures in their crystalline form. In order to apply the OCTSs to electronic devices, they should be made junction structures. Although there are established ways and advanced methods for doping and fabrication of junction structures in the current industrial techniques for the silicon devices, few of them are applicable to the OMs due to totally different chemical and physical properties between inorganic and organic materials. In this comment, we would like to discuss a new method for simultaneous realization of doping and junction structures beginning with the single crystals of the OCTSs. The method utilizes a photo-induced chemical reaction, and produces a stable solid state composed of well-defined different parts of different conducting/magnetic properties. With reference to our recent and previous work as well as related studies of other groups, discussion will briefly cover experimental methods, preparation of materials, examination of irradiation conditions and resultant solids’ characterization, outline of mechanism of this photochemical modification, and remaining problems to be explained or overcome.

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Introduction Why Organic Materials for Electronic Devices? *1

Molecular or organic materials (OMs) now attract the growing interest of a wide range of researchers for application in electronic devices such as transistors and displays [1,2]. Their advantages over other materials are practically unlimited choice of compounds, accumulated knowledge of material design along with simulation/calculation of the properties using computers, and possible low cost due to industrial large-scale production. Ambient conditions in temperature and pressure is required or adequate for dealing with them, which makes many of the facilities and the procedures for them simple and inexpensive. In addition to experimental techniques such as lithography, one can use wet chemistry, i.e., a variety of procedures of moving, removing, purifying, mixing, and forming desired/undesired organic substances in solution. The products equipped with organic devices can be small or large, depending on their purposes, but anyway they are expected to be light, thin, flexible, and can be suitable for portable equipment [3–5]. Low cost may provide disposable electronic products in the future; in such cases organic devices will not seriously suffer from fragility. Organic electroluminescence (EL) displays, for example, have recently been in practical use in cellular phones. It is also advantageous in usage of the OMs that nanowires and other functional nanostructures are often easily obtained as bulk products owing to their controllable self-assembly properties and well-established synthetic chemistry. Since nanowires, nanotubes, and nanoparticles are considered to be key materials for future electronic and bio-devices such as sensors, the OMs will keep attracting strong attention for the next decade or more.

Beginning with Thin Films A substantial body of research on development of organic devices utilizes various OMs as thin films on appropriate substrates [3–5]. This trend is partly because they are comparatively easily made thin films of desired shapes and dimensions, while it is generally extremely difficult to make their single crystals, like silicon and gallium arsenide, with sufficiently large dimensions for industrial use. Additionally, the chemical, thermal, and mechanical properties of the OMs are quite different from those of silicon and other typical inorganic semiconductors; the OMs do not stand severe conditions such as elevated temperatures and ion beam bombardment required in the current device fabrication processes for the inorganic semiconductors. As a result most organic devices are commonly studied and fabricated on the basis of their unique processes (sublimation and/or various methods of wet chemistry) partly in combination with the current silicon technology (lithography and so on). For example, different kinds of organic and inorganic materials are layered on a plastic substrate using ink jet, vacuum sublimation and other techniques [3–5], functioning as a fieldeffect transistor (FET). Each different kind of material plays its own role such as an insulator, a dielectric, an electrode, and so on. Such an assembled structure with interfaces between

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materials with qualitatively different electrical properties is called a junction structure and is necessary for any kind of device function. The current device fabrication technology has been established mainly upon doping methods to control the electrical properties of the materials in addition to the lithography techniques to make elaborate junction structures by combining different materials.

Different Strategies Required Besides the differences in chemical, thermal, and mechanical properties, it is the high anisotropy (low dimensionality) mentioned in detail below that we must keep in mind as an intrinsic difference between inorganic semiconductors and OMs. This means that electrical conduction is seriously disturbed in a disordered/incoherent arrangement of molecules/domains/grains such as amorphous, thin films, powders and so on. In particular, electrical properties generally depend more on the grain boundaries than on the intrinsic electrical properties inside each grain in highly conductive materials. In other words, in a given OM, the electrical conduction is higher and better-defined in a single crystalline state than any other solid state. The ill-defined electrical conduction pathways in a disordered molecular assembly result not only in extrinsically higher resistivity but also in ambiguous electrical behavior. For example, one might observe a much smaller current ratio between on and off states with higher leak current in an FET than it should have. Two sheets of thin films of the same material never exhibit identical electrical behavior, because their mesoscopic structures, i.e., the sizes and the distribution of grains in the films, differ from each other. Therefore utilization of single crystals may improve or modify the performance of many kinds of organic electronic devices based on thin films. Now we have a problem; how can we fabricate devices from single crystals instead of thin films? The OMs in crystalline forms have more difficulties in processing than inorganic materials and thin films of the OMs. Examples include small dimensions (≤ 1 mm), and low plasticity (every crystal has its own shape and is fragile). In particular, the crystalline organic charge transfer salts (OCTSs) mentioned below additionally suffer from low solubility (typically almost insoluble) and low vapor pressure (practically non-volatile). These properties lead to the following problems mainly on doping methods when we apply the known strategies to them. 1) Few of the standard doping methods of the current silicon technology can be applied. 2) The only known doping method applicable to the OMs is simple mixing of foreign chemical species in the synthesis, yet it is very difficult to control precisely the dopant content, homogeneity, and the part to be doped using the chemical method. 3) Inclusion of dopants might seriously lower the crystal quality and/or might cause unexpected changes of the crystal structures, both of which certainly deteriorate the resultant device performance. 4) One can not use the device fabrication methods suitable for organic thin films based on the wet chemistry or sublimation.

*1

Molecular materials may include metal complex molecules and clusters as well as organic molecules. In this chapter the two terms, the molecular materials and the organic materials, are used in an identical context and not distinguished unless such usage is misleading.

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In this chapter we will discuss the problems of the crystalline OCTSs in the fabrication process of electronic devices, and will propose a new strategy. It involves spatially resolved control of electrical resistivity of molecular crystals using photochemical solid state reactions, which practically transforms the single crystals to junction structures.

Discussions π-Conjugated Planar Radical Ion Molecules The OMs can be divided into groups in terms of their component molecules: neutral or ionized molecules, those with or without unpaired electrons, π-conjugated planar molecules or not, and so on. Hereafter we will focus our attention on the OMs including π-conjugated planar radical ion molecules. Such materials are generally called charge transfer salts and often have high conductivity, magnetism, and well-defined nanostructures, especially in their crystalline form. Compared with the inorganic semiconductors and the OMs of neutral molecules, the OCTSs have unique advantages for electronic devices as follows. a) multi-component: different chemical species can simultaneously play different roles in a single material, which may enable multi- or advanced functions of resultant devices. b) anisotropy: different (usually incompatible) electrical properties appear in different directions. This and the next features (b and c) can be favorable for many kinds of organic devices, because the electric current is expected to flow through some particular parts in given directions and should not flow otherwise. c) nanostructure: in a given crystal structure, the electrical conduction pathways can be clearly defined within an atomic scale. d) low concentration of the carriers: many OCTSs have relatively high conductivity in spite of lower concentration of carriers than naturally occurring metals (about oneorder lower number density per unit volume). This and the next (d and e) features enable effective control of carriers or conductivities by applying electric field and thermodynamic methods. e) high sensitivity to perturbations: their electrical behavior markedly varies under external fields and different thermodynamic conditions. f) correlated electrons: many of them have strongly correlated electron systems, which often lead to superconductivity, ferromagnetism, and other unusual as well as useful physical phenomena. Here we should take a look at an example of the OCTSs; M(R1,R2-DCNQI)2 (R1,R2DCNQI = 2-R1-5-R2-N,N’-dicyanoquinonediimine, R1,R2 = halogen atoms and organic substituent groups, abbreviated as DCNQI below; M = cations of lithium, silver, copper and so on) [6-8]. Figures 1 and 2 show the molecular structure of DCNQI and the crystal structure of M(DCNQI)2, respectively. In the crystal the planar DCNQI molecules exist as radical anion species, and stack to form a one-dimensional (1D) columnar structure. Such structure can stabilize the unpaired electrons, which are originally delocalized within each DCNQI molecule, in the most effective way by fully delocalizing them within each DCNQI column. Except for the case of Cu(DCNQI)2, the unpaired electrons cannot travel between columns

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through the cations. As a result this material exhibits metallic conductivity only in the stacking (// c axis) direction, and behaves as a dielectric (or an insulator) in other directions (thus called 1D conductor). Such a structure can be regarded as a single crystal of nanowires composed of the DCNQI molecules. In this case, the cations do not play a role in the electrical conduction except for insulation of the neighboring columns. However, in the case of M = Ag it was thought, by analogy with photography, that the conduction (unpaired) electrons might be reduced in number by photo-induced charge- (electron-) transfer from the DCNQI0.5- radical anions to the Ag+ cations. This phenomenon may enable direct control of the electrical resistivity. Further detail is discussed below. N C N R1

R2 N C N

Figure 1. Molecular structure of R1, R2-DCNQI (R1, R2 = halogen atoms and organic substituent groups).

Figure 2. Continued on next page.

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φ ~ 0.5 nm

Figure 2. Crystal structure of Ag(DBr-DCNQI)2, where DBr-DCNQI = 2,5-dibromo-N,N’dicyanoquinonediimine. Views down along the c axis (upper), and perpendicular to the c axis (lower). Thin lines indicate the unit cell. Colored spheres indicate silver (pink), nitrogen (light blue), carbon (grey), hydrogen (white), and bromine (red) atoms, respectively. White vertical arrows in the lower figure indicate the conduction pathways through the DBr-DCNQI columns (// c axis) of ca. 0.5-nm thick. Other related salts designated as M(R1, R2-DCNQI)2 (M = metal and organic cations) take isostructures with the structure shown herein.

Basic Problems to be Solved For making an electronic device from a single crystal of a given OM, two basic problems are to be solved. 1) How can we obtain a suitable single crystal with a desired crystal structure and a desired electrical property? 2) How can we make a junction structure from the single crystal? The first problem appears to require a fine and accurate control of the crystal structure as well as precise prediction of the electrical property of the resultant material, either of which is impossible at present. Fortunately, instead of developing a novel material with perfectly desired crystal and electronic structures, one can select a known material with an appropriate crystal structure [9] only if one can control its electrical property. The second problem can be solved in some ways. The straightforward way is to carry out the similar procedures with those for the organic thin film devices. For example, combined

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with other materials, single crystals of OMs can be actually made FETs [10,11]. Here, we would like to discuss a different strategy, which can simultaneously solve both problems [12]. Using photochemical reactions, it is thought that a part of a single crystal could be modified in the electrical properties, remaining the rest of the crystal intact (Figure 3). In other words, an organic single crystal itself can be made a junction structure after irradiation on a part of the crystal. This method, tentatively named “optical doping”, can be performed under mild conditions and suitable for the OMs.

Figure 3. Schematic presentation of “optical doping” on the single crystal of Ag(DM)2. Arrows with broken lines in the lower figure indicate the conduction pathways through each nanowire.

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Photochemical Control of Resistivity—Initial Idea By way of illustration, details are explained with reference to Ag(DM)2 (DM = DMeDCNQI = 2,5-dimethyl-N,N’-dicyanoquinonediimine). In order to realize the idea above, the crystal must contain the following two kinds of chemical species. The first species (Ag+) are inert without irradiation, but are reductive or oxidative under irradiation. The second species (DMn-; 0 ≤ n ≤ 2) are multistep redox systems like transition metal atoms, where some different oxidation states are (meta-)stable to occur. When the material composed of the two species is exposed to light of appropriate wavelengths, a photo-induced electron transfer (redox reaction) is expected to occur between Ag+ and DM0.5-, and the numbers of valence electrons of both species should change. In Ag(DM)2 the electrical conduction is dominated by the unpaired electrons on the DM radical anions, and the Ag+ ions may become strongly oxidative only when they are irradiated with ultraviolet (UV) and/or visible (Vis) light. What is more, the coordination bonds between Ag+ and DM0.5- prove that there is interaction between them. Therefore UV-Vis irradiation is expected to change (reduce) the number of unpaired electrons on the DM columns, and thus should change (increase) the resistivity only at the irradiated part of the sample. Because the resistivity of Ag(DM)2 is metallic, i.e. low (~ 0.01 Ωcm above 100 K) before irradiation, it may be changed to a wide range of values by this method. Based on this idea, we tried to arrange for single crystals of this kind of salts to be continuously irradiated with UV-Vis light.

Main Experimental Results The crystalline and powder samples of Ag(DM)2 were synthesized following the reported procedures [6-8]. The detail of irradiation conditions, measurement and analysis equipments are described in previous reports [12-14]. Various irradiation conditions (temperatures and atmospheres during irradiation, duration, wavelengths and intensity of light) were examined [15]. The results clearly depended on the intensity and the sample temperature, however, did not depend on the wavelengths or on the atmospheres. All the samples in this study were stable under room and spectroscopic lights of various wavelengths with a standard intensity. They were also stable in air and in water. The pristine Ag(DM)2 was thermodynamically and chemically stable below 155ºC, at which temperature it irreversibly turned into amorphous solid retaining the original chemical formula. Whether there were thermal and/or irradiation damages/decompositions on/in the samples were examined using matrix-assisted laser desorption ionization (MALDI) mass, infrared (IR), UV-Vis, solid state NMR, Auger electron, and X-ray photoelectron spectra (XPS), X-ray absorption fine structure (XAFS), elemental analysis, differential scanning calorimetry (DSC), thermogravimetry differential thermal analysis (TG-DTA), scanning electron and atomic force microscopes (SEM, AFM). No damage/decomposition was detected in the samples of the data below. The resistivity of the single crystals was measured with a standard 4-probe method along the stacking (c) axis. All the spectra were measured on the powder samples; the single crystals of Ag(DM)2 were finely ground in an agate mortar and irradiated in a similar manner with that the single crystals were irradiated. Figure 4 shows the irradiation effect of the resistivity of the single crystal of Ag(DM)2 using UV light of 250-450 nm (filtered from 200 W Hg/Xe lamp equipped with a multimode

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quartz optical fiber (length = 1 m, core diameter = 5 mmφ)) [14]. The resistivity measurement and irradiation were simultaneously carried out in air. The temperature was measured with a silicon diode thermometer (LakeShore DT-470-SD fixed with Apiezon N), which was also put in the UV light immediately beside the sample on a brass plate and thermally in contact with it. The resistivity increased in accordance with irradiation time. Sufficiently long interval (75 min) of the irradiation let the sample cool down, which enabled quantitative estimation of the effect of heating involved with the UV irradiation. After ca. 10000 s (~ 3 h) the sample temperature became nearly constant (~ 45ºC). Accordingly, nearly linear increase in resistivity after ~ 10000 s can be attributed to an irreversible photochemical reaction. Similar experiments were carried out in liquid nitrogen using a different light source (200-1100 nm, 5 W tungsten-halogen and 30 W D2 lamps) and a quartz optical fiber (length = 1 m, core diameter = 800 μmφ) (Figure 5) [12]. It should be noted here that only a part of the sample (a single crystal) was irradiated in order to examine whether the electrical properties can be modified only at the part of the crystal. The temperature dependence of resistivity before irradiation (the pristine sample) exhibited metallic behavior from room temperature (RT) down to~ 100 K, where a metal-to-insulator (MI) transition occurred and resistivity rapidly *2 increased below this temperature. After irradiation the resistivity clearly increased; the longer the irradiation became, the more the resistivity increased [12,16]. The observed temperature dependence of the resistivity can be explained as a sum of the two different types of behavior: one is semiconducting behavior, of which resistivity increases with decreasing temperature, and the other is metallic behavior, of which resistivity decreases with decreasing temperature. The behavior around RT and MI transition was nearly retained after irradiation, which means that a part of the sample retains its original electrical behavior after irradiation. This is consistent with the fact that the sample was only partially irradiated. In order to confirm that the silver ions are reduced after irradiation and also to confirm that the non-irradiated part of the sample retains its original chemical state, only the center of the pressed powder pellet (10 mmφ) of Ag(DM)2 was irradiated for four days through the abovementioned (800 μmφ) optical fiber with putting the end of the fiber directly upon the sample and then examined with XPS by scanning along a diameter of the pellet and measuring the spectra at every 100 μm [12]. The result showed that the formal charges (oxidation states) of C, N, and Ag atoms in this material changed after irradiation only at the irradiated part of the sample with a clear borderline, while the rest of the pellet retained its original chemical state. On the surface of Ag(DM)2 in the irradiated part all the silver ions were reduced to be metallic silver, while the carbon and the nitrogen atoms in DM were *3 oxidized [17]. This method has a spatial resolution. Simply by covering half of the single crystal with Al foil during irradiation, a sharp borderline between the irradiated and the nonirradiated parts was observed in Ag(DM)2 as well as in Ag(DI-DCNQI)2 (DI-DCNQI = 2,5diiodo-N,N’-dicyanoquinonediimine) (Figure 6). Identical spectra (XPS) were obtained after *2

*3

Due to fluctuation and metal instability characteristic to the low-dimensional conductors, an MI transition was observed at ~ 100 K with fluctuation (gradual increase in resistivity with decreasing temperature) below ~ 180 K. In this case the silver ions in the pristine sample were fully reduced to be bulk silver and the irradiated part (surface) of Ag(DM)2 was decomposed. This turned out to be due to the elevated temperature (≥ 155ºC) of the sample during irradiation. After the examination of the photochemical products under various irradiation conditions, it proved that one could control the resistivity without decomposition by retaining the sample temperature around RT.

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half a year using the same pellet kept at RT in air, which means that the change is irreversible and the resultant state is stable without particular treatment. The formal charge (n) of DMnwas found to change from n = 0.5 to n ~ 0.35-0.40 progressively with irradiation on the basis of the Raman spectra [12].

Figure 4. Change in the resistivity of Ag(DM)2 during UV irradiation. (a) Overview, and (b) Close view of the initial stage [encircled part of (a)]. The irradiation commenced at ca. 700 s after the resistivity measurement started. Before the irradiation (< 700 s) the resistivity was constant within experimental error. When the UV light was turned on, the resistivity suddenly increased (Δρheat). This could be attributed to heating effect involved with the irradiation. Reproduced and modified from Figure 2 in ref [14] with copyright permission.

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Figure 5. UV-VIS illumination effects on the electrical behavior of the single crystal of Ag(DM)2. a) Temperature-dependence of the resistivity of Ag(DM)2. b) Top: Schematic description of the resistivity measurement by a standard 4-probe method after illumination. The curves with arrows indicate possible different current routes. Bottom: Conceptual section views of the partially illuminated single crystal taking domain structures with different electrical properties at different temperatures. Reproduced and modified from Figure 1 in ref [12] with copyright permission.

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Figure 6. (Upper) Scanning electron micrograph (SEM) of single crystals of Ag(DM)2 after partial irradiation with UV-Vis light using Al foil as a mask. Because of the clear difference in electrical resistivity of the two parts, one can distinguish them owing to the difference in charge-up effects under electron beams. (Lower) Optical micrograph of Ag(DI-DCNQI)2 (DI-DCNQI = 2,5-diiodo-N,N’dicyanoquinonediimine) fixed on the conducting carbon tape after partial irradiation with UV-Vis light using Al foil as a mask. The irradiated part turned bluish hue, while the masked part retained its original brownish reflection. Because of the clear difference in electronic structures of the two parts, one can distinguish them owing to the difference in reflectance spectra. (Inset of lower figure) Atomic force micrograph (AFM) of the area (40×40 nm2; around the yellow square) on the borderline of the irradiated and non-irradiated parts. The AFM shows as fairly smooth surface as the crystalline solid should have, and in fact roughness of the surface on the borderline is less than 1 nm, which suggests that the irradiation does not cause mechanical/thermal damage on the surface of the single crystal.

Now that this photochemical modification of resistivity is found to serve as a chemical fabrication method of junction structures, we should discuss whether a prototype of some device can be made. After irradiating on a half of the single crystal of Ag(DM)2 (long needles with typical dimensions of 1-10 mm in length and 0.01-0.1 mm thick), the current (I)-voltage (V) property curves were measured at RT across the borderline of the irradiated and the nonirradiated parts in the longitudinal direction of the needle-shape crystal, which is parallel with the conducting direction (// c axis). The result (Figure 7) indicated rectifying behavior, a characteristic property of a diode, which is the simplest device with a single junction interface. The similar measurements at lower temperatures ( ≤ 100 K) resulted in poor

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asymmetry of the I-V curves with higher resistivity, which is consistent with the fact that the pristine Ag(DM)2 becomes an insulator at T ≤ 100 K.

Figure 7. Current-Voltage property curve of the single crystal of Ag(DM)2 after UV-Vis illumination upon only half of it for ~21 days; observed (red) and theoretical fit (blue). The theoretical fit was based on the following formula;

ηeV

I = I S {exp(

kT

) − 1} ,

where I, Is, η, e, V, k and T denote observed

current (mA), saturation current (mA), reducing factor, elementary electric charge (C), applied voltage (mV), Boltzmann constant (JK-1) and temperature (K). The parameters which give the best fit are IS = 0.166, η = 0.016, when T = 290.47. Reproduced and modified from Figure 4 in ref. [12] with copyright permission.

Mechanism We will discuss here an outline of the mechanism of the above photochemical modification. Here, for simplicity, we should call the semiconducting photochemical product β, and the starting material α [18]. α produces β when it is irradiated with UV and/or Vis light of an arbitrary wavelength, intensity and atmosphere as long as the sample temperature is retained well below 155ºC during the irradiation. By varying duration of irradiation on α, the resistivity of β can be continuously controllable. It should be noted here that β indicates any solid derived from α by irradiation around RT for various duration. Thus β may be an inhomogeneous mixture and may exhibit different physical properties from another β produced under different irradiation conditions. Accordingly we had carefully checked reproducibility and consistency of the data from different samples/measurements. β has an identical appearance with α, dark blue or black fine needles (Figure 8). The measurements of TG-DTA, DSC, IR and MALDI mass spectra, elemental analysis, and X-ray powder diffraction (XRD) consistently indicated that α and β share the chemical formula of Ag(DM)2. High resolution solid state 13C-NMR of β by a cross polarization magic angle spinning (CPMAS) method indicated that the molecular electronic state of DM in β progressively changed with irradiation [14]. It is consistent with the fact that the electrical conduction of Ag(DM)2 is dependent on the DM molecules, and the way the NMR spectra changed depending on the irradiation time qualitatively agreed with the way the electrical resistivity and the Raman spectra did. Similarly the magnetic susceptibility measurements of β with different irradiation times indicated that more and more part of the sample became

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semiconducting (Curie-like behavior) progressively with irradiation (Figure 9). Between α and β what makes their electrical behavior differ from each other? The answer was given by a series of XRD patterns of β with different irradiation times (Figure 10). Within the first ~ 20 h of irradiation, the crystal structure of α gradually changed so that the diffraction peaks did not change in their positions (2θ) but did change in their relative intensities. However, beyond ~ 60 h of irradiation, a crystalline state with a different structure suddenly appeared and rapidly increased in intensity with rapid decrease of the original powder pattern. Taking the data above into consideration, we can put forward a hypothesis on the mechanism of transforming metallic α into semiconducting β by UV-Vis irradiation. It appears to be more complicated than we had anticipated. (Step 1) Immediately after irradiation, Ag and DM are excited to be nearly neutral species by transferring electrons between them. (Step 2) During relaxation from the abovementioned excited states, a part of Ag and DM species remain as lattice defects, which is consistent with the increase in resistivity and paramagnetic susceptibility, particularly evident at early stages of irradiation. (Step 3) Increase and accumulation of such lattice defects make the crystal structure of α unstable, and finally a structural transition occurs. This interpretation is corroborated by the powder XRD data. (Step 4) The newly appeared solid state exhibits semiconducting behavior. When the structural transition finishes, the irradiation effect on the resistivity is expected to saturate. This is consistent with the observed electrical behavior and the electronic states of DM determined by various spectroscopic methods [14,15,17].

Figure 8. Micrographs of α-Ag(DM)2 and β-Ag(DM)2.

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Figure 9. Temperature- and irradiation-time-dependence of paramagnetic susceptibility of α- (0 h) and β-Ag(DM)2 (12 and 40 h). Measured on the powder (polycrystalline) sample after irradiation (250-450 nm) with the Hg/Xe lamp. The jump and peak behavior around 50 K in the data of 40 h is due to impurity (residual oxygen on the surface of the sample).

Figure 10. Continued on next page.

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Figure 10. (a) Change in the powder XRD patterns of Ag(DM)2 during UV irradiation from 0-40 h. (b) Close-up view of high 2θ-angles in (a). (c) (from top to bottom) Ag plate (99.95%), 60, 50, 12 and 0 h with background [Scotch tape (broad peaks at 2θ = 10-20 deg) and base Cu plates (*)]. Black arrows in (b) indicate the diffraction peaks which changed intensity during irradiation. Black and red arrows in (c) indicate the diffraction peaks which progressively weakened (black) or emerged (red) during the irradiation. Reproduced from Figure 3 in ref. [14] with copyright permission.

Remaining Problems By examination of the electronic and structural features of Ag and DM, the change in the electrical property was connected with the changes in the chemical states of the DM molecules and the crystal structure of Ag(DM)2. Yet the details are still to be examined to obtain experimental evidence so that the mechanism can be established. Open questions include the following. 1) How is the actual structure of β? 2) What kind of relaxation stabilizes the lattice defects produced by irradiation in the early stages of the crystal structure of β until the structural transition occurs? 3) Why β is so stable? 4) How is the interface between α and β in a single crystal?

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Clearer understanding of the mechanism will contribute to enhance/improve the performance of devices fabricated by this method. An important and interesting problem lies in the mechanism of the function of devices based on the low-dimensional, strongly correlated, conducting OCTSs with marked fluctuation. Some organic superconductors exhibit an unusual electronic state at higher temperatures (RT-ca.40 K), where there is no distinction between metals and insulators because it corresponds to a supercritical liquid of a strongly correlated metal and a Mott insulator [19]. If one makes a junction structure with these kinds of materials in crystalline forms, how will the resultant devices behave? Will the junctions belong to known categories such as pn-junction and Schottky junction, or totally unknown categories? Because of the anisotropy, their performance must depend on the orientation of the single crystal relative to the direction of junction layers and the voltage applied, but how? Theoretically or experimentally, there is no prediction at present. Recently a series of intriguing experimental results have appeared one after another [20-24], which may indicate a bright prospect of the devices based on the crystalline OCTSs.

Conclusion It has become possible to modify the conducting properties of a part of an OM by controlled irradiation. The OMs are now widely recognized as important candidates for advanced electronic materials, and many applications could be realized if an effective, controllable, and versatile way of doping is available. The new method discussed here is simple, requiring only irradiation as long as the material contains a photosensitive oxidizing/reducing chemical species in addition to a multistep redox active species such as πdonor/acceptor molecules. The latter condition is usually satisfied by most conducting OCTSs, while the former can be easily met considering the many choices of the OMs. The photochemical process discussed here is obviously irreversible, and the resultant state survived at least for several months after completion of irradiation. The irradiation can be ceased or resumed at any time, and the resistivity can be continuously controlled simply by (total) duration of the irradiation. Thus, by combining this method with current advanced technology of controlling light beams for the manufacture of integrated circuitry, a new path towards manufacturing various or novel kinds of organic electronic devices will be opened.

Acknowledgment This work is carried out mainly in collaboration with Professors T. Inabe and K. Asakura, a post-doctor and students of their laboratories (Dr. H. Niimi, Mr. H. Sugawara, Mr. A. Kakizaki and Mr. T. Miyamoto) at Hokkaido University.

References [1] Facchetti, A. Materials Today 2007, 10, 28-37. [2] Hung, L. S.; Chen, C. H. Materials Science & Engineering R-Reports 2002, 39, 143-222.

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[3] Sekitani, T.; Takamiya, M.; Noguchi, Y.; Nakano, S.; Kato, Y.; Sakurai, T.; Someya, T. Nat. Mat. 2007, 6, 413-417. [4] Noguchi, Y.; Sekitani, T.; Someya, T. Appl. Phys. Lett. 2006, 89, 253507 (3 pages). [5] Someya, T.; Kato, Y.; Iba, S.; Noguchi, Y.; Sekitani, T.; Kawaguchi, H.; Sakurai, T.; IEEE Trans. Electron Devices 2005, 52, 2502-2511. [6] Aumüller, A.; Erk, P.; Klebe, G.; Hünig, S.; Schütz, J. U. V.; Werner, H.-P. Angew. Chem. Int. Ed. Engl. 1986, 25, 740-741. [7] Aumüller, A.; Hünig, S. Liebigs Ann. Chem. 1986, 1, 142-164. [8] Aumüller, A.; Hünig, S. Liebigs Ann. Chem. 1986, 1, 165-176. [9] For example, consult the world’s largest crystallographic database at Cambridge Crystallographic Data Centre (http://www.ccdc.cam.ac.uk/). [10] Reese, C.; Bao, Z. J. Mater. Chem. 2006, 16, 329-333. [11] Hasegawa, T.; Mattenberger, K.; Takeya, J.; Batlogg, B. Phys. Rev. B 2004, 69, 245115 (6 pages). [12] Naito, T.; Inabe, T.; Niimi, H.; Asakura, K. Adv. Mater. 2004, 16, 1786-1790. [13] Naito, T.; Sugawara, H.; Inabe, T.; Kitajima, Y.; Miyamoto, T.; Niimi, H.; Asakura, K. Adv. Func. Mater., 2007, 17, 1663-1670. [14] Naito, T.; Sugawara, H.; Inabe, T. Nanotechnology, 2007, 18, 424008(8 pages). [15] Naito, T.; Sugawara, H.; Inabe, T.; Miyamoto, T.; Niimi, H.; Asakura, K.; Mol. Cryst. Liq. Cryst., 2006, 455, 311-316. [16] Yamamoto, H. M.; Ito, H.; Shigeto, K.; Tsukagoshi, K.; Kato, R. J. Am. Chem. Soc., 2006, 128, 700-701. [17] Naito, T.; Sugawara, H.; Inabe, T.; Miyamoto, T.; Niimi, H.; Asakura, K. J. Non-Cryst. Solids, 2006, 352, 2628-2630. [18] Miyamoto, T.; Niimi, H.; Chun, W.-J.; Kitajima, Y.; Sugawara, H.; Inabe, T.; Naito, T.; Asakura, K. Chem. Lett., 2007, 36, 1008-1009. [19] Naito, T.; Yamada, Y.; Inabe, T.; Toda, Y. J. Phys. Soc. Jpn, 2008, 77, 064709(6 pages). [20] Sawano, F.; Terasaki, I.; Mori, H.; Mori, T.; Watanabe, M.; Ikeda, N.; Nogami, Y.; Noda, Y. Nature, 2005, 437, 522-524. [21] Tajima, N.; Fujisawa, J.; Naka, N.; Ishihara, T.; Kato, R.; Nishio, Y.; Kajita, K. J. Phys. Soc. Jpn., 2005, 74, 511-514. [22] Iimori, T.; Naito, T.; Ohta, N. Chem. Lett., 2007, 36, 536-537. [23] Iimori, T.; Naito, T.; Ohta, N. J. Am. Chem. Soc., 2007, 129, 3486-3487 [24] Iimori, T.; Ohta, N.; Naito T. Appl. Phys. Lett., 2007, 90, 262103(3 pages).

In: Nanotechnology… Editors: C.J. Dixon and O.W. Curtines, pp. 293-308

ISBN: 978-1-60692-162-3 © 2010 Nova Science Publishers, Inc.

Chapter 8

FABRICATION OF ELECTRICAL CONTACTS ON INDIVIDUAL METAL OXIDE NANOWIRES AND NOVEL DEVICE ARCHITECTURES

a

Francisco Hernandez-Ramirez1 a,b,c, Juan Daniel Prades a, Roman Jimenez-Diazb, Olga Casals b, Albert Cirera b, Albert Romano-Rodriguez b, Joan Ramon Morante a,b, Sven Barthd and Sanjay Mathur d,e

M-2E / XaRMAE, Catalonia Institute for Energy Research (IREC), Barcelona, E-08019 Spain b Departament d´Electrònica, Universitat de Barcelona, C/Martí i Franquès 1, Barcelona, E-08028, Spain c Electronic Nanosystems S.L. (e-nanos), Barcelona E-08028, Spain d Nanocrystalline Materials and Thin Film Systems, Leibniz Institute of New Materials, Saarbruecken, D-66123, Germany e Institute of Inorganic Chemistry, University of Cologne, Cologne, D-50923, Germany

Abstract Metal oxide nanowires exhibit novel properties due to their high surface-to-volume ratio and high surface stability. For this reason, they are considered excellent candidates to be incorporated into a new generation of devices with improved performance. Nevertheless, reaching complete control of their physical, chemical and electrical properties is needed before they can be widely used in our everyday life. This objective can be only fulfilled if reproducible electrical measurements on individual nanomaterials are performed. However, the fabrication of electrical nanocontacts in a fast and well-controlled process is still an unsolved issue. In this chapter, the main nanofabrication techniques that are commonly used to electrically access individual metal oxide nanowires, and to study their intrinsic properties are presented. Advantages and limitations of these methodologies are discussed in detail. By integrating bottom-up and top-down techniques, the first functional prototypes based on 1

E-mail address: [email protected]

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Francisco Hernandez-Ramirez, Juan Daniel Prades, Roman Jimenez-Diaz et al. individual nanowires have already been implemented, paving the way to the future developments of nanoscale electronics, optoelectronics and chemical sensing devices.

1. Introduction Since the 1950s, the miniaturization of the electronic components in integrated circuits (IC) has become the pillar of the huge growth and success of high-tech industries. During the last decades of the 20th century, the most popular electronic systems have penetrated into every aspect of modern life, and considerable business opportunities have arisen around them: the market of consumer electronics in the European Union is estimated to be approximately one billion euro per year [1,2]. The beginning of the 21st century has represented a new revolution of this miniaturization process and, according to some authors, this has been driven by recent advances in nanoscience and nanotechnology, which are defined by a unit of length, the nanometer (1 nm = 10-9 m) [3]. Nanoscience is a young field of research that involves several disciplines of traditional science and engineering, and it is attracting increasing public and private interest. The main goal of nanoscience is to build innovative materials and devices with new physical, chemical and electrical properties derived from the phenomena characteristic of such tiny scales [4-7]. Although the synthesis and characterization of nanomaterials has progressed in the last years, this work is still far from being finished, since there are a large number of technological issues associated with handling and manipulation of these small structures; which remains an enduring task. In first approximation, nanowires have their electrons confined in two of their spatial dimensions, while they are free to move along the third one. As a result of this geometry, new electrical transport properties with potential applications in multiple fields have been observed in these materials [7-9]. Metal oxide (MOX) nanowires are among the most promising 1D systems due to their low cost and economic relevance [10]. The majority of works devoted to the electronic applications of individual metal oxide nanowires have focused on zinc oxide (ZnO) [13-19], which is a large bandgap material (Eg ≈ 3.4 eV) with promising properties in visible-blind UV optoelectronics, such as photodetectors, photodiodes, and energy harvesting applications. Furthermore, metal oxide nanowires such as SnO2 [11,12,20-23] have also been studied, mainly motivated by their application in gas sensing [24]. In this chapter, the fabrication of devices based on individual metal oxide nanowires and their applications are briefly reviewed. The first part deals with the challenges of fabricating good and small electrical nanocontacts to individual nanowires. This issue is crucial to develop reproducible and reliable devices with superior performance. The second part is a short overview of the most significant examples of the use of MOX-nanowire systems as the key elements of a new generation of functional devices, such as UV photodetectors and gas sensors. Although most of these prototypes still do not fulfill the requirements for becoming commercial products, they are significant demonstrations that the here-presented technologies could become competitive alternatives in the near future. Finally, some important conclusions are summarized and the main research guidelines that should be followed before obtaining new and better devices are identified.

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2. Fabrication of Electrical Nanocontacts to MOX Nanowires The electrical characterization of nanowires is a maturing subject. Although the first works were based on the study of bundles of nanowires located between two microelectrodes [25,26], the lack of reproducibility of this type of measurements forced the repetition of the experiments in systems much easier to model. In other words, the scientific community realized that good control over the intrinsic properties of nanowires could only be achieved if the first experiments were performed with individual nanowires. Nevertheless, the fabrication of electrical contacts to one nanowire is not a trivial process, and even nowadays it remains one of the most challenging issues [27]. Here, a short review of the different approaches developed in the last decade to overcome the bottleneck of nanofabrication is presented. Although this book chapter focuses on MOX nanowires, most of the techniques presented here are similar to those employed with other 1D nanomaterials [1,28-31].

2.1. First Approaches The first attempts to electrically access individual 1D nanomaterials made use of very simple techniques that did not require any sophisticated equipment to develop singlenanowire devices. The nanowires were randomly deposited onto prefabricated microelectrodes using a suspension of inorganic material and organic solvents [32]. This method was not reproducible, but some of the attempts were always positioned touching at least two microelectrodes, and thus, enabling their electrical characterization in a fast way. Thanks to these first prototypes, it was possible to demonstrate the feasibility of performing basic studies on individual metal oxide nanostructures. However, the poor quality of the electrical contacts formed between the nanowire and the pre-patterned microelectrodes, and the limitation of extending this technique to large-scale processes emphasized the development of new nanofabrication alternatives. To date, this methodology has been replaced by more versatile approaches, such as e-beam and FIB lithography.

2.2. E-Beam, UV and Shadow Mask Nanolithography Techniques The need for well-controlled and reproducible electrical nanocontacts motivated researchers to develop alternatives to the simple fabrication method described in the previous section. On one hand, the most promising technology was e-beam lithography, which is considered a powerful technique for operating at the nanometer scale [1]. On the other hand, the suitability of using traditional techniques like UV- or resist free lithography as feasible alternatives to obtain measurable prototypes was demonstrated as well [33,34]. E-beam lithography is a nanofabrication technique with a lateral precision below 100 nm [1,35,36] and thus, metal depositions with well-defined shapes in the nanometer range can be obtained with commercial equipment, enabling the formation of high-quality electrical contacts with MOX nanowires by a well-established process. Nevertheless, e-beam

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lithography requires several time-consuming steps before reaching the final device and, as a consequence, it is only appropriate for prototyping purposes. The typical protocol necessary to contact a (MOX) nanowire with e-beam equipment [37] is summarized in figure 1. First, some nanowires are dispersed onto a silicon substrate with predefined alignment marks to locate the position of the nanowires. Afterwards, an electronbeam resist is spin coated onto the substrate and the e-beam lithography (sensibilization of some regions of the resist by impinging on them a controlled electron beam) is performed to define two or more electrodes touching the nanowire. Finally, a metallization followed by a liftoff process is done, resulting in a contacted nanowire ready to be measured [38].

Figure 1. e-beam lithography methodology to electrically contact a single nanowire. (a & b) The nanowire is dispersed onto a silicon substrate. (c) A resist is spin coated onto the substrate. (d) e-beam irradiation is performed. (e) Metallization and lift-off process of the metal contacts (f) Final device ready to be measured.

Figure 2. Shadow mask lithography methodology to electrically contact a single nanowire. (a & b) The nanowire is dispersed onto a silicon substrate. (c) Metallization is performed using a mask to protect the nanowire (f) Final device ready to be measured.

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It is noteworthy that most of the works devoted to the electrical studies on individual MOX nanowires exploit the high degree of freedom provided by e-beam lithography to define the final position and shape of the nanoelectrodes. However, its suitability to contact (MOX) nanowires has been questioned by some authors [39], who argue that coating the nanowires with a polymeric resist may irreversibly modify any of their intrinsic properties, changing the performances of the final devices. Thus, e-beam and by extension other techniques such as UV photolithography techniques, all of them based on the use of resists, have been gradually replaced by less aggressive alternatives like shadow mask nanolithography [39-42]. In this resist-free approach, the evaporation of metal micropads is directly performed onto the nanowire. The exact position of this deposit is determined by a shadow mask (figure 2) containing the appropriate pattern. This technology, which can be neither extended to largescale processes, has allowed reaching some of the most significant results in the characterization of individual MOX nanowires in spite of being less flexible and precise than e-beam and UV lithography.

2.3. Focused Ion Beam (FIB) Nanolithography Techniques Focused Ion Beam (FIB) is a technique developed in the 1970s for the patterning and deposition of different materials with nano-scale resolution. Their usual applications are focused on circuit edit, mask repair, material characterization and fabrication of micro-nano systems [43-48]. Its basic principle is scanning the sample surface with a focused ion beam (generally Ga+ ions accelerated at tens of keV) that sputters the exposed area. On introducing a metalorganic precursor in the beam path, accelerated electrons decompose this compound, and a significant part of this material is deposited onto the sample surface [49-52]. Thus, metallic and isolating patterns are easily obtained without the need of masks. That is the main reason why FIB has emerged as a competitive alternative to e-beam and other techniques [1]. First attempts of contacting individual nanomaterials with FIB soon demonstrated that it was possible to fabricate nanodevices in a fast and simple way. Nevertheless, the use of this technique was restrained during years because of the damage and the structural modification of these nanomaterials originated by ion bombardment. Later on, technical improvements on FIB equipments allowed fabricating nanocontacts by using low ion currents in the proximity of the nanomaterials to reduce undesired modifications. Although ion exposure was not completely eliminated, these innovative works inspired subsequent studies on the use of FIB nanolithography [53-59]. In them, well-established strategies such as nanowire dispersion on organic solvents and the use of pre-patterned microelectrodes were combined with FIB metal depositions, improving the overall quality of the electrical contacts between the nanowire and microelectrodes [60-62]. The development of the so-called Dual Beam FIB systems, which comprise a conventional FIB microscope (ion beam) and a Scanning Electron Microscope (electron beam) in a single instrument [63], facilitated the possibility of performing both ion- and electron-assisted depositions without masks, renovating the interest on this technique. In this configuration, electrons can be used for imaging the sample and then, depositing conductive materials with well-defined shapes.

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Figure 3. Sketch: (a) Microelectrodes distribution onto a silicon wafer; (b) detail of the four microprobes; (c) nanowire placed near the microelectrodes using the manipulator inside of the focused ion beam, FIB; (d) electron beam FIB deposited Pt stripes; (e) ion beam FIB deposited Pt stripes. Reprinted with permission from Ref.[115] F. Hernandez-Ramirez, et al., Sensors Actuators: B 121, 3, 2007. Copyright Elsevier (2007).

Thanks to the fact that the interaction between electrons and the sample is less destructive than using ions, the structural modification and contamination introduced during the nanofabrication process is significantly reduced if first, electron-assisted depositions are performed onto and in the proximity of the nanowire to be contacted, and then, the rest of the metal strips are fabricated using ion-assisted lithography without imaging the nanowire [27,64] (figure 3). The process described here, whose suitability to contact individual nanomaterials was previously proved with carbon nanotubes (CNTs) [65], has been successfully used in different works devoted to the fabrication of individual metal oxide nanowire-based devices [38,66,67]. Although this nanolithography methodology is neither useful for large scale-processes, it is currently one of the most versatile alternatives to contact and perform fundamental studies on individual nanowires.

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2.4. Self-assembly Approaches Despite the advances in nanolithography techniques, which enable the fabrication of electrical contacts on individual nanowires, they are merely suitable for prototyping and academic purposes. To extend the use of nanowires to low cost and large scale fabrication processes, self-assembling techniques must be taken into account. In this direction, the first steps had already been made to self-align one-dimensional (MOX) nanostructures by means of dielectrophoretic techniques [68-71]. Dielectrophoresis is an attractive alternative for the positioning and alignment of nanowires thanks to its low-cost, simplicity and flexibility [68,72]. This method is based on the well-known forces that appear when dielectrically polarized materials are in a medium in which a non uniform electric field is applied [73,74]. It has been demonstrated for different nanomaterials such as single- [75-77] and multi-walled carbon nanotubes [77,78], polymeric [79], metal [80,81] and semiconductor nanowires [69,82], and of course MOX nanowires [6871]. Dielectrophoresis can be applied to the fabrication of a new generation of nanodevices and it can be easily combined with other techniques like e-beam or conventional photolithography. If the appropriate design of electrodes is used, nanowires are not only aligned but also positioned at any desired position, and thus, the time necessary to fabricate a device is significantly reduced. This advantage can be applied to the fabrication of the simplest electronic elements, such as rectifying junctions [82] and transistors [76], paving the way for the development of novel electronic devices exclusively based on nanostructured semiconductors materials. Nevertheless, before reaching this high control of self-assembly techniques, first of all the performances of hybrid designs which combine conventional components integrated in silicon and nanowire devices must be investigated in order to achieve novel microsystems with enhanced capacities [83].

3. Applications and Functional Devices The development of nanofabrication techniques has been motivated by two main aims: (1) developing research prototypes to perform basic studies devoted to demonstrate novel properties of MOX nanowires and, (2) evaluating their potential applications. Nowadays, it is already possible tackling the second one in order to achieve functional nanodevices better than their macro and micro counterparts. In this direction, many research efforts have been recently undertaken to obtain the first low-cost and portable prototypes based on nanowires with high response and accuracy. Since the sole apparition of these devices justifies all the efforts invested in this field, we shortly review some of the first nano-class functional devices appeared in the last years. The application of nanowires in a new generation of electronic devices requires a cost-effective solution to interface the contacted nanowire with a processing decision-maker unit. Here, we first review the recent advances in low-cost multipurpose platforms to monitor the electrical properties of individual nanowires followed by the description of sensing applications. Among the large family of MOX nanowires synthesized to date, ZnO and SnO2 nanowires are considered excellent candidates to demonstrate the feasibility of monitoring

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their electrical response to different external stimuli (e.g., gas sensing) [10]. In particular, the possibility of developing UV photodetectors based on ZnO nanowires is presented. The high crystalline quality of the nanowires, which are almost defect-free 1D nanocrystals, enables to reach superior optoelectronic performances when compared to their bulk or thin film counterparts. Furthermore, the use of individual SnO2 nanowires as conductometric gas sensors, whose detection principle is the change of their electrical resistance after the exposure to gases, is also reviewed.

3.1. Low-Cost Multipurpose Platforms Most of the studies with nanowires make use of expensive and large lab instruments, and thus, the effective transfer of this technology to our everyday life remains as an unfeasible goal. For this reason, there is an increasing interest in demonstrating that low-cost portable devices with integrated nanowires can be developed to operate as functional elements [84,85]. Recently, the feasibility of monitoring the electrical properties of individual nanowires with portable cost-effective and consumer-class electronics (figure 4) was demonstrated [85]. This low-cost instrument, compared to lab equipments, was able to detect and quantify the response of individual nanowires to UV light pulses and various gas species with long-term stability due to the low current injected by the platform to the nanowire, which prevented any undesired damage of the nanocontacts due to Joule self-heating [85]. Besides the electronic instrumentation, there are other important issues concerning the appropriate operating conditions of the nanowires for certain applications. For instance, it is well-known that metal oxide materials need to be heated at a specific temperature to maximize their response to a specific target [86]. Therefore, the use of a heater becomes a fundamental tool to modulate the final performance of these materials. For this reason, both bottom-up and top-down fabrication techniques have been successfully integrated in a single process; nanowires are electrically contacted to a micro-hotplate with an integrated heater [85] (figure 4). This set up allows modulating the effective temperature of the wire as function of the power dissipated at the heater in a fast and completely reproducible manner. It is noteworthy that this solution combined with a good electronic interface, which integrates the thermal control of the nanowire, is extremely useful in manifold sensing applications [85]. Other architectures are also being explored to solve one of the major issues of sensors: the lack of selectivity to interfering stimuli. Typical examples are photodetectors sensible to a wide range of light wavelengths or gas sensors responsive to parasitic species, like moisture. For this reason, brand-new studies are attempting to develop electronic systems based on arrays of different individual MOX nanowires [87]. According to this approach, their responses are monitored in parallel, and the specific sensing characteristics of each one are determined and electronically recorded following one of the strategies previously described. Later, the data are processed by pattern recognition software to determine the composition of the external stimuli [88]. Although these studies are currently ongoing, they are the most promising solution to overcome the lack of selectivity, which is characteristic of metal oxide nanowires.

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Figure 4. Low-cost electronic interface and detail of a micromembrane with an integrated microheater. Additional magnification (inset) is necessary to see the nanowire.

3.2. Photodetectors MOX nanowires are gaining a great interest as photodetectors due to their potential applications in optoelectronics [89-91]. Among all MOX nanomaterials, the response of ZnO and SnO2 nanowires has been widely studied [92-94]. The energy of the absorbed light is used to promote an electron from the valence band to the conduction band of the material (for example in a MOX semiconductor) leaving behind a hole in a process known as photogeneration of charge pairs [95-97]. Consequently, the conductance of the material increases after exposure to light and, in a first approach, the minimum detectable wavelength corresponds to the band gap energy of the material. Since this phenomenon is essentially a bulk process, the main advantage of using nanosized materials is that the entire bulk contributes to the final response because of their high surfaceto-volume ratio [98]. Besides these simple photoconductors, there are a large number of other macro and micro devices based on the photogeneration of charge pairs that are conventionally used in real applications [95-97,99]. Most of them are based on the charge separation of the photogenerated pairs (electron and hole) by a built-in electric field. In these cases the junction of different materials (metal-semiconductor, heterojunction and homojunction of semiconductors) is used to generate this internal field. Figure 5 illustrates some these devices architectures. Recently, many of these configurations have been successfully reproduced using nanowires (diodes in individual nanowires [100,101], core-shell nanowires [102,103], nanowire-wafer unions [104]), paving the way to further improvements in photodetection applications using nanowires as the basic component. Finally, it must be mentioned that similar nanostructures and devices can be also used in solar energy harvesting applications, in which the same principle of separation of photogenerated pairs is used to provide useful power (voltage and current) to an external electrical load [95,105,106]. Particularly, the integration of nanosized materials in costeffective dye-sensitized solar cells is necessary to efficiently collect the charge generated by

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light in the organic dye layer [107-109]. Another promising achievement is the integration of a fully functional solar cell in a single nanowire [100,101,110].

Figure 5. Schematic structure of different photodetectors based on semiconductors such as MOX.

3.3. Gas Nanosensors The response of MOXs to gases is mainly due to a transfer of electrical charge between gas molecules and atoms located at their surfaces. In rough approximation, oxidizing species lead to a rise of the resistance in MOXs, whereas reducing molecules cause a drop of this parameter. Additionally, these reactions are highly influenced by oxygen vacancies inside the bulk of these materials [111] which are created during their synthesis, and other experimental parameters such environmental moisture [112] or temperature. The general theory, which explains the sensing capabilities of this type of sensors, has been reviewed in many literature reports [86,113]. Once the first individual MOX nanowires were electrically contacted using any of the techniques presented in this chapter, researchers focused on monitoring the changes of the nanowires’ resistance as a function of the surrounding atmosphere. It was soon demonstrated that the observed resistivity change in the presence of different gas species was in good agreement with the general theory for metal oxide sensors [114]. Moreover, low gas concentrations could be easily detected [115]. According to these preliminary results, the main advantage of using individual nanowirebased gas nanosensors in comparison to traditional MOX microsensors is the excellent recovery of their initial resistance baseline when the target gas is removed from the chamber, due to the fact that nanowires are single crystalline materials with well-defined and stable surfaces, in which adsorption and desorption of molecules may occur in a fast and reversible way. On the contrary, MOX microsensors are composed of a thick layer of nanoparticles agglomerated onto a substrate, whereby uncontrolled gas diffusion occurs through them, and long-term drift processes are experimentally observed [116]. It is noteworthy that the recovery of the initial resistance baseline of an individual nanowire-based gas sensor can be

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enhanced, illuminating the system with UV light pulses in order to clean their surface thanks to a gas photodesorption induced process [117]. The main disadvantage of MOX nanowires in performing gas measurements is the lack of selectivity to different gas species. This issue is common to other MOX-based sensors, and it is worsened because of the extremely high sensitivity of MOXs to environmental moisture, which dramatically changes both the nanowire resistance baseline and the cross-sensitivity to other gases, such as CO [38]. For this reason, many research efforts are being devoted to developing new devices able to overcome this intrinsic limitation [87,88]. Some of these challenging difficulties are currently being tackled with the use of doped [118] or functionalized [119] nanowires.

4. Conclusion Determining and modeling the novel properties of individual MOX nanowires requires working with only a single nanomaterial in order to reduce the parameters that have an influence on their experimental responses. For this reason, particular attention has been paid to the techniques necessary to fabricate electrical contacts with a precision of only a few nanometers. Traditional processes in the field of microelectronics, such as e-beam, UV- and shadow mask lithography have demonstrated a high rate of success in developing prototypes based on individual MOX nanowires. Nevertheless, all of these techniques have intrinsic disadvantages which in some cases even restrain their use. FIB nanolithography has recently emerged as a competitive alternative to electrically contact individual nanomaterials in a fast and reproducible way, becoming one of the best solutions for the development of research prototypes. Although the mastering of all of these technologies is crucial to obtain future breakthroughs in nanotechnology studies, they are not suitable for large-scale processes; thus, novel self-alignment techniques are currently being evaluated. On the other hand, it has been demonstrated that MOX nanowires are excellent candidates for integration into functional devices, like UV photodetectors and gas sensors. Although significant progress in this field has been achieved during the last years, the transfer of this technology to real devices is still scarce due to the high costs and other technological issues. All of the methodologies presented here to interface with nanomaterials should pave the way for obtaining new and enhanced technological applications in the near future.

Acknowledgments We cannot conclude this chapter without thanking the outstanding researchers who have actively participated in this piece of work. A considerable part of this scientific research consisted of collaborations with them. For this reason, my thanks go to all of these people we have worked with, such as our present and past colleagues at the University of Barcelona and the friends of the Leibniz Institute of New Materials. We would also like to thank Dr. C. de Mairena for the critical comments to the earlier versions of the manuscript.

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In: Nanotechnology… Editors: C.J. Dixon and O.W. Curtines, pp. 309-330

ISBN: 978-1-60692-162-3 © 2010 Nova Science Publishers, Inc.

Chapter 9

FUNCTIONALIZATION OF NANOPARTICLES, NANOTUBES AND NANOWIRES BY SURFACEINITIATED ATOM TRANSFER RADICAL POLYMERIZATION Jinying Yuan*, Mi Zhou and Yingwu Yin Key Laboratory of Organic Optoelectronics & Molecular Engineering of Ministry of Education, Department of Chemistry, Tsinghua University, Beijing 100084, People’s Republic of China

Abstract The inorganic-polymer hybrid nanomaterials have many excellent properties. So they are becoming increasingly important for various applications ranging from biomaterials to semiconductors in many fields and arouse much interest of scientists all over the world. This chapter highlights the development of surface-initiated living radical polymerizations from the inorganic materials, including nanoparticles and one-dimensional (1D) nanostructures, by surface-initiated atom transfer radical polymerization (SI-ATRP). The emphasis is put upon the new developments of SI-ATRP taken to prepare hybrid nanomaterials in the recent years.

Keywords: inorganic nanomaterials; surface initiated polymerization; atom transfer radical polymerization (ATRP); surface-initiated atom transfer radical polymerization (SIATRP)

Organic/inorganic composite materials have both the strongpoints of organic polymers and inorganic materials, and could be expansively applied in mechanism, optics, electronics, separation, catalysis, and biology in the future, and now have become a new hotspot in material science. In recent years, considerable interest exists in inorganic-organic hybrid compounds due to their excellent properties combining inorganic and organic materials. *

E-mail address: [email protected]

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Among them, functionalization or modification of gold, silicon dioxide, titanium oxide, ferric oxide nanoparticles and carbon, silicon dioxide nanotubes, zinc oxide nanowires, lanthanum hydroxide nanowires etc., has attracted special attention for their applications in medicine or electronics. It is necessary to find a convenient method for modifying inorganic nanomaterials. Up to now, synthetic polymers or biomacromolecules have been grafted or assembled onto the surface of the nanostructures via covalent bonding, electrovalent bonds or chemisorption. However, it is difficult to control the thickness of the functionalization layer. To combat this, the “living” polymerization technologies, such as atom transfer radical polymerization (ATRP)[1-3] can meet the challenge. Recently, living radical polymerization techniques were applied to surface-initiating graft polymerization. Many researchers successfully prepared a core-shell hybrid nanostructure though the covalent bonds of carboncarbon bond or silicon-carbon bond. Using this new method of functionalization on the surface of inorganic nanomaterials, the solubility in organic solvent has increased, and it affords a new application of the inorganic nanomaterials in the multiplex materials. This offers potential applications in optoelectronic, nanoscale devices, materials science and biotechnology. In this chapter, the progress of functionalization of nanoparticles, nanotubes and nanowires by surface-initiated atom transfer radical polymerization (SI-ATRP) was reviewed.

1. Functionalization of Nanoparticles by ATRP Considerable attention has been focused on the surface functionalization of inorganic nanoparticles by a polymeric shell with well-defined architecture due to the improvement of the properties of the nanoparticles, such as the dispersion and stability in various solvent. More important, the nanoparticles functionalized by polymer could combine the advantages of inorganic nanoparticles and polymer materials. Progress in polymerization has made it possible to produce polymer chains or brushes on a surface with controlled length and structure. Polymers of various architectures (block, comb, graft, hyperbranched, star, etc.) have been synthesized by living radical polymerizations. Successful examples of living radical polymerization include nitroxide-mediated radical polymerization, ATRP, reversible addition-fragmentation chain transfer polymerization, and so on. ATRP does not require stringent experimental conditions, as in the case of cationic and anionic polymerization. Growing polymer brushes on the surface of substrates by ATRP is a mature technology and a hot point. We consider it is mature, because ATRP was discovered by Matyjazewski in 1995, and the conditions of ATRP, including initiators, catalyst ect., have already been studied systemically. Another key is the chemistry of surface of substrates had been researched deeply and how to connect the polymer chains and the substrates is clearly. It is a hot field, because ATRP is a controlled and “living” radical polymerization. It has at least three strongpoints: (1) many kinds of monomers can be polymerized by ATRP, including acrylate, styrene etc. This affords the possibilities of the functions of the polymers. (2) the polymers that synthesized by ATRP have narrow polydispersity. This is an aim of almost all the scientists of polymer field. (3) ATRP is a “living” polymerization, so block copolymers are synthesized easily by ATRP. This affords the possibilities that many functions are presented at one substrate.

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Among the numerous inorganic–organic hybrid materials, silica–polymer hybrid materials are the most commonly reported in the literature. This may be attributed to their wide use and the ease of particle synthesis. Silica nanoparticles have been used as fillers in the manufacture of paints, rubber products, and plastic binders. Silica particles coated with organic modifiers are used in applications that include stationary chromatography phases, heterogeneous supported catalysts, and in the automotive, electronics, appliance, consumer goods, aerospace, and sensor industries. In general, two approaches, “grafting from” and “grafting to” can be utilized to prepare silica or other inorganic nanoparticles covalently protected with polymers. The “grafting-from” approach allows polymers to be built up at the surface of the silica naonoparticles using surface-initiated polymerization. End-functionalized monolayers, regarded as macroinitiators, may be used to initiate living/controlled polymerization, such as ATRP, directly on the surface of silica. Hult et al[4] synthesized organic-inorganic hybrid materials consisting of nanosized silica particles with surface grafted PS or PS-b-PMMA using ATRP. The surface of the nanosized silica particles was first functionalized by the reaction of silanol groups with 3-Aminopropyltrimethoxysilane (APTMS) (Fig. 1). These hybrid materials were used in the fabrication of highly-ordered isoporous membranes. Optical characterization revealed that the membranes consisted of hexagonally ordered pores of uniform size. The combination of an open pore structure and high surface area makes isoporous membranes into materials of high interest in fields as biotechnology and photonics. OH

O

OMe

OH

MeO

Si

NH2

THF

O

OMe

OH Br

Br O

Si

NH2

O

O O O

Si

N H

Br O

Figure 1. Surface Functionalization of Silica Particles[4]. Copyright Copyright Wiley-VCH (2005).

Fukuda et al[5] polymerized an oxetane group-carrying methacrylate, 3-ethyl-3(methacryloyloxy)methyloxetane (EMO), via copper-mediated ATRP initiated from the surface of monodisperse silica particles (SiPs). The polymerization proceeded in a living manner producing SiPs grafted with well-defined poly(EMO) (PEMO) of target molecular weight up to about 400K with a graft density as high as 0.36 chains/nm2. The surface-initiated ATRP of methyl methacrylate (MMA) with PEMO-grafted SiPs as macroinitiator afforded SiPs grafted with block copolymer of the type PEMO-b-PMMA ((PEMO-b-PMMA)-SiPs). The PEMO layer of (PEMO-b-PMMA)-SiPs, located between the PMMA shell and the SiP core, was cross-linked by cationic ring-opening reaction of the oxetane groups of the EMO moieties. The removal of the SiP core of the cross-linked (PEMO-b-PMMA)-SiPs by HF etching gave polymeric hollow spheres having sizeuniformity and good dispersibility in organic solvents.

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SiO2-g-PDMAEMA was synthesized successfully by our group[6], with APTMS and 2Bromo-propionyl bromide linked in sequence, and ATRP of (2-Dimethylamino)ethyl methacrylate) (DMAEMA) initiated at surface to get the organic/inorganic hybrid nanoparticle SiO2-g-PDMAEMA (Fig.2). Core/shell structure was obtained, and the hybrid nanoparticle showed better dispersion in THF than primary silica nanoparticle. Under the same temperature, the average hydrodynamic diameter of the PDMAEMA-grafted particles in acidic environment was much larger than that in basic environment. As temperature increased from 20 oC to 50 oC, Rh became obviously larger, indicating LCST between 20 oC and 50 oC, which led to the aggregation behavior of particles. These results indicate that the particles with PDMAEMA brushes are double-responsive, which provides potential applications in biomedicine and biotechnology. O H 2N

Si O

O

O

O Si

SiO2

NH2

O Br O O Br

O Si O

O O

N

O N H

O

O O Si O

Br

Br NH

n

O

O

N

Figure 2. Synthesis of SiO2-g-PDMAEMA nanoparticles[6]. Copyright Elsevier (2008).

Fe3O4 magnetic nanoparticles have been used in various fields such as sealing, oscillation damping, information storage and electronic devices. One of the rapidly developing applications of Fe3O4 magnetic nanoparticles in recent years is in biomedical areas, including rapid biologic separation and drug delivery. Surface functionality of Fe3O4 magnetic nanoparticles with polymeric shell is of current research interest owing to the flexibility in the controls of the chemical structures, compositions, and function of the polymers. Magnetite nanoparticles coated with polymers in the aqueous phase have the advantages of high stability, easy engineering of the surface properties and functional groups on the nanoparticle surface and have gained an increasing interest in a lot of fields such as hyperthermia, magnetic resonance imaging, DNA seperation, enzyme purification, and targeted drug delivery.

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Fukuda et al[7] reported the synthesis of magnetite nanoparticles coated with a welldefined graft polymer. The magnetite nanoparticles with an initiator group for ATRP, 2-(4chlorosulfonylphenyl) ethyltrichlorosilane (CTCS) chemically bound on their surfaces were prepared by the self-assembled monolayer-deposition method. The surface-initiated ATRP of methyl methacrylate (MMA) was carried out with the CTCS-coated magnetite nanoparticles in the presence of free (sacrificing) initiator, p-toluenesulfonyl chloride (Fig.3). Polymerization proceeded in a living fashion, exhibiting first-order kinetics of monomer consumption and a proportional relationship between molecular weight of the graft polymer and monomer conversion, thus providing well-defined, low polydispersity graft polymers with an approximate graft density of 0.7 chains/nm2. The molecular weight and polydispersity of the graft polymer were nearly equal to those of the free polymer produced in the solution, meaning that the free polymer is a good measure of the characteristics of the graft polymer. The graft polymer possessed exceptionally high stability and remarkably improved dispersibility of the magnetite nanoparticles in organic solvent. PMMA Immobilization of CTCS

ATRP of MMA

Fe3O4 O O Si O

SO2Cl

Immobilized-CTCS

Figure 3. Surface Functionalization of Magnetite Nanoparticles [7]. Copyright Elsevier (2004).

Schmidt et al[8] reported the preparation of Fe3O4/poly(ε-caprolactone) (Fe3O4/PCL) core-shell particles. Schmidt et al[9] also described the synthesis of nanoscopic magnetitecored polymeric brushes by surface-initiated ATRP and the properties of novel thermoreversible magnetic fluids based thereon. They presented the first results on the synthesis and characterization of novel thermoreversible magnetic fluids based on magnetite (Fe3O4) coated with a covalently anchored, polymeric shell of poly(2-methoxyethyl methacrylate) (PMEMA). The presented combination of thermoresponsive polymers with the properties of magnetic fluids, together with tailorable hydrodynamic diameter and critical temperature, contributes to the development of easily recoverable polymer-supported magnetic separation kits and catalytic systems. Li et al[10] reported the synthesis of magnetic magnetite nanoparticles coated with amphiphilic block copolymers of poly(ethylmethacrylate)-block-poly(2-hydroxyethyl methacrylate) for use as new potential carriers for hydrophobic drug delivery (Fig.4). The results show that a new core-shell-corona structural material is obtained with a very narrow molecular weight distribution of the hydrophobic segment. UV-Vis results show that 37% of progesterone is released from the nanoparticles after 22 h, much slower than free release (99% after 14 h), which demonstrates that the presence of the hydrophobic segment can effectively control the release of hydrophobic drugs.

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HO

O

Br

ATRP of EMA

HO

ATRP of HEMA

m

O

O

Fe3O4

O HO

O

Br n

m

O O

O

O

O

HO

Figure 4. Synthesis of an amphiphilic block polymer on magnetite nanoparticles [10]. Copyright WileyVCH (2006).

Our group studied a facile synthesis of magnetite nanoparticles coated with homopolymers DMAEMA as noble potential carriers for targeted drug delivery. Different from Li’s work, there is no block copolymers linked to magnetite nanoparticles so that the synthesis steps are fewer. Also, the preparation of special initiators is not required. These Fe3O4/PDMAEMA nanoparticles with core-shell structure are able to load drugs into the shell, and the release rate of drugs is approximately steady-going and can be effectively controlled by altering the pH value. Furthermore, the hybrid nanoparticles are stable to dilution because of the inherent strong interaction between the interface of the magnetite nanoparticles and the carboxyl groups of the initiators, and still have superparamagnetism which is important for drug delivery. Metal nanoparticles have attracted continuous interest owing to their unusual properties and potential uses in electronics, optics, magnetics, catalysts, and sensors. As a well-known noble metal, gold is widely investigated due to its specific impact in the fields of biotechnology and bioscience. Choi et al[11] investigated the formation of thermoresponsive gold nanoparticle/poly(Nisopropylacrylamide) (AuNP/PNIPAAm) core/shell hybrid structures by surface-initiated ATRP in aqueous media and the effect of cross-linking on the thermoresponsiveness of the AuNP/PNIPAAm hybrids (Fig.5). The disulfide containing an ATRP initiator was attached onto AuNPs and the monomer, N-isopropylacrylamide (NIPAAm), was polymerized from the surface of AuNPs in the absence or presence of a cross-linker, ethylene diacrylate, in aqueous media at room temperature. They expect that these hybrids could be useful as a stimuliresponsive optical device, such as surface plasmon resonance-based sensing materials, because of the combination of optical properties of AuNPs and control over the interparticle distance using thermoresponsiveness of PNIPAAm.

Initiator

Initiator

ATRP of NIPAAm

S(CH2)11OOCC(CH3)2Br

2

Figure 5. Surface Functionalization of gold nanoparticles[11]. Copyright Wiley-VCH (2005).

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Li et al[12] fabricated novel nanocomposites of gold nanoparticle and poly(4vinylpyridine) (Au@PVP) through surface-initiated atom-transfer radical polymerization (SIATRP) at ambient conditions. The citratestabilized gold nanoparticles were first modified by a disulfide initiator for ATRP initiation, and the following polymerization of 4-vinylpyridine (4VP) occurred on the surface of gold particles. The assembled Au@PVP nanocomposites are pH-responsive because of the pyridyl groups, which are facially protonated and positively charged. At low pH (<3.2), the polymer chains attached on gold nanoparticles are expanded by electrostatic forces, and the polymer layer is loosely swelled; hence, the Au@PVP composite particle displays a comatulid-like nanostructure in 3-D AFM images. However, at a relatively high pH (>3.2), the polymer chains shrink and wrap around the gold particle surface, which results in the aggregation of gold nanoparticles with a thin shrunken polymer layer under TEM observation. Such assembled Au@PVP nanocomposites as a smart supporter can entrap transition metal ions by their efficient coordinating segments, and subsequently, the metal ions can be reduced in situ to construct novel bimetallic nanocomposites, which are regarded as intelligent catalysts with environmental stimuli activity.

2. Functionalization of One-Dimensional Nanostructures by ATRP Ever since the discovery of carbon nanotubes (CNTs) by Iijima[13], there has been great interest in the synthesis and characterization of other one-dimensional (1D) nanostructures[14-29]. Nanowires, nanotubes, nanorods nanobelts and 1D array constitute an important class of 1D nanostructures, which provide many new and promising fields including nanofabrication [23,30-32], nanodevices [33-38], nanobiology [39], nanocatalysis [40], etc. Six families of inorganic 1D nanostructures which are more than one hundred have been synthesized up to now. The current list is as follows[14]: (1) carbon nanotubes; (2) transition metal oxide and chalcogenide 1D nanostructures; (3) transition metal halogenous 1D nanostructures; (4) mixed-phase and metal-doped 1D nanostructures; (5) boron- and siliconbased 1D nanostructures; and (6) metal 1D nanostructures. But the surface energy of the 1D nanostructures is too high to disperse the nanostructures in organic solvents, and make them difficult to explore and understand the chemistry of the nanostructures at the molecule level. For conquering this problem, two methods are used by the scientists. The first string is developing the method for growing inorganic 1D nanostructures; the second string is debasing the surface energy of the nanostructures by modifying them. Carbon nanotube is the first kind of 1D nanostructures which are discovered, and the properties of CNTs are studied deeper than other 1D nanostructures[41-45]. So the problem how to disperse the CNTs in organic solvent was studied at first. In 1998, the paper from Hadden’s group that was published in science[46] reported a method that the long-chain molecule octadecylamine was added to the open ends of shortened SWNTs via formation of the amide functionality, and the modified CNTs could be disperse in CS2 easily. This is an exciting work that is considered as a milestone of the field of modifying 1D nanostructures,

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because they had “demonstrated a methodology for preparing solutions of naked carbon metals and semiconductors in organic solutions”, and explored the chemistry of single-walled CNTs primarily. Almost at the same time, Professor Tang’ group reported other method to prepare soluble poly(phenylacetylene)-wrapped CNTs by “in situ” polymerization[47]. This is the first paper that was reported how to make the soluble polymer-CNTs complex consciously, and greatly influences the field of making polymer-CNTs complex. But these two papers still leave challenges. For Hadden’s work, although the CNTs can be diepersed in organic solvent, the solubility is low. And for Tang’s work, the solubility is increased, but it is still a “black box” that nobody knows why the polymer can wrap CNTs, and how strong the force between the chain of the polymer and the nanotubes is. How to open the “black box”? This problem has puzzled the scientists until 2003. From the end of 2003 to the beginning of 2004, at least three groups reported the same model of reaction almost at the same time: to make the polymer-CNTs complex by atom transfer radical polymerization (ATRP) on the surface of the CNTs[49-51].

Figure 6. Grafting PnBMA from the surface of CNTs using ATRP[49]. Copyright American Chemical Society (2004).

All of the three groups chose the same kind of initiator, alkyl bromides, and the same kind of monomer, acrylate (Fig.6-8). But the details of the functionalization were a litter different. Ford’s group and Yan’s group chose the similar method to make the initiators: first, carboxyl-contained CNTs were prepared by oxidation of the crude nanotubes; second, the initiators were prepared by two esterifiable steps. The difference between their methods was the order of the two esterifiable steps. Adronov’s group chose a different procedure to prepare the initiator involving first 1,3-dipolar cycloaddition to introduce phenol unctionalities, followed by an esterification with 2-bromoisobutyryl bromide.

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Figure 7. Synthesis of an ATRP macroinitiator using a 1,3 dipolar cycloaddition[50]. Copyright American Chemical Society (2003).

The second difference among the three groups’ works is the aim. We consider that Ford’s work leans to increase the solubility and stability of CNTs in organic solvent; Adronov’s work prefers to change the structures of polymer on the surface by easy reaction to control the solubility of nanotubes in organic solvent or water; and Yan’s work selects to synthesize block copolymer on the surface to prove the chemistry excellence of ATRP.

Figure 8. Grafting PMMA from the surface of CNTs using ATRP[51]. Copyright American Chemical Society (2004).

However, the three group’s works opened the “black box” which puzzled the scientists who researched in the polymer-nanotube field. The polymer chains are grown from the initiators on the surface, and the key became how to connect the initiator molecules onto the nanotubes. But this key has been solved already. These important works hearten the field of CNTs-polymer complex. The following scientists, Warren T. Ford from Oklahoma State University, Alex Adronov from McMaster University and Deyue Yan from Shanghai Jiao Tong University, are the precursors. There is another interesting thing that the first authors of the three paper are all from China.

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Nowadays, different structures of carbon -polymer complex have been prepared by ATRP[48]: CNTs-poly(n-butylmethacrylate) (CNTs-PnBMA)[49], CNTs-poly (methyl methacrylate) (CNTs-PMMA)[50-52], CNTs-poly(methyl methacrylate) -b-polystryene (CNTs-PMMA-b-PS)[52], CNTs-polystryene (CNTs-PS)[52-46], CNTs-polystryene-bpoly(methyl methacrylate) (CNTs-PS-b-PMAA)[52], CNTs-poly(tert-butyl acrylate) (CNTsPtBA)[50], CNTs-poly(acrylic acid) (CNTs-PAA)[50,68], CNTs- polystyrene/ poly(Nisopropylacrylamide) (CNTs-PS/ PNIPAAm)[54], CNTs-poly(methyl methacrylate)-bpoly(hydroxyethyl methacrylate) (CNTs-PMMA-b-PHEMA)[50], CNTs-polystryene- bpoly(tert-butyl acrylate) (CNTs-PS-b-PtBA)[53,55], CNTs-polystryene- b- poly(acrylic acid) (CNTs-PS-b-PAA)[53,55], CNTs-poly(N-isopropylacrylamide) (CNTs-PNIPAAm)[57,58], CNTs-poly(sodium 4-styrenesulfonate) (CNTs-PSS) [59,60], CNTs-poly(2,2diethylaminoethyl methacrylate) (CNTs-PDEAEMA) [61,62], CNTs-poly(2,2dimethylaminoethyl methacrylate) (CNTs-PDMAEMA) [63], CNTs-poly(glycerol monomethacrylate) (CNTs-PGMA) [64], CNTs-poly (3-(trimethoxysilyl) propyl methacrylate) (CNTs-TMSPMA) [65], CNTs- polystyrene-b-polyacrylonitrile (CNTs-PS-bPAN)[66], CNTs-poly(Methacryloyloxyethyl phosphorylcholine) (CNTs-PMPC)[67], CNTspoly(lactobionamidoethyl methacrylate) (CNTs-PLAMA)[67], CNTs-poly(3-O-methacryloyl1,2:5,6-di-O-isopropylidene-D-glucofuranose) (CNTs-PMIG) [68], and so on. Among them, Yan’s systemic and in-depth works are deserved to emphasize especially. They not only prepared many structures of the CNTs-polymer complexes by ATRP, but also found another field called functionalization and re-functionalization of CNTs. The work published at first was to prepare CNTs-PNIPAAm [67]. “The resulting hybrid molecular nanotubes showed temperature switching assembly and disassembly behaviors in water, because of the hydrophilic and hydrophobic transformation of the bonded PNIAAm chains at 30-35oC ”.

Figure 9. SEM and TEM images of PNIPAAm-modified ACNTs. SEM images viewed from the a) top and b) side. c) The magnified image of (b). d) TEM image of a single CNT taken from the ACNT film with PNIPAAm modification.[68] Copyright Wiley-VCH (2004)

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Jiang&Li and their co-workers successfully grafted PNIPAAm onto aligned carbon nanotubes by ATRP combined the technology of aligned carbon nanotubes and temperature responsiveness of PNIPAAm chains (Fig.9)[68]. A water droplet exhibits different spreading behaviors on the film surface that had been modified, which shows different wettability under low and high temperatures. The good reproducibility of this effect shows a high degree of control over the responsive wettability of the ACNT film. Gao and his co-worker synthesized CNTs-PMIG[68], R.Narain’ group prepared CNTsPMPC, CNTs-PLAMA[67]. Both the two works, the resulting hybrids can be found many applications in biomedical fields, because the polymers that be used to coat CNTs are biocompatible. Zhang’s group prepared a steady core-shell nanostructures due to the hydrolysis and polycondensation of PTMSPMA[65].

Figure 10. The layer-by-layer self-assembly procedure on CNTs surface.[59] Copyright Elsevier (2005).

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Then Yan’s group prepared polyelectrolyte-CNTs by ATRP and layer-by layer selfassembly (Fig.10)[59]. It is an interesting work. The polyanion coated carbon nanotubes was prepared by ATRP and used as the substrates for layer-by-layer (LBL) deposition of PDMAEMA or hyperbranched poly(sulfone amine) (HPSA) and then deposition of poly(sodium 4-styrenesulfonate) (PSS). The cycle was repeated and the core-shell heterostructures had been accomplished to produce. Almost at the same time, Ford and his coworkers did the similar work[60]. Another important function that had been exploited is some kinds of polymer that contain a large number of hydroxyl groups or carboxyl groups coated carbon nanotubes could sequester metal ions and produce metal nanoparticles. Gao and his co-worker prepared CNTsPGMA[64]. And the hydroxyl groups of the polyGMA chains grafted onto the MWNTs are still highly active and can be further reacted with succinic anhydride to be converted into carboxylic acid groups. The complex was successfully loaded with Ag+, Co2+,Ni2+, Au3+, Y3+, and La3+ ions (Fig.11-13). In the case of Ag+ ions, the loading capacity of the complex is higher than that of pristine or oxidized CNTs by 1-2 orders of magnitude that had been reported. Farther, Ag+ ions were reduced to zerovalent Ag nanobead-like structures by the residual hydroxyl groups on the esterified polyGMA chains. These nanobeads with a diameter of 3-10 nm on the surface of the nanotubes were tightly enwrapped by the polymer chains, forming stable necklace-like structures.

Figure 11. Functionalization of CNTs with PGMA by ATRP, Esterification of the Hydroxyl Groups of CNTs-PGMA, and Metal Sequestration/Reduction by the Grafted Polyacid Chains.[64] Copyright American Chemical Society (2005).

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Figure 12. Transmission electron microscopy (TEM) images of CNTs-PGMA/metal (ion) nanocomposites: (a, b) CNTs-PGMA/Ag with a silver concentration of ca. 0.46 g per gram of CNTsPGMA, (c) CNTs-PGMA/Ag with a silver concentration of ca. 0.15 g per gram of CNTs-PGMA, (d) CNTs-PGMA/Ag hybrid nanowire, (e) CNTs-PAA/Ag nanowire section, and (f) CNTs-PGMA/La3+ nanocomposites.[64] Copyright American Chemical Society (2005).

Figure 13. High-resolution transmission electron microscopy (HRTEM) images of CNTs-PGMA/Ag (a, b).[64] Copyright American Chemical Society (2005).

Further research extended the concept of sequestering metal nanoparticles within the polymer chain. Gao and his co-worker prepared CNTs-PDMAEMA and CNTs- PDEAEMA[61-63]. And the resulting composites were treated with methyl iodide, quaternizing the amino groups and forming a cationic polyelectrolyte modifying CNTs. Due to high density of ammonium groups on the surface of nanotubes, higher amounts of CdTe or Fe3O4 can be adsorbed. The optical properties of quantum dots and the magnetic properties of the Fe3O4 nanoparticles in the tri-composes hybrids were maintain. The author also showed the sheep blood cells that

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could be attached with magnetic nanotubes in a buffer solution can be selectively rotated or conveyed in a magnetic field (Fig.14, Fig.15), because the magnetic nanotubes can be easily manipulated in the magnetic field.

Figure 14. The procedure to synthesize the magnetic CNTs hybrids.[62] Copyright American Chemical Society (2006).

Figure 15. Continued on next page.

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Figure 15. Representative TEM images of the magnetic nanotube sample with 6.3 wt % of iron (a, b) and the sample with 22.5 wt % of iron (c-e). HRTEM image of an iron oxide nanoparticle attached on tube surface, showing the lattice of iron oxide crystal (f). STEM images of the magnetic nanotube samples with 6.3 (g) and 22.5 (h) wt % of iron. The scale bars in a, b, c, d, e, and f correspond to 100, 20, 100, 50, 20, and 2 nm, respectively.[62] Copyright American Chemical Society (2006).

Figure 16. Electrophilic addition on SWNTs to produce an ATRP macroinitiator.[69] Copyright Springer(2005).

Another developmental direction of polymer-CNTs complex by ATRP is to prepare new initiators on the surface. Many methods have been reported. Direct electrophilic addition of chloroform to SWNT sidewalls, followed by hydrolysis in methanolic KOH to produce hydroxylated SWNTs was reported by Paik’s group[69]. The ATRP initiators which could be used for polymerization of stryene were prepared by esterification of the hydroxylated SWNTs with 2-chloropropionyl chloride (Fig.16). Terrones and co-workers also utilized a free radical process for the surfacefunctionalization of nanotubes[70]. This was accomplished by thermolysis of benzoyl peroxide at 105oC to produce phenyl radicals, which efficiently coupled to the surface of Ndoped MWNTs (Fig.17). Subsequently, the covalently grafted aromatic rings on the surface of the tubes, were brominated by Br2/FeBr3. The doped tubes did not require any acid treatment to protect the structure of CNTs farthest.

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Figure 17. The functionalization process of CNTs by radical reaction.[70] Copyright Elsevier (2006).

Because the research of the chemistry of CNTs and the experience that ATRP growing chain ends have shown reactivity toward C60 via radical addition reactions. Liu and his coworker found that the ATRP active species could transfer to CNT surfaces (Fig.18)[64], after an addition reaction. With the interesting result, linear PS and V-shaped PS-b-PNIPAAm polymer chains were successfully incorporated on CNT through the “grafting-from” and sequential “grafting-on” and “grafting from” method, respectively. With the functionalization, CNTs-PS/PNIPAAm showed amphiphilic behavior.

Figure 18. Incorporation of Amphiphilic V-Shaped Poly(styrene)-b-poly(N-isopropylacrylamide) (PSPNIPAAm). Polymers onto Carbon Nanotube Surface through Sequential “Grafting-To” and “GraftingFrom” Techniques.[64] Copyright American Chemical Society (2007).

Hong and Pan reported a method to grow hyperbranched macromolecules on the surface of CNTs by a self-condensing vinyl polymerization (SCVP) strategy via ATRP[71]. It is a good attempt for development more kinds of polymer structure, and it is good for research the chemistry of the CNTs.

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Figure 19. Grafting PMMA from the surface of Si/SiO2 core/shell nanowires using ATRP.[72] Copyright American Chemical Society (2005).

Figure 20. (a) SEM image of the pyrolysis product from a surface-initiated polymerization of PMMA on ZnO. (b) TEM micrograph of the sample in a. (c) TEM micrograph of the pyrolysis product from a sample of dipcoated PAN on ZnO.[73] Copyright American Chemical Society (2006).

Just like the discovery of 1D nanostructures, the scientists dissatisfy to modify carbon nanotubes only. Because of the particularity of carbon element, the methods used to carbon nanotubes are not suitable to other 1D nanostructures. But the key idea for polymer-1D nanostructures, how to prepare the initiators on the surface, was inherited. According to the

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experience to prepare polymer-inorganic nanoparticles hybrids, Arnold and his co-workers from University of California[72,73], Berkeley and our group from Tsinghua University[74] reported the methods to synthesize functional polymer from inorganic 1D nanostructures successively. Arnold and his co-workers took the lead in demonstrating a simple method to produce coaxial poly(methyl methacrylate) silica / silicon oxide/ poly(methyl methacrylate) core-shell nanostructures, then the Si cores were etched using XeF2 giving polymer nanotubes supported on silicon oxide (Fig.19)[72]. The polymer layer makes the outer surface of these tubes hydrophobic, while the SiO2 on the inner surface is hydrophilic. Later, they reported ATRP on ZnO nanowires[73]. By controlling catalyst and reaction conditions, polymerizations of methyl methacrylate, phenyl methacrylate, ethylene glycol dimethacrylate, and styrene on ZnO were all successfully performed (Fig.20). And these hybrid organic-inorganic materials were used as precursors for graphitic carbon tubes. MeO MeO

Si

NH2

O

MeO

OH

*

n

O

Si

NH2 n

toluene,reflux,36h

O

La(OH)3 O O O

Br

Br Br

*

O

Si

Et3N,CH2Cl2,24h

N H

n

O O O

H2 C

Styrene,CuBr,PMDETA 100oC, 24h

*

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Si

N H

H C

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Figure 21. Grafting PS from the surface of lanthanum hydroxide nanowires using ATRP. [74] Copyright IOP(2007).

As a consequence of their unique electronic structures and the numerous transition modes involving the 4f shell of their ions, lanthanide compounds usually have outstanding optical, electrical, and magnetic properties, and have been widely used as high-quality phosphors[75], up-conversion materials[76], catalysts[77], and time-resolved fluorescence (TRF) labels for

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biological detection[78]. Our group demonstrated a method for preparation of polystyrene shells based on surface of lanthanum hydroxide nanowires (Fig.21, Fig.22)[74].

Figure 22. (a) TEM image of La(OH)3 nanowires (bar = 200 nm), (b) Electron diffraction pattern of a single La(OH)3 nanowire, (c) HRTEM image of a single La(OH)3 nanowire, (d) TEM image of Sample 1 (bar = 20 nm), (e) HRTEM image of Sample 1 (bar = 10 nm), (f) HRTEM image of Sample 2 (bar = 10nm), (g) HRTEM image of Sample 3 (bar = 10 nm).[74] Copyright IOP(2007).

The key step of the synthetic process is the initiator of the polymerization for the process, using the hydrolyzation of the silane that the amino substituted, the amidogen was inducted on the surface of the nanowires; using 2-bromo-propionyl bromide or 2-bromo-2-

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methylpropionyl bromide to react with the amidogen on the surface. The only difference is the order. Recently, Zhao’s group reported hyperbranched polymers were grafted onto halloysite nanotubes by SCVP via ATRP[79]. The initiator was prepared by adsorption of 2bromoisobutyric acid on the surface of the nanotubes, and then, polymers were synthesized with the same method that Hong and Pan had used[71]. In conclusion, inorganic nanomaterials have caused for widespread concern. However, they could not be dispersed well in water or organic solvents, due to their particularly surface areas. Thus the application of inorganic nanomaterials is limited. Meanwhile, polymers have excellent properties that inorganic compounds are difficult to replace, such as processability and biocompatibility. The inorganic-organic hybrid nanomaterials emerge. The controlled polymerization, especially the surface-initiated atom transfer radical polymerization, was used to improve the inorganic nanomaterials. Using this new avenue of functionalization on the surface of inorganic nanomaterials, the solubility in organic solvent increased affording new applications in the multiplex materials. A variety of new types of nanomaterials, combined the properties of inorganic compounds and polymer, were prepared. These new inorganic-organic hybrid nanomaterials are expected to use as sensors, biological probes, and so on.

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Jinying Yuan, Mi Zhou and Yingwu Yin H. Kong, C. Gao, D. Y. Yan, J. Am. Chem. Soc. 2004, 126, 412-413. D. Baskaran, J. W. Mays, M. S. Bratcher, Angew. Chem. Int. Ed. 2004, 43, 2138-2142. H. Kong, C. Gao, D. Y. Yan, Macromolecules 2004, 37, 4022-4030. Y. L. Liu, W. H. Chen, Macromolecules 2007, 40, 8881-8886 H. Kong, C. Gao, D. Y. Yan, J. Mater. Chem. 2004, 14, 1401-1405. S. H. Qin, D. Q. Qin, W. T. Ford, D. E. Resasco, J. E. Herrera, Macromolecules 2004, 37, 752-757. H. Kong, W. W. Li, C. Gao, D. Y. Yan, Y. Z. Jin, D. R. M. Walton, H. W. Kroto, Macromolecules 2004, 37, 6683-6686. T. L. Sun, H. Liu, W. L. Song, X. Wang, L. Jiang, L. Li, D. B. Zhu, Angew. Chem. Int. Ed. 2004, 43, 4663-4666. H. Kong, P. Luo, C. Gao, D. Y. Yan, Polymer 2005, 46, 2472-2485. S. H. Qin, D. Q. Qin, W. T. Ford, Y. J. Zhang, N. A. Kotov, Chem. Mater. 2005, 17, 2131-2135. W. W. Li, H. Kong, C. Gao, D. Y. Yan, Chinese Sci. Bull. 2005, 50, 2276-2280. C. Gao, W. W. Li, H. Morimoto, Y. Nagaoka, T. Maekawa, J. Phys. Chem. B 2006, 110, 7213-7220. W. W. Li, C. Gao, H. F. Qian, J. C. Ren, D. Y. Yan, J. Mater. Chem. 2006, 16, 18521859. C. Gao, C. D. Wo, Y. Z. Jin, W. W. Li, S. P. Armes, Macromolecules 2005, 38, 86348648. Q. Xiao, S. J. He, L. W. Liu, X. Z. Guo, K. Y. Shi, Z. J. Du, B. L. Zhang, Compos. Sci. Technol. 2008, 68, 321-328 A. M. Shanmugharaj, J. H. Bae, R. R. Nayak, S. H. Ryu, J. Polym. Sci. Polym. Chem. 2007, 45, 460-470. R. Narain, A. Housni, L. Lane, J. Polym. Sci. Polym. Chem. 2006, 44, 6558-6568. C. Gao, S. Muthukrishnan, W. W. Li, J. Y. Yuan, Y. Y. Xu, A. H. E. Müller, Macromolecules 2007, 40, 1803-1815. J. H. Choi, S. B. Oh, J. Chang, I. Kim, C. Ha, B. G. Kim, J. H. Han, S. Joo, G. Kim, H. Paik, Polym. Bull. 2005, 55, 173–179. B. Fragneaud, K. Maseneli-Varlot, A. Gonzalez-Monitiel, M. Terrones, J. Cavaillé, Chem. Phys. Lett. 2006, 419, 567-573. C. Y. Hong, Y. Z. You, D. C. Wu, Y. Liu, C. Y. Pan, Macromolecules, 2005, 38, 26062611. M. J. Mulvihill, B. L. Rupert, R. R. He, A. Hochbaum, J. Arnold, P. D. Yang, J. Am. Chem. Soc. 2005, 127, 16040-16041. B. L. Rupert, M. J. Mulvihill, J. Arnold, Chem.Mater. 2006, 18, 5045-5051. M. Zhou, J. Y. Yuan, W. Z. Yuan, Y. W. Yin, X. Y. Hong, Nanotocnology 2007, 18, 405704. J. P. Cotter, J. C. Fitzmaurice, I. P. Parkin, J. Mater. Chem. 1994, 4, 1603-1609. J. A. Capobianco, F. Vetrone, J. C. Boyer, A. Speghini, M. Bettinelli, Opt. Mater. 2002, 19, 259-268. M. S. Palmer, M. Neurock, M. M. Olken, J. Am. Chem. Soc. 2002, 124, 8452-8461. A. H. Peruski, L. H. Johnson, L. F. Peroski, J. Immunol. Methods 2002, 263, 35-41. B. Mu, M. F. Zhao, P. Liu, J. Nanopart. Res. 2008, 10, 831-838.

In: Nanotechnology… Editors: C.J. Dixon and O.W. Curtines, pp. 331-370

ISBN: 978-1-60692-162-3 © 2010 Nova Science Publishers, Inc.

Chapter 10

SYNTHESIS AND APPLICATIONS OF NANO-SIZED FERROELECTRICS VIA MECHANOCHEMICAL ACTIVATION L.B. Konga, Z. Xub and T.S. Zhangc a

Temasek Laboratories, National University of Singapore, 10 Kent Ridge Crescent, Singapore 119260 b Electronic Materials Research Laboratory, Xi’an Jiaotong University, Xi’an 430069, P. R. China c Institute of Materials Research and Engineering, 3 Research Link, Singapore 117602

Abstract Ferroelectric materials have been found to be promising candidates in applications of a wide range of electronic devices, such as high-dielectric constant capacitors, piezoelectric sonar or ultrasonic transducers, pyroelectric security sensors, medical diagnostic transducers, electrooptical light valves, and ultrasonic motors, and so on. Ferroelectric materials were conventionally fabricated via solid-state reactions at relatively and sometimes extremely high temperatures for calcining and sintering. Due to the presence of volatile components, such as lead (Pb), bismuth (Bi) or lithium (Li), in most ferroelectric compounds, high temperatures processing would brought out the problems of losing of the elements, which often resulted in the deteriorations in microstructures and thus electrical performances of the ferroelectric materials. To reduce the fabrication temperatures of ferroelectric ceramics, it is necessary to use ultrafine powders. High-energy mechanochemical technique, as an alternative method, has been used to synthesize nanosized ferroelectric powders directly from their oxide and other precursors. This chapter serves as an overview of progress in the synthesis of various ferroelectric materials by using various mechanochemical milling facilities. In addition, applications of nanosized ferroelectric powders in materials preparation and device fabrication will be also be included.

Keywords: Mechanochemical synthesis, mechanochemical activation, oxide, materials, nanosized powders, ferroelectrics

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1. Introduction Ferroelectric materials possess spontaneous polarizations that are reversible under an external applied electric field [1-3]. The distinct dielectric, piezoelectric, pyroelectric and insulating properties of ferroelectric materials have made them useful for a wide range of applications, such as sensors, actuators and transducers. As a result, ferroelectrics have been one of the most intensively studied topics in scientific researches. Properties and performances of materials are determined by both their intrinsic and extrinsic characteristics [3, 4]. Intrinsic properties include types of elements and the combinations of them, while extrinsic ones are closely related to the processing parameters. Parameter-related properties include microstructures, grain size and size distribution, density and porousity, and so on. One of the most important processing parameters is synthesis technique. It has been shown that synthesis method of ferroelectric powders has played a significant role in determining the microstructural, electrical and optical properties of the final ferroelectric ceramics. Ferroelectric powders were conventionally prepared via solid-state reaction process, starting from constituent oxides, hydroxides or carbonates. Solid-state reaction process usually requires relatively high temperature, which results in powders with large grains or particles. Due to the rough grains or particles, these powders can only be fully sintered at even higher temperatures to achieve ferroelectric ceramics with desired performances. Many ferroelectric materials contain lead (Pb) component that has a high volatility and is likely to be lost during the high temperature process, which in turn leads to worsened properties. To reduce the sintering temperature, it is necessary to use powders of with fine grain size and narrow size distribution. During the last decades, a lot of attempts have been made to synthesize submicron or even nanosized ferroelectric powders. Various wet-chemistry methods, including chemical coprecipitation, sol-gel process, hydrothermal synthesis, microemulsion, combustion, thermal pyrolysis spray, molten salt, etc. have been developed for this purpose. Besides the significant achievements of these chemical methods, they also have various unavoidable disadvantages. For example, sol-gel process uses metal alkoxides as the starting materials, which are very expensive and extremely sensitive to the environmental conditions such as moisture, light and heat. Moisture sensitivity makes it necessary to conduct the experiment in dry boxes or clean rooms. Co-precipitation processes involve repeated washing in order to eliminate the anions coming from the precursor salts used, making the process complicated and very time consuming. Furthermore, it is difficult to produce large batches by using most of the chemical solution processing routes. Therefore, exploring alternative methods for the preparation of ferroelectric ceramics is still of technological as well as scientific significances. Mechanochemical synthesis, which is also known as mechanical alloying, high-energy mechanical milling, high-energy milling, high-energy activation and many others (they are not differentiated in the present review unless otherwise stated), was initially invented to prepare oxide-dispersed metallic alloys for structural applications and subsequently applied to extensions of metallic solid solubility, synthesis of intermetallics, disordering of intermetallics, solid-state amorphization, nanostructured materials, and mechanochemical synthesis of nanosized oxides or metal powders [5, 6]. Up to now, this technique has been used to synthesize various ferroelectrics [7]. The mechanical technique is superior to both the conventional solid-state reaction and the wet-chemistry-based processing routes for several

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reasons [7, 11, 23]. First of all, it uses cost-effective and widely available oxides as the starting materials and skips the intermediate temperature calcination step, leading to a simpler process. Secondly, it takes place at room temperature in well-sealed containers, thus effectively alleviating the loss of the volatile components, such as lead, bismuth and lithium. This is of particular interest to the synthesis of ferroelectric materials, since most ferroelectric ceramics contain lead (Pb), bismuth (Bi) or lithium (Li). Furthermore, due to their nanometer scale size and very high homogeneity, the mechanochemically derived ceramic powders demonstrate much better sinterability than those synthesized by the conventional solid-state reaction and wet-chemical processes. Also, the high-energy milling can greatly improve the reactivity of precursors by reducing the phase formation temperatures of some ferroelectric materials that cannot be directly synthesized, such as BaTiO3 [76-81] and many Arivillius family ferroelectrics [86-88, 91-102]. A thorough review on the progress in synthesis of nanosized ferroelectric powders and ceramics has been published recently [7]. This chapter aims to provide a briefly overview on the progress in synthesis and moreover to introduce several applications of the nanosized ferroelectric powders produced by the mechanochemical activation techniques [120-133].

2. Experimental Procedure 2.1. High-Energy Mechanical Milling Excellent reviews on the application of high-energy ball milling process to various metallic materials can be found in the open literatures [5, 6], where detailed experimental procedures have been discussed. Here, only a brief introduction is given to each type of the milling techniques. There are various types of high-energy milling machines that have been used and reported in the literature [5]. Different types of mills have different efficiencies and capabilities. The productivity of high-energy mills can be from several grams to as much as thousands of kilograms. High-energy mills that have been widely used for research purpose are vibrational shake mills (SPEX), planetary mills and attritor mills. Other equipments, for example, multi-ring-type mill (Model MICROS: MIC-0, Nara Machinery, Tokyo, Japan), are also used in mechanochemical synthesis of some of the ferroelectric powders. Either stainless steel or tungsten carbide milling media were used in the experiments of high-energy milling.

2.1.1. Vibrational Shake Mills SPEX vibrational shake mill is one of the most widely used equipments in the research community of mechanochemical synthesis. The common variety of SPEX shaker mill has one vial, containing the sample and grinding balls, secured by a clamp and swung energetically back and forth several thousand times a minute. The back-and-forth shaking motion is combined with lateral movements at the ends of the vial, so that the vial appears to be following a figure “8” or infinity sign as it moves. With each swing of the vial, the balls impact against the sample and the end of the vial, both milling and mixing the sample. Due to the amplitude (~5 cm) and speed (~1200 rpm) of the clamp motion, the ball velocities are very high (on the order of 5 m/s) and consequently the force caused by the ball’s impact is very large, hence, it is a high-energy mill. The only shortcoming of this type of mill is its

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relatively small throughput, but this problem has been addressed by designing two-vial equipment [5]. Another type of shake mill is Fritsch Pulverizette 0 (provided by Fritsch GmbH in Germany). Both types of equipments, SPEX 8000 [9, 20, 37-42, 46-54] and Fritsch Pulverizette 0 [57], have been employed to synthesize ferroelectric nanosized powders.

2.1.2. Planetary Ball Mills Planetary ball mills have better capability than the SPEX mill. Vials are arranged on a rotating support disk and a special drive mechanism causes them to rotate around their own axes. The centrifugal force produced by the vials rotating around their own axes and that produced by the rotating support disk both act on the vial contents, which include the materials to grind and the grinding balls. Since the vials and the supporting disk rotate in opposite directions, the centrifugal forces acted alternatively in the same and opposite directions. This causes the grinding balls to run down the inside wall of the vial (the friction effect) followed by the material being ground and grinding balls lifting off and traveling freely through the inner chamber of the vial and colliding against the opposing inside wall (the impact effect). A planetary mill, combined with high-density milling media such as stainless steel and tungsten carbide, also can provide high-energy ball milling. The most popular planetary mills that are reported in the literature are supplied by Fritsch and Retsch in Germany. Various ferroelectric powders have been synthesized using planetary mills [11, 12, 24-26, 30-36, 43, 56].

2.1.3. Attritor Mills Although attritor mill has been rarely reported to be used as a tool to synthesize ferroelectric powders, it is of significance to give a brief description. A conventional planetary ball mill consists of a rotating horizontal drum half-filled with small steel balls. As the drum rotates the balls drop on the powders that are being ground; the rate of grinding increases with the speed of rotation. When speeds are too high, however, the centrifugal force acting on the flying balls exceeds the force of gravity, and the balls are pinned to the wall of the drum. As a result, the grinding action stops. An attritor consists of a vertical drum with a series of impellers inside it. Set progressively at right angles to each other, the impellers energize the ball charge, causing powder size reduction because of impact between balls, between balls and container wall, and between balls, agitator shaft, and impellers. Some size reduction appears to take place by interparticle collisions and by ball sliding. A powerful motor rotates the impellers, which in turn agitate the balls in the drum. The most important advantage of attritor mills is their capability of processing large quantities of powders (from about 0.5 to 40 kg). The operation of an attritor is different from vibration and planetary mills. Powders to be milled are placed in a stationary tank together with the grinding media. During milling process, the mixture is agitated by a shaft with arms, rotating at a speed of about 250 rpm. This causes the media to exert both shearing and impact forces on the material [5].

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2.1.4. Process Parameters Mechanical milling process involves a number of variables that can be adjusted to achieve different performances [5]. The important parameters include type of mill, materials used for the milling vial and balls, milling speed, milling time, ball-to-powder weight ratio, milling environment, process control agent (PCA) used, temperature controlling and application of electrical or magnetic field during milling. Typically, for a SPEX shaker mill, 5 g of starting mixture is milled, using a stainless-steel cylindrical vial of 40 mm in diameter and 40 mm in length, with a milling ball of 12.7 mm in diameter. The milling speed is ~900 rpm. Milling time is dependent on the formation abilities of the designed compounds. The planetary mill can mill 20 g of powder at a time. A tungsten carbide vial of 250 ml and tungsten carbide balls with various diameters are used as the milling media. Typical ball-topowder weight ratio is 10-50:1. Milling speed is 200-400 rpm in most cases. The milling temperature is not controlled intentionally. However, the milling really causes temperature to rise during the milling process. Till now, it remains impossible to monitor the temperature of high-energy milling. However, the raised temperatures should be much lower than the calcination temperatures used in the conventional solid-state reaction process. Unless otherwise specified, milling is generally carried out in air.

3. Direct Phase Formation of Ferroelectrics 3.1. Normal Ferroelectrics and Antiferroelectrics 3.1.1. Lead Titanate and Lead Lanthanum Titanate Lead titanate (PbTiO3 or PT) ceramics is a typical ferroelectric material, with a phase transition temperature (Curie temperature or TC) of ~490°C. The unique properties of PT ceramics, such as high transition temperature, low ratio for the planar-to-thickness coupling factor, low aging rate of the dielectric constant and low dielectric constant, make them useful to a variety of applications. For example, PT ceramics are very good candidates as stable pyroelectric and piezoelectric devices for high temperature or high frequency applications. It is well known that PT ceramics cannot be prepared via the conventional ceramic process without the addition of other elements. This is because the anisotropic thermal expansion caused by the phase transition from cubic paraelectric to tetragonal ferroelectric (with a relatively large c/a ratio of ~1.065) during cooling from a high sintering temperature induces large internal stresses, which destroys the ceramics by creating numerous microcracks. Modification with other dopants allows for the fabrication of dense PT ceramics using conventional ceramic process, but it tends to worsen their electrical properties. Therefore, there is very rare report on preparation of undoped, dense PT ceramics. Spontaneous cracking in polycrystalline ceramics results from internal stress among the grains, which is caused by incompatible strains from thermal expansion anisotropy during the phase transformation. The area in which microcracking is likely to initiate is defined by a certain grain size limit. The stress per unit grain boundary area is proportional to the grain volume. PT has been reported to crack spontaneously when the average grain size is larger than ~3 μm and to completely disintegrate into powders when the grains are larger than 10

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μm. Therefore, crack-free PT ceramics can only be prepared as the grain seizes of the sintered samples are well controlled below 3 μm.

Figure 1. XRD patterns of the mixture of PbO and TiO2 powders milled for different times, using a SPEX shaker-mill operated at ~900 rpm (open circle: PbO, filled rhombus: TiO2, filled circle: PT) (Reproduced with permission from [9] Xue et al, Mater Lett 1999; 39: 364, Copyright Elsevier 1999)

Crack-free PT ceramics has been successfully prepared via a sol-gel process. By carefully controlling the calcination temperature of the derived gels and the sintering parameters of the sol-gel derived powders, the grain sizes of the PT powders and thus the sintered PT ceramics could be less than 0.2 and 1.8 μm, respectively. Since the anisotropic stress due to the phase transition was totally buffered by the grain boundaries, crack-free PT ceramics could be achieved, which are consistent with the prediction mentioned above. Preparation of crack-free PT ceramics has also been realized by other methods, such as seeding-assisted sol-gel process and refined ceramic process. Additionally, dense PT ceramics with submicrometer grains can be obtained by spark plasma sintering (SPS). The SPS process is able to sinter a compact powder at relatively low temperature in a very short time (a few minutes), which is very

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effective in suppressing exaggerated PT grain growth. However, it is found that special caution must be given to control the grain size of the PT powders when sol-gel processes are used. The SPS process uses expensive facilities with relatively low productivity. Comparatively, mechanochemical process is able to synthesize nanosized PT powders directly from oxide precursors.

Figure 2. TEM image of the PT powders synthesized via a mechanochemical process with SPEX shaker mill. (Reproduced with permission from [38] Wang et al, Solid State Ionics 1999; 124: 271, Copyright Elsevier 2001).

Indensity (a. u.)

0.75 0.85 1.00 1.20 1.40

20

30

40

50

2 Theta

60

70

Figure 3. XRD patterns of the PT powders derived from mixtures with different Pb/Ti molar ratios. (After [7] Kong et al, Prog Mater Sci 2008; 53: 207, Copyright Elsevier 2008).

Synthesis of PT powders via mechanochemical milling has been reported by several research groups, with different types of milling equipments [7-17]. Wang et al [9, 10] used SPEX vibrating mill to synthesize PT powders from oxides (PbO and TiO2) and amorphous Pb-Ti-O precursor derived from a coprecipitation process. Kong and co-workers [11, 12] used Fritsch Pulverisette 5 planetary high-energy ball mill for the preparation of PT powders from different TiO2 precursors. Fig. 1 shows the XRD patterns of the mixture of PbO and TiO2 powders milled for different durations, showing representatively the evolution of PT phase formation as a result of the high-energy milling [9]. Typical TEM image of the 20-h-milled

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sample is shown in Fig. 2 [38]. It is found that PT powders have a particle size of 20-30 nm. Experimental results indicated that the phase formation of PT is also affected by the properties of starting materials, especially during the early stage of phase evolution. There is also a fairly significant effect of PbO/TiO2 ratio on the phase formation of PT via the highenergy mechanical milling. Fig. 3 shows the XRD patterns of the mixtures with PbO/TiO2 (anatase) ratios from 1.4 to 0.75 [7, 17]. It is found that single phase PT can be formed from the mixtures with different PbO/TiO2 ratios. However, the tetragonality (c/a ratio) of the PT phases obviously decreases from ~1.03 for PbO/TiO2=1.4 to ~1.01 for PbO/TiO2=1.00. The samples from the mixtures with PbO/TiO2 ratio less than 1 are of cubic structure. It is important to note that the PT powders have almost the same grain size, although they have different tetragonalities. It means that high-energy ball milling technique is able to synthesize PT compound with compositions far away from the 1:1 stoichiometric state. Such a wide range of nonstoichiometry has not been reported for PT powders prepared by the conventional solid-state reaction method and wet-chemistry processing routes. This result is expected to be applicable to other lead containing ferroelectric compounds. More significantly, crack-free PT ceramics can be obtained from the synthesized PT powders [11, 12]. The reason is that the average grain size of the PT ceramics is about 1 μm and decreases gradually with increasing milling time, which is much less than the critical size of 3 μm, below which crack-free PT ceramics will be produced. The formation of crack-free PT ceramics can be attributed to the fact that PT powders synthesized via the high-energy mechanical milling process are of nanometer scale grains. Such powders possess very high sinterability, and hence, can be sintered at relatively low temperature. The low sintering temperature effectively suppressed the grain growth of the PT ceramics.

Figure 4. DTA curves of the PT ceramics derived from the nanosized powders. (Reproduced with permission from [12], Kong et al, Ferroelectrics 1999; 230: 281, Copyright Taylor & Francis 1999).

It is worth mentioning that the crack-free PT ceramics is not due to the contamination of tungsten carbide that comes from the vials and balls during the high-energy milling process, which has been conformed by TEM element analysis. The results indicated that the content of any impurities in the as milled powders in the sample milled for 80 h was less than 1 at% [17], which is similar to that reported by Jiang et al [19], who examined the tungsten

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contamination in ZrO2 powder milled for as long as 120 h with a tungsten carbide milling media and found that the tungsten content is less than 1 at%. Additionally, no second phases could be detected by the XRD measurement of the PT ceramics. This, together with the noncontamination statement, has also been supported by the differential thermal analysis (DTA) results of the PT ceramic pieces. As shown in Fig. 4, the DTA curve of the PT ceramics produced from the 20-h-milled powder clearly shows an endothermic peak at 491ºC during heating up and an exothermic peak at 471ºC during cooling down, respectively. The endothermic peak at 491ºC is the Curie temperature TC of PT, corresponding to the phase transformation from tetragonal ferroelectric state to cubic pyroelectric state. The exothermic peak at 471ºC is 20ºC below the endothermic peak, due to the hysteresis of the phase transition of PT. Almost similar DTA curves were observed for the samples made from powders milled for longer durations. If there were impurities that have been incorporated into the PT ceramics, their Curie temperatures would have been changed. In summary, high-energy mechanical milling has been shown to be a unique method in the preparation of PT powders and crack-free PT ceramics. As compared to most chemical processing routes such as sol-gel and hydrothermal method, high-energy mechanical milling technique is much simpler and more cost-effective.

3.1.2. Lead Zirconate Titanate

Figure 5. XRD patterns of the mixtures of PbO and TiO2 powders milled for different time durations. (Reproduced with permission from [23] Kong et al, J Mater Res 2001; 16 (6): 1636, Copyright Materials Research Society 2001).

Lead zirconate titanate [Pb(ZrxTi1-x)O3 or PZT] ceramics, solid solutions of PT and PZ, are important to a variety of applications such as transducers, sonars, micropositioners, rotary actuators and pyroelectric sensors, etc, due to its outstanding ferroelectric, piezoelectric, pyroelectric, and opto-electronic properties, and have been extensively and intensively studied for several decades [1-4]. Although PZT ceramics of different compositions have

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various functions, one of the important features of this solid solution system is the existence of almost temperature-independent phase boundary around x=0.52-0.53, which separates a rhombohedral Zr-rich phase from a tetragonal Ti-rich phase. The dielectric constant, piezoelectric constant and electromechanical coupling coefficient all exhibit a pronounced maximum value for the composition corresponding to this phase boundary, which is generally referred to as the morphotropic phase boundary (MPB) in the literature [1, 2]. This is mainly attributed to the existence of a mixture of phases at the boundary and the presence of a larger number of reorientable polarization directions existing in the MPB mixed-phase region [1, 7]. It is found that the microstructural and electrical characteristics of PZT ceramics made from the powders milled for different durations are significantly different. Using SPEX shaker mill, 10-h-milling leads to the formation of perovskite PZT with small amount of the unreacted oxides, while single phase PZT is obtained after milling for 25 h. The Milling also enhanced the reactivity of the precursors. For example, perovskite phase formation of the mixture milled for 10 h can be fully realized at 700ºC, while 800ºC is required by the unmilled sample. At the same time, there are intermediate phases, such as PT, appearing in the unmilled mixture calcined at relatively lower temperatures (600ºC-700ºC). These observations indicate that PZT phase formation was greatly enhanced as a result of the highenergy mechanical activation. The enhanced PZT phase formation can be readily attributed to the refinement of the oxide precursors as a result of the high-energy milling. Furthermore, PZT ceramics demonstrated better properties than those derived from the unmilled precursors.

o

Sintering rate (%/ C)

Linear shrinkage (%)

5 0 4h 8h 15 h 24 h

-5 -10 0.8 0.0 -0.8 -1.6 400

600 800 1000 o Temperature ( C)

Figure 6. Thermal mechanical properties of the PZT powders. (Reproduced with permission from [23] Kong et al, J Mater Res 2001; 16 (6): 1636, Copyright Materials Research Society 2001).

Similar results have been reported for PZT ceramics prepared by using a planetary mill. Combining with thermal mechanical analysis, it is found that the reaction temperature to form

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a perovskite structure can be greatly reduced as a result of high-energy milling. As shown in Fig. 6, the volumetric expansion temperatures, which is attributed to the formation of PZT from the constituent oxides, are gradually decreased from 771°C to 763°C and to 752°C for the powders milled for 4, 8 and 15 h, respectively [23]. The PZT ceramics prepared in this way have properties comparable with those fabricated via the conventional ceramic process. The high-energy milling technique has also been applied to PZT ferroelectric materials doped with various components [28]. Also, rapid formation of PZT is possible by simply increasing the milling intensity, i. e. high milling speed and large ball-to-powder weight ratio. In this way, PZT phase can be obtained by milling for just 1 h.

3.1.3. Lead Lanthanum Zirconate Titanate

Figure 7. XRD patterns of the mixtures for PLZT8/65/35 milled for different time durations: (a) 4 h, (b) 15 h and (c) 36 h. (Reproduced with permission from [31] Kong et al, Mater Lett 2002; 52: 378, Copyright Elsevier 2002)

Lead lanthanum zirconate titanate (PLZT) ceramics, with variable dopant concentrations of lanthanum and different Zr/Ti ratios, exhibit a variety of ferroic phases such as ferroelectric (FE), antiferroelectric (AFE) and paraelectric (PE) phases. As a result, PLZT ceramics are widely investigated during the last decades. The general formula of PLZT is (Pb1-yLay)(Zr1-xTix)1-y/4O3, usually being shorted as PLZT100y/1-x/x. A typical roomtemperature phase diagram of PLZT solid solution is quite complicated, where various phases exist with different compositions. The phases of great importance include antiferroelectric orthorhombic (AFEO), ferroelectric rhombohedral (FERh), ferroelectric tetragonal (FETet), relaxor ferroelectric (RFE) and paraelectric cubic (PECubic). Several PLZT ceramics, including PLZT8/65/36, PLZT7/60/40, PLZT2/95/5 and PLZT15/65/35, have been synthesized by the high-energy milling technique, showing the feasibility of the technique in synthesizing multicomponent materials [30-34]. Compared to PZT, the phase formation of PLZT is more difficult, which is because the presence of La element. As shown in Fig. 7, single phase

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formation of perovskite structure requires a time duration of 36 h. However, similar to PZT, PLZT ceramics with desired electrical properties can be well achieved from the incompletely reacted mixtures. For example, transparent PLZT8/65/35 ceramics can be obtained from the mixtures milled for 4, 15 and 36 h [31]. It is also found that excessive PbO that must be used in conventional solid-state reaction process is not a crucial requirement when PLZT ceramics are prepared from mixtures treated by a high-energy ball milling process. This conclusion should also be applicable to other PbO-containing materials.

3.1.4. Antiferroelectrics In antiferroelectrics, neighbour spontaneous polarization dipoles are anti-parallel to each other. The phase transformation of antiferroelectric ceramic from antiferroelectric (AFE) to ferroelectric (FE) phase can be induced by a sufficiently high electric field, which is usually accompanied by a large volumetric change since the unit cell of FE phase is larger than that of AFE phase. As a result, a very large longitudinal strain is associated with the AFE to FE phase transformation, as compared with the strains achieved by ferroelectric materials. This characteristic of antiferroelectric ceramics makes them good candidates for high displacement electromehanical actuator applications. In addition, the transformation of antiferroelectrics from AFE to FE phase leads to a significant energy storage, which can be used in energy storage applications. Antiferroelectric ceramics are also studied because of their unique pyroelectric, electrooptical and many other useful properties. It is difficult to synthesize antiferroelectrics, especially for those with multicomponent compositions, using chemical processing methods, such as coprecipitation and sol-gel process. Antiferroelectric materials synthesized via high-energy mechanical milling process include lead zirconate (PbZrO3 or PZ) [35], lead lanthanum zirconate titanate (PLZT2/95/5) [30], and La-/Nb-doped lead zirconate titanate stannate [Pb(Zr,Ti,Sn)O3 or PSZT] [36]. This success is especially is of special significance when more components are needed to modify the properties of the materials.

3.2. B-site Perovskite Relaxor Ferroelectrics and Their Derivatives Another important group of ferroelectric materials are lead-containing relaxor ferroelectrics also with the perovskite structure [1]. They have a general formula of Pb(B'B'')O3, where B’ is a low-valence cation, such as Mg2+, Zn2+, Fe3+ or Sc3+, and B” is a high-valence cation, such as Nb5+, Ta5+ or W6+ [1]. Relaxors are characterized by their high dielectric constant, broad and frequency-dispersive temperature dependence of dielectric constant. These properties can be ascribed to the so-called local compositional heterogeneity caused by the lack of ordering among the cations within the B-site sublattice. Their high dielectric constant, together with excellent electrostrictive response, makes relaxor ferroelectrics very useful for a wide range of applications, such as high dielectric constant capacitors, various sensors, transducers and actuators. The main problem for the synthesis of relaxors is the presence of pyrochlore, which seriously degrades the performance of relaxor ferroelectrics. To address this problem, two-step processing routes (e. g. columbite or Wolframite) are used to synthesize single phase or pyrochlore-free perovskite relaxors, via the conventional solid-state reaction process. Many chemistry-based methods, such as co-

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precipitation, sol-gel, gel combustion and molten salt, have also been developed to produce phase-pure ultra-fine powders of relaxor ferroelectrics. Comparatively, high-energy mechanochemical activation has been shown to be an alternative and effective process to prepare a wide range of phase-pure relaxor ferroelectric materials. Typical relaxor ferroelectrics include lead magnesium niobate [Pb(Mg1/3Nb2/3)O3 or PMN], lead zinc niobate [Pb(Zn1/3Nb2/3)O3, or PZN], lead iron niobate [Pb(Fe1/2Nb1/2)O3 or PFN] and lead scandium tantalate [Pb(Sc1/2Ta1/2)O3 or PST] (PMN), etc. Base on these basic relaxors, modifications by addition of other components such as PT, PZT, BT (BaTiO3) and ST (SrTiO3) or combination with each other have led to tremendous new types of materials with binary or ternary compositions.

3.2.1. Monophase Pyrochlore-free nano-sized PMN powders have been directly synthesized from the mixture of PbO, MgO and Nb2O5 powders, using SPEX shaker-mill (stainless steel) [37-42] and planetary high-energy ball mill (tungsten carbide) [43-46]. Nanosized PMN powders can be produced either from oxide precursors or coprecipitation-derived amorphous precursors. The formation mechanism of phase-pure PMN via the high-energy ball milling process is different from that via the solid-state reaction of oxides. The reaction is very complex in the solid-state process, where perovskite phase of PMN is not directly formed from the oxides. PbO and Nb2O5 firstly react at about 500°C resulting in a cubic pyrochlore (Pb3Nb4O13). The cubic pyrochlore compound further reacts with PbO leading to a rhombohedral pyrochlore (Pb2Nb2O7) at about 600°C. Pb2Nb2O7 then reacts with MgO to form perovskite PMN, with the appearance of Pb3Nb4O13 at higher temperature of 800°C. Although the amount of Pb3Nb4O13 phase can be reduced by adding excessive amount of PbO and MgO, it is very difficult to obtain single phase PMN by the conventional solid-state process. However, this problem is readily addressed by using the high-energy ball milling technique. The highenergy ball milling technique provides the milled system with energies high enough to trigger the reaction directly, avoiding the formation of pyrochlore phases, which are usually formed at high temperatures. Therefore, the mechanism that governs the formation of PMN via the mechanochemical activation is different from that of the solid-state reaction. Slightly different from the observations reported earlier, a recent study indicated that the reaction sequences of the formation of PMN from oxide precursors are relatively complicated [45]. Under the milling conditions used, it was found that pyrochlore-type compounds were also formed together with perovskite PMN. However, after milling for 60 h, almost single phase perovskite was obtained, with a very small amount of pyrochlore phase (~1 wt%). The formation of PMN was found to experience the following reactions. During the initial stage of milling, PbO massicot was transformed into PbO litharge and traces of hydrocerussite PbCO3·PbO·(H2O)2 were detected by XRD, while orthorhombic Nb2O5 and cubic MgO were kept unchanged in terms of phase structure. The presence of the complex compound could be due to the absorption of CO2 during the milling. At the same time, a particle size reduction and the formation of an amorphous phase, as a result of the high-energy milling, were observed in the mixtures. The degree of amorphization of PbO and Nb2O5 was found to be more significant than that of MgO, which means that the latter is “harder” than the former in this case. The amount of amorphous phase decreased with increasing milling time from 57 to

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42 wt% as the milling time increased from 8 to 48 h. Interestingly, the composition of the amorphous phase changed with milling time. With increasing milling time, the relative amount of PbO decreased until the milling was carried out up to 48 h, while the relative amount of Nb2O5 increased up to 32 h and after further milling up to 60 h it remained constant. The relative amount of MgO decreased from 8 to 32 h of milling and after that its relative amount increased. After milling for 8 h, pyrochlore phase was predominant. Only a small amount of Pb(Mg1/3Nb2/3)O3 were detected by the quantitative XRD analysis. The quantity of Pb(Mg1/3Nb2/3)O3 increased almost linearly with increasing milling time, which was accompanied by the consumption of MgO. The amount of MgO decreased slowly with increasing milling time. A small amount of MgO was still present in the powder after milling for 60 h. The reduction in the relative amount of PbO and Nb2O5 was due to the formation of the lead-niobium pyrochlore phase and the presence of them in the amorphous phase. As the lattice parameter of the pyrochlore decreased with milling time, it was highly suggested that an isostructural PMN pyrochlore was formed, which reacted with PbO and MgO into perovskite. The maximum concentration of pyrochlore was reached after milling for 16 h. The formation of perovskite was significantly accelerated by the high-energy milling process. Perovskite phase was also promoted by thermal annealing from the mixtures milled for short times [45]. This difference is understandable, since the outcome of a high-energy milling could be determined by a number of factors. It is of particular interest to produce single-phase PZN powder via a mechanical milling process because no one has succeeded in trying to synthesize phase-pure PZN powder by either the conventional solid-state reaction process or chemistry-based routes. So far, molten salt flux is the only technique to grow phase-pure PZN in the form of single crystal. This is because that the conventional solid-state reaction process and most chemistry routes require calcination temperatures of 600-900ºC. In this temperature range, PZN is unstable. The fact that PZN crystal can be grown in PbO flux at high temperature is due to the unequilibrium state. Phase formation of PZN was reported by Shinohara et al [66], where the authors used a soft-mechanochemical route to prepare PMN-PZN ceramics. Although they noted the formation of perovskite PZN by the mechanical milling, they didn’t pay much attention to the observation and made no attempt to synthesize phase-pure PZN powder. The synthesis of perovskite PZN via mechanochemical process was formally reported by Wang et al [47, 48]. They found that nanocrystalline PZN powders could be synthesized either from oxide mixture of PbO, ZnO and Nb2O5 or a mixture of PbO and ZnNb2O6. The phase evolution of PZN is similar to that of PT, PZT and PMN, as discussed above. The formation of single phase PZN indicates again the difference in formation mechanism between mechanochemical synthesis and conventional solid state reaction process. However, due to its poor thermal stability, PZN powders cannot be used to prepare phase-pure PZN ceramics. Lead iron niobate [Pb(Fe1/2Nb1/2)O3, or PFN) and lead iron tungstate [Pb(Fe2/3W1/3)O3 or PFW] are important relaxor ferroelectric materials. Due to their high dielectric constants, broad ferroelectric-paraelectric phase transition, and especially relatively low sintering temperature, PFN and PFW have been acknowledged to be promising candidates for application of multilayer ceramic capacitors (MLCs). The low sintering temperature of PFN and PFW makes it possible to use low melting temperature inexpensive inner electrodes, leading to a significant reduction in the cost of MLCs products. Also, these iron containing ferroelectric materials are of particular interests due to their possible magnetic properties. For

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example, PFN has a diffuse ferroelectric phase transition at ~380 K and an antiferromagnetic phase transition at ~145 K. PFW ceramic sample demonstrates a Curie temperature (TC) of ~175 K and antiferromagnetic-to-paramagnetic at ~340 K. These multiferroic magnetoelectric materials simultaneously show electric and magnetic polarizations. The coupling between ferroelectric and magnetic activity offers the possibility to manipulate the magnetic properties through electric field and vice versa, making these materials potential for a wide range of applications in spintronics, multiple state memory elements or memory devices, which use electric and/or magnetic field for read/write operations. The synthesis of iron containing relaxor ferroelectric phases, via high-energy mechanical milling process, has been demonstrated by Wang et al [49-52] and other authors [64, 65]. Unlike PMN and PZN, the formation of PFN and PFW follows a relatively complicated reaction sequence. It has been shown that single phase PFN powder can be formed from oxide mixture via a mechanical activation. But PFN ceramics derived from this PFN powder demonstrated poor microstructural and electrical properties. High performance PFN ceramics can only be obtained from the columbite precursor of PbO and FeNbO4. PFW cannot be directly synthesized from oxide mixture of PbO, Fe2O3 and WO3 by using a mechanochemical activation [49]. Instead, nanocrystalline lead tungstate (PbWO4) and pyrochlore (Pb2FeWO6.5) are formed in the oxide mixture as a result of the mechanical activation. However, the subsequent thermal treatment showed that the reactivity of the milled oxide mixture is enhanced for the formation of PFW perovskite phase. This is not surprising because the precursor have been significantly refined as a result of the high-energy milling. Although the reduction in calcination temperature leading to the formation of perovskite PFW phase is not very significant, the milled powders result in PFW ceramics with much better electrical properties as compared to those from the unmilled samples. PFW ceramics, with high density, uniform microstructure and better dielectric properties, are obtained from the milled powders, while the ceramics derived from the unmilled mixture have high porosity, rough microstructure and poor electrical properties. This observation is similar to that observed in PFN. Single phase PFW can also be formed from oxide mixture of PbO, Fe2O3 and WO3 with 0.4 mol PFW powder as seeding, via a high-energy mechanical activation [49]. This seeding level is much higher than those for solution processing, which is because the diffusion rate of species in a liquid solution or in a solid state at high temperature is much higher than that in the case of the high-energy milling, since the milling was carried out at relatively low temperatures. Another way to synthesize PFW is to use pre-formed constituents, such as Fe2WO6, PbWO4 and Pb2FeWO6.5 [52]. Lead scandium tantalate [Pb(Sc1/2Ta1/2)O3 or PST] is a typical candidate for investigating the relationship between structural ordering and electrical characteristic of relaxor [53]. Ordered PST behaves as a normal ferroelectric material, whereas disordered PST is a relaxor having properties such as frequency dependent dielectric constant and broad phase transition temperature range. The degree of the B-site ordering in PST can be controlled by thermal annealing, which has been acknowledged to be a very useful way to tailor the electrical properties of PST. The definition and characterization of B-site ordering and disordering will be discussed later. Single phase PST of perovskite structure has been synthesized from oxide mixture via a mechanical milling [53]. 20-h-milling leads to the formation of single phase PST. However, although the PST phases derived from the mechanical activation and the solid-state reaction

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both possess the perovskite structure, they have different ordering characteristics. As shown in Fig. 8, the superlattice reflection (111), an indication of ordering characteristic, is observed in the sample derived from the Wolframite precursor via the conventional solid-state reaction process, while no such peak is detected in the mechanochemically synthesized phase. The former has an ordering parameter of 0.729, in contrast to a totally immeasurable order parameter of the later. Due to the high sinterability, the mechanical activated derived PST sample has a relative density of 97.1%, as compared to 92.6% of the ceramics made from the Wolframite precursor. After sintering at 1200ºC for 2 h, their ordering factors increase to 0.764 and 0.23, respectively. The difference in the disordering characteristics between the PST ceramics is also confirmed by the difference in their dielectric properties [53]. The 1200ºC-sintered PST ceramic sample derived from the 20-h-activated oxide mixture demonstrates a typical relaxation behavior. A maximum dielectric constant of ~14000 was measured at 0.1 kHz, with a dielectric loss tangent of 0.032, which are better than the dielectric properties of its Wolframite counterpart. The most significant aspect of these observations is that the disordering character in the mechanical activated PST can be retained after sintering at high temperatures. Therefore, the mechanochemical technique offers us an effective and simple way to prepare disordered PST ceramics, without any special treatments used in the conventional solid-state reaction process.

Figure 8. XRD patterns of the PST phases derived from 20-h-milled oxide mixture and from Wolframite precursor. (Reproduced with permission from [53] Lim et al, Mater Chem Phys 2002; 75: 157, Copyright Elsevier 2002)

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3.2.2. Binary-Phase Numerous binary systems have been created through the combination of relaxors with ABO3 compounds, such as BT, ST and PT. Combination of relaxor and relaxor is also an effective way to produce materials with desired electrical properties. Representative binary phases, such as PMN-PT [54-58, 63], PZN-BT [59-62] and PMN-PFN [67], will be discussed to show the feasibility of high-energy mechanical activation in synthesizing ferroelectric materials. Three subsections, i. e. PMN family, PZN family and relaxor-relaxor combination, will be discussed. (1-x)PMN-xPT binary phases with different compositions demonstrate various dielectric, ferroelectric and piezoelectric properties. 0.9PMN-0.1PT is a pronounced candidate to replace BT (BaTiO3) in MLCs (multilayer ceramic capacitors) because it has lower sintering temperature as well as higher dielectric constant than BT. It is also one of the best candidate materials for devices of actuators because of its large electrostrictive and piezoelectric coefficients. (1-x)PMN-xPT system has also a morphotropic phase boundary (MBP), similar to that in PZT family [1]. The piezoelectric properties of the (1-x)PMN-xPT system, however, are much higher that of PZT system. 0.65PMN-0.35PT has been shown to exhibit the highest piezoelectric properties among various ferroelectric materials, thus making it important for applications of many actuators and sensors. The first example to synthesize 0.9PMN-0.1PT using mechanochemical process was reported by Baek et al [55], with a primary object to improve the performance of the ferroelectric ceramics. More recently, it has been reported that single phase nanocrystalline 0.9PMN-0.1PT powders can be directly synthesized from oxide precursors via high-energy mechanical milling process [54-58]. The phase evolution of PMN-PT follows a similar pattern of PMN, as discussed above. The presence of pyrochlore phase is not observed. As summarized in Table 6, the 0.9PMN-0.1PT ceramics, derived from the high-energy mechanically synthesized nano-sized powders, demonstrated very promising dielectric and ferroelectric properties. Compositions with high contents of PT have also been reported to be synthesized using high-energy milling technique. Similar to PZT, the synthesis of PMN-PT powders can be readily speeded up by increasing the milling rates. As discussed above, although single phase PZN powder has been synthesized from oxide precursor via mechanical activations, it is impossible to have phase pure PZN ceramics from the synthesized powder since it is not stable at high temperatures. However, it has been reported that perovskite structure of PZN can be stabilized by the addition of BaTiO3 (BT), SrTiO3 (ST) and PbTiO3 (PT) to form binary solid-state solutions of PZN-BT, PZN-ST or PZN-PT, via the conventional ceramic process or single-crystal growth. PZN based binary materials have many outstanding properties. For example, 0.91PZN-0.09PT, which is near the morphotropic phase boundary (MPB) at room temperature shows a surprisingly large dielectric and piezoelectric constant and higher electromechanical coupling coefficient than the PZT family of ferroelectrics. Among various additives, BT is the most effective additive to suppress the formation of pyrochlore phase in PZN binary ceramic materials. (1-x)PZNxBT with x=0.05-0.30 have been synthesized via high-energy mechanical milling process [54-58]. The results indicated that the mechanochemical technique is advantageous over some of the chemical processing routes for the preparation of PZN based ceramics. For example, special carefulness must be taken to obtain high percentage of perovskite 0.75PZN-0.25BT.

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Although high chemical homogeneity was achieved in the synthesis method, the maximum weight percentage of perovskite was 96%, which was only achieved under proper thermal treatment (14 min calcination at 800ºC after at fast firing rate of 50ºC/min). Comparatively, the mechanochemical process is much simple and effective. There is also report about the synthesis of other PZN based binary compositions, such as 0.92Pb(Zn1/3Nb2/3)O3-0.08PbTiO3 (0.92PZN-0.08PT). Relaxor-relaxor combination has shown many particular interests. For example, PMN has a Curie temperature TC of ~50ºC, where as PZN’s TC is ~140ºC. It is, therefore, possible to make a relaxor with TC close to room temperature by combining PMN and PZN. This kind of material is expected to have very high room temperature dielectric constant, which is important for the fabrication of small dimension ceramic capacitors with high values of capacitance. Although normal ferroelectrics such PT, BT or PZT can be used to modify the properties of relaxors, they alter the relaxor characteristics at the same time. For example, with increasing concentration of BT, PZN-BT will become normal ferroelectric gradually. As discussed above, single phase PFW cannot be directly synthesized from oxide mixture via a mechanochemical process, but the mechanically activated mixture can be used to fabricate PFW ceramics because single phase is formed at high temperature. In contrast, PZN has been successfully derived from the oxide mixture. However, PZN is not stable at high temperature. It is expected that an appropriate combination of PFW and PZN might lead to a relaxor that can be obtained from oxide mixture via a mechanochemical activation and is also stable at high temperatures. so that PZN-PFW ceramics Successful applications of mechanochemical synthesis to these kinds of relaxor binary systems will be presented as following. A soft-mechanochemical processing was used to treat precursors of PbO, Mg(OH)2, Nb2O5 and 2ZnCO3·3Zn(OH)2·H2O, in order to synthesize PMN-PZN ceramics [66]. It was found that perovskite phase was formed even after milling for only 30 min and became major phase after 3 h of milling. About 80% perovskite phase was achieved for (1-x)PMN-xPZN with x≤0.9 and ~60% for pure PZN. The thermal stability of the as-milled PMN-PZN powder is greatly dependent on the content of PZN. With increasing PZN, the thermal stability of the PMN-PZN decreases. Phase-pure PMN-PZN ceramics can only be obtained as the concentration of x<0.7. Another example is (1-x)PFW-xPZN.

3.2.3. Ternary-Phase The electrical properties of relaxor ferroelectrics can be significantly modified using multi-component compositions. Representative ternary systems prepared by the high-energy mechanochemical process are 0.54PZN-0.36PMN-0.1PT and 0.48PFN-0.36PFW-0.16PZN, reported by Wang’s group [68-71]. The results suggested that it is not necessary to use the columbite precursor to synthesize PZN-PMN-PT powder via a high-energy mechanochemical milling process and the PFN-PFW-PZN ceramics derived from the 20-h-milled mixture, had better dielectric properties, as compared to that from its unmilled counterpart (<500).

3.2.4. Order-Disordering Transition Induced by Mechanical Activation A schematic diagram, showing an ordered and a disordered crystal structure of A(B'1/2B''1/2)O3, is demonstrated in Fig. 9 [7]. In the ordered structure (Fig. 9 (a)), sublattice B' or B'' can be readily identified, while in the disordered one (Fig. 9 (b)), no sublattice can be

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established, because B' and B'' ions are distributed randomly. The degree of ordering (disordering) is dependent mainly on the difference in size of the two B-site ions and the difference in their valence states. A large difference in size and valence state is favorable to B-site ordering. Depending on their scales of dimension, B-site ordering can be classified into three groups: (i) random cation distribution for coherence length below 2 nm, (ii) short coherence long-range ordering for nanoscale of 2-50 nm and (iii) long coherence ordering of above 100 nm. It has been found that an order-disorder transition can be triggered by highenergy activations in the conventionally synthesized ordered powders of monophase PST [72] and binary system 0.4Pb(Mg1/3Nb2/3)O3-0.6Pb(Mg1/2W1/2)O3 (0.4PMN-0.6PMW) [73, 74]. Due to the fact that the grain sizes of the PST ceramics were far beyond nanometer scale, the disordering and dielectric behaviors has been considered to be associated with the presintering mechanical activation in the sintered PST ceramics, and not related to particle size. This finding offers us an opportunity to tailoring the disordering of PST and PMN-PMW and other relaxor ferroelectric ceramics through proper combinations of mechanical activation and subsequent sintering.

Figure 9. Schematic diagram of crystal structure for A(B′1/2B″1/2)O3: (a) ordered and (b) disordered (B′: larger open circle and B″: smaller filled circle). (Reproduced with permission from [7] Kong et al, Prog Mater Sci 2008; 53: 207, Copyright Elsevier 2008)

3.3. BaTiO3 and CaTiO3 Barium titanate (BaTiO3 or BT) is the first ferroelectric ceramics [1, 75-83], which is a good candidate for a variety of applications, such as piezoelectric actuators, multiplayer ceramic capacitors (MLC) and positive temperature coefficient resistors (PTCR), due to its excellent dielectric, ferroelectric and piezoelectric properties [1]. BaTiO3 powders were conventionally synthesized by solid-state reaction between BaCO3 and TiO2 at temperatures higher than 1200°C [1]. The high calcination temperature required by solid-state reaction process leads to many disadvantages of the BaTiO3 powders, such as large particle size, wide size distribution and high degree of particle agglomeration. In this regard, it is desired to lower preparing temperature in order to get BaTiO3 powder with fine and homogenous structures. Various chemistry based methods have been developed to

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prepare BaTiO3 at low temperatures. Examples include chemical co-precipitation, sol-gel process, hydrothermal, molten salt, microemulsion and auto-combustion. It was reported by Kong et al [77] that single phase BT cannot be synthesized directly from BaCO3 and TiO2 mixture by high-energy activation, but it brought out a significant reduction in phase formation temperature of BT from the mixture. As shown in Fig. 10, well developed single phase BT is observed in the sample after calcining at 800ºC for 2 h. This temperature is much lower than that required by a standard solid-state reaction process. A later finding showed that single phase BT can be synthesized from oxide mixture of BaO and TiO2 via a mechanical milling only under nitrogen environment or vacuum. The phase evolution of perovskite BT is similar to those of PT, PZT and PMN and the BT powder obtained is of nanometer size and narrow size distribution. High-energy mechanical milling can also reduce the phase formation temperature of BT by altering the reaction process. Recently, however, there was a report showing that perovskite BT can be directly synthesized from BaO and TiO2 via mechanical milling in air. There are also reports of synthesis of nanosized Ba1-xSrxTiO3 (BST) with a whole composition, using mechanochemical technique [83, 123]. CaTiO3, with a similar structure to BaTiO3, have also been synthesized via the high-energy technique [84, 85].

Figure 10. XRD patterns of the mixtures of BaCO3 and TiO2 powders milled for 10 h calcined at different temperatures: (a) 600ºC, (b) 700ºC, (c) 800ºC and (d) 900ºC for 2 h. (Reproduced with permission from [77] Kong et al, J Alloy Comp 2002; 337: 226, Copyright Elsevier 2002)

3.4. Aurivillius Ferroelectrics Aurivillius type structure compounds have a general formula [Bi2O2][An-1BnO3n+1], which are built up by n pseudo-perovskite [An-1BnO3n+1]2- layers alternating with [Bi2O2]2+ layers. These materials have received great interest due to their high Curie temperatures and excellent piezoelectric properties. Various ferroelectric materials of Aurivillius family have

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been synthesized via high-energy mechanical milling process. Some of them can be produced directly from oxide mixtures, while others are crystallized from amorphous of oxide precursors created by high-energy activations.

3.4.1. Bi4Ti3O12 [86-89] Bismuth titanate (Bi4Ti3O12, BiT) is the most famous Aurivillius type ferroelectric material with n=3. It is a good candidate for high-temperature piezoelectric applications, memory storage, and optical displays because of its high Curie temperature (675°C) and good electro-optical switching behavior. Nano-sized BiT powders have been synthesized from oxide mixture via a high-energy planetary ball milling process. BiT ceramics with good dielectric and pyroelectric properties are obtained from the nano-sized BiT powders. Although BiT nanosized powders have been successfully synthesized via high-energy mechanochemcal activations, no attempt has been made to examine and explain the structural change in detail. The structural variation of the system 2Bi2O3·3TiO2, as a result of highenergy milling, was carefully studied by Zdujic et al in a recent article [89]. To have a better understanding of the effect of high-energy milling on the structural evolution, the authors also milled BiT powders synthesized via a reactive sintering method separately. Both the oxide mixture (Bi2O3 and TiO2) and the BiT powder were milled using a Fritsch Pulverisette 5 type planetary ball mill. Stainless steel vials of 500 ml and hardened-steel balls of 13.4 mm in diameter were used as milling media. The ball-to-powder weight ratio was 20:1. Two sets of milling experiments were used to check the effect of milling power on phase evolution of BiT. In the first set, the powders were milled up to 15 h initially at an angular velocity of the basic disc of 180 rpm and vials of 225 rpm. After that, the angular velocities of both basic disc and vials were increased to 317 rpm and 396 rpm, respectively. The samples were milled for extra time durations of 1, 3 and 5 h, after which additional 10 h milling was applied at basic disc and vial velocities of 180 rpm and 225 rpm (the initial speed). The second milling set conducted at the maximum speed (317 and 396 rpm) for 20 h, followed by an additional 10-h-milling at low speed (180 and 225 rpm). It was found that the formation of BiT phase in the 2Bi2O3·TiO2 mixture took place through the crystallization of an intermediate phase Bi2(CO3)O2, which was not observed by other researchers [86-88]. The formation of this intermediate compound was found after milling for 1 h. Further milling resulted in the formation of amorphous BiT phase. The mechanochemical reaction between 2Bi2O3 and TiO2 was triggered due to the creation of highly reactive fresh surfaces of particles and interfaces between the reactants. An additional high intensity milling induced a partial crystallization of the amorphous BiT phase. Prolonged milling for another 10 h at low speed led to amorphization of the crystalline BiT. The structural evolution of BiT powder as a result of the high-energy milling was similar, with only slight differences in degree and time of amorphization and crystallization transition. In this study, the kinetics of the reaction and phase transition have been systematically investigated.

3.4.2. Other Aurivillius Type Ferroelectrics Besides Bi4Ti3O12, many other ferroelectric materials belonging to the Aurivillius family have also been produced using high-energy milling techniques. They include bismuth

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vanadate (Bi2VO5.5 or BiV) [90-94], bismuth molybdate tungstate (Bi2Mo1-xWxO6 or BiMW) [95], calcium and strontium bismuth titanate (CaBi4Ti4O15 and SrBi4Ti4O15 or CBiT and SBiT) [96-98], bismuth titanate niobate (Bi3TiNbO9 or BiTN) [99-101] and (1-x)Bi2SrNb2O9xBi3TiNbO9 [(1-x)SBN-xBTN] [102-105]. Some compositions were made into ceramic samples and characterization results indicated that the ceramics made from the nanosized powders exhibited good dielectric and piezoelectric properties.

3.5. LiNbO3 and NaNbO3 Lithium niobate (LiNbO3 or LN) is a ferroelectric material with a variety of attractive properties, such as high pyroelectric, piezoelectric electro-optical and photo-elastic coefficients, large acoustic-optic figure-of-merit, and significant photorefractive effects. Conventional synthesis of LN requires high temperatures. Similar to lead containing ferroelectric materials, the high temperature required causes loss of Li element. The loss of stoichiometry by Li2O evaporation, together with the formation of Nb-rich grain boundaries is seriously harmful to the planar coupling coefficient (Kp) and piezoelectric coefficient (d33) of LN ceramics. Although ultrafine LN powders can be produced by wet-chemical routes, such as sol-gel, coprecipitation and Pechini’s method, problems are similar to those encountered by lead containing ferroelectrics. Therefore, high-energy milling is an alternative way to synthesize LN powder. The work reported by de Figueiredo et al [106] is probably the only example. The starting materials used were Nb2O5 and Li2CO3 and milling duration varied from 2 to 42 h. A planetary mill (Fritsch Pulverisette 5) and stainless steel milling media were used [106]. Sodium niobate (NaNbO3 or NN) is another example of niobate ferroelectric materials synthesized via mechanochemical milling, reported by Castro et al [107-110]. Unlike LN, NN crystalline phase cannot be directly synthesized from the mixture of Na2CO3 and Nb2O5, using a vibrating mill (Fritsch Pulverisette 0). A post thermal anneal, at a temperature though much lower than that required by the conventional ceramics process, is necessary to facilitate the phase of NN. NN ceramics prepared from the powder possessed better piezoelectric properties.

4. Mechanisms Phase formations by thermal activation in the conventional solid-state reaction and wet chemistry-based process are through various interfacial reactions or diffusions at boundaries between/among the precursor components, where one or more intermediate phases were usually formed preferentially prior to the formation of desired compounds. In contrast, the direct production of nano-sized ferroelectric phases activated by various high-energy mechanochemical milling involves nucleation and subsequent growth of the ferroelectric crystallites, without the occurrence of interfacial reactions and diffusions. However, reactions to form ferroelectric phases, caused by high-energy mechanochemical activation, could a complicated process, which is significantly different from the solid-state reactions at high temperatures or those in solutions with a molecular dispersion. It has been accepted that the phase formation of ferroelectric materials via highenergy milling experiences two stages. Firstly, the grain/particle sizes of the precursors are

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greatly refined. The fragmentation results in high defect densities, shorter diffusion distances, and more intimate contacts of precursors. At the same time, large amounts of fresh/cleansed surfaces/interfaces are created. All these factors will increase the reactivity of the precursors. After a certain period of activation, depending on the materials involved and facilities used, spontaneous reactions towards desired ferroelectric phases start to occur at some spots where the reactants are sufficiently reactive enough or where localized temperature is high enough to trigger the reactions. The phases formed serve as nucleation and cites of grain growth. Amorphization is occasionally also observed in the synthesized powders, depending on milling conditions and/or materials characteristics [39, 41, 62, 89, 123]. Also, the very high localized-pressure created by the strong collision during the milling process could be an additional factor to facilitate the reactions [62, 112]. It is reported that, when a planetary mill is used to synthesize lead containing ferroelectric compounds, the final products are usually hard agglomerates, instead of loose powders [11, 22-26, 30-36, 119]. The hard layers are strongly stuck at the bottom of tungsten carbide vials. However, the agglomerates can be easily pulverized into powders. It has been found, from the surface SEM images, of the pieces of PbZr0.7Ti0.3O3 (PZT) agglomerates derived from oxide mixture of PbO, ZrO2 and TiO2, after milling for 20 hours, that there are areas with two types of morphologies, dense and porous [26]. The dense area consists of rodlike particles with a length of 0.2-0.5 μm and a thickness of less than 100 nm. The appearance of such hard layers means that sintering of the formed PZT powder occurred during the milling process. This is a strong support to the suggestion of high localized-temperature caused by the high-energy ball milling [9]. In this case, PZT phase is first formed incidentally at the bottom of the WC vial where the flying balls could collide at all time. After a certain number of nanosized PZT crystallites were formed, they would act as nuclei from which grain growth takes place. High local temperature and high pressure produced by the collision could cause the growth and densification of the PZT grains.

Figure 11. Schematic drawing of free energy for Pb(Zn1/3Nb2/3)O3 (PZN). (Reproduced with permission from [114] Wakiya et al, Mater Res Bull 1995; 30: 1121, Copyright Elsevier 1995)

The formation of PZN nanosized powder from oxide constituents, via high-energy activation, is a miracle, since there is never such report via solid-state reaction or chemical synthesis methods. It is well known that single phase PZN crystals can only be formed in excess PbO at high temperature [114], which has been successfully explained by the Ostward’s step rule, as shown schematically in Fig. 11 [114]. The PZN single crystal is precipitated from supercooled melt under unequilibrium state. The variation in free energy in

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the molten system follows the thick line from high temperature. At a certain temperature, the supercooled melt becomes supersaturated so that crystallization will take place in the supercooled melt. According to the Ostwald’s step rule, the phase precipitated from the supercooled melt is that with smaller free energy difference from the supercooled melt. Therefore, PZN perovskite, rather than pyrochlore, is precipitated from the molten state, as indicated by the arrow in the figure. Since the free energy of perovskite PZN is higher than that of pyrochlore, the precipitated PZN is in a metastable state. As a result, the PZN crystals grown from excess PbO flux will be decompose to pyrochlore phase, a state of lower free energy. However, Fig. 11 also indicates, the free energy-temperature curves of perovskite and pyrochlore cross each other at a point. Below this point, perovskite has a lower free energy than pyrochlore. According to Jang et al [113], this temperature is ~600ºC. Therefore, the synthesis of PZN via the high-energy mechanochemical process takes place at a temperature at least not high than 600ºC. Otherwise, PZN cannot be obtained via a high-energy milling. This result also suggests that the localized temperature during the high-energy ball milling should not be too high (>800ºC). Of course, this suggestion is not applicable to the cases of milling induced combustion [111]. In summary, most of the lead-containing ferroelectric powders can be directly synthesized from their oxide precursors as a result of high-energy mechanochemical process. Because the high-energy milling effectively suppresses the grain growth of the synthesized phases at the same time, ferroelectric powders produced in this way possess nanosized crystalline grains, thus having very high sinterabilities. Ferroelectric ceramics derived from the nanosized powders have demonstrated promising electrical, dielectric, ferroelectric, piezoelectric and pyroelectric properties. Some ferroelectric compounds cannot be directly obtained via a high-energy milling process, but their phase formation temperatures of these compounds can be significantly reduced as compared to that required by conventional solid-state reactions from unmilled precursors. This can be called activation-assisted phase formation. This assisting effect of a high-energy milling is readily attributed to the refinement of the precursors by the mechanical activation. The refinement not only means the reduction in grain or particle size, but also means the creation of defects, dislocations, lattice distortions and microstrains in the precursor lattice. An extreme case of the refinement is amorphization of the precursors. Amorphous metallic alloys or metallic glasses are solid alloys, with a liquid-like or noncrystalline atomic structures, have been extensively investigated [115-118]. Three critical requirements for the production of metallic glasses, via rapid solidification method, are (i) multiconmponent systems with three or more constituent elements, (ii) significantly different atomic size ratios typically with difference exceeding ~13% and (iii) negative heats of mixing among constituents. Recently, mechanochemical activation is found to be very effect way to produce amorphous metallic alloys. Metallic alloys can be even made in the systems with a positive heat of mixing. Mechanical alloying can produce amorphous alloys with wider composition ranges than those made by rapid solidification process. The mechanochemical activation has also been applied to amorphization of semiconductors, such as Se [116], Si [117] and SiC [118], oxides, such as mullite and zeolite. High-energy milling can induce not only accumulation of vast lattice and point defects but also chemical disordering of lattice. If the rate of dynamic recovery is less than the rate of defect production, the accumulation of the topological and chemical disorders will lead to a collapse of crystalline structure.

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5. Applications of Nanosized Ferroelectrics 5.1. Powder Refining Besides the direct synthesis and assisted synthesis of ferroelectric powders, high-energy milling has also been used to treat commercial and lab-synthesized ferroelectric materials, such as PZT [119, 120], BT [121] and BST [122]. It is found that high-energy milling is an effective way to modify the morphology, reduce the grain/particle sizes and thus enhance the densification behavior of ferroelectric powders. For example, a commercial PZT powders, consisting of large particles (tens of micron) with grains of about 1 µm, can be refined to less than 100 nm, after milling for 10 h, using a planetary high-energy mill [119]. The milled PZT powders can be sintered at a temperature of 100°C-200°C lower than that required by the unmilled powders. The enhanced sinterability of commercial PZT powders is of interest to thick film and multilayer PZT devices. Thick films are usually fabricated by screen-printing technique with PZT paste consists of PZT powders, sintering aids and organic binders. The sintering aids, which are glasses with low melting temperature, are used to lower the sintering temperature of thick films by forming liquid phase during sintering process. These glasses consisting of non-ferroelectric phases, however, are harmful to the performance of the final devices. As a result, PZT powders that can be sintered at low temperature (<1000°C) without any glasses additives will be attractive to thick films PZT based devices. Such powders are also important to multilayer PZT devices for at least two reasons. First consideration is cost reduction. Multilayer PZT devices are fabricated in such a way that the PZT layers and electrode layers are alternatively stacked via tape casting and screen-printing and then co-fired (sintered) at high temperatures. If the sintering temperature is too high (>1000°C), only expensive metal like Pt, Au, Pd and Ag can be used as electrode layers, leading to higher cost devices. The reduction of device cost needs to use cheaper metals like Ni, Cr, and Cu as electrode materials. In this case, the sintering temperature of ferroelectric materials must be sufficiently low. Additionally, with the development of modern microelectronics, multilayer structured ferroelectric devices have to be able to work at low-driving voltage, to meet the requirement of miniaturization and hybridization. The pre-requisition is to reduce the thickness of the PZT active layer from hundreds (currently) to tens of micrometers. As the thickness of ferroelectric layer is reduced to 20 μm or less, the grain size of ferroelectric materials should not be larger than 1 μm. Otherwise, the electrode will be very easily shorting through the ferroelectric layers and hence the yield of the device will drop. Therefore, low temperature sintering is also desired to prevent grain growth.

5.2. Thick Films There has been report that nanosized ferroelectric powders can be used to fabricate thick films using screen-printing technique [123]. Nanosized 0.65Pb(Mg1/3Nb2/3)O3-0.35PbTiO3 powders were synthesized from oxide constituents, using a planetary high-energy mill [45]. Analysis of sintering behavior indicated that the nanosized PMN-PT powder started to shrink at 800°C. Due to the presence of ~30 wt% amorphous phase, the synthesized powders were

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annealed at 700°C to promote a full crystallization before being used to prepare thick films. The presence of the amorphous phase is probably because the milling speed used to synthesize the PMN-PT powder is too high (300 rpm). The contents of amorphous phase in the ferroelectric powders, reported by other researchers, are much lower [54, 56].

Figure 12. Schematic drawing of free energy for Pb(Zn1/3Nb2/3)O3 (PZN). (Reproduced with permission from [123] Kosec et al, J Eur Ceram Soc 2007; 27: 3775, Copyright Elsevier 2007).

Figure 13. Dielectric properties of the PMN-PT thick films annealed at different temperatures. Inset (a) shows the real part of permittivity of the film annealed at 850ºC. Inset (b) shows the imaginary permittivity of the thick films annealed at different temperatures. (Reproduced with permission from [123] Kosec et al, J Eur Ceram Soc 2007; 27: 3775, Copyright Elsevier 2007).

PMN-PT thick films were fabricated on Al2O3 substrate using the screen-printing technique with paste made from the nanosized powders. Fig. 12 shows the SEM images of the

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thick films annealed at 850°C, 900°C and 950°C, respectively. It is found that thickness decreases with increasing annealing temperature, due to the densification of the films. At the same time, significant grain growth can be observed. Almost fully dense thick film with a thickness of about 35 µm is obtained after annealing at 950°C for 2 h. Fig. 13 shows the dielectric constant of the thick films annealed at different temperatures. Room temperature dielectric constant of the film is 4100, while the peak value of dielectric constant at TC is as high as 41000. The 950°C-annealed sample has a remanent polarization of 26 µC/cm2, a coercive field of 5.7 kV/cm, a piezoelectric constant of 170 pC/N. All these properties of the PMN-PT thick film are comparable with those of bulk ceramics, making it to be potential candidates for applications of micro-piezoelectric devices.

5.3. Nanocomposites A very interesting approach, called nanocomposite processing route, was proposed to incorporated high-energy ball milling technique with sol-gel process to deposit thick films, which cannot be realized using the typical sol-gel with solution precursors [124-128]. This approach has combined advantages of high-energy milling and the sol-gel process.

Figure 14. Flow chart of the processing sequence to prepare uniformly dispersed suspensions or slurries. (Reproduced with permission from [124] Wang et al, Mater Chem Phys, 2002; 75: 71, Copyright Elsevier 2002).

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Fig. 14 shows the processing sequence of the nanocomposite technique [125]. In this processing, commercial PZT powder was milled using a high-energy mill to produce nanosized PZT powder, which similar to that reported in Ref [119]. The nanosized PZT powder was then milled for one more time with appropriate dispersants, which are used to modify its surface characteristics. The modified PZT nanosized powder can be readily mixed with sol-gel solutions without the presence of agglomerations or precipitations. This kind of mixture can be used as same as the typical sol-gel solutions to deposit thick films on various substrates. The slurries may also directly used to other processing, such as tape casing, screen printing and molding [125]. The outstanding features of this technique are summarized in Fig. 15 [125]. Due to these basic characteristics, the technique can be used for various purposes. For example, if a sol-gel solution has the same composition as the nanosized powder, the resulting production will have single phase, while if different compositions of solution and powder are used, multiphase composites will be produced.

Figure 15. Features of the nanocomposite processing technique. (Reproduced with permission from [125] Zhu et al, Jpn J Appl Phys 2002; 41: 6969, Copyright The Physical Society of Japan 2002).

PZT ceramics have been reported to be prepared from the nanocomposite precursor at much lower temperature than that required by the conventional ceramic process [125]. Fig. 16 shows representative SEM images of PZT ceramics derived from nanocomposite precursors. Together with measured density results, it is found that fully dense PZT ceramics can be achieved after sintering at a temperature of as low as 800ºC. Another interesting feature is that no obvious grain growth is observed in the samples sintered at temperatures from 850ºC to 1000ºC, which means that the grain growth was stopped at ~850ºC. The absence of grain growth at high temperatures has been attributed to the uniform size distribution and non-

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agglomeration of the nanocomposite precursors. This feature is particularly of interest if this kind of precursor is used to fabricate multilayer structured devices in which small thickness of single layer is required.

Figure 16. SEM images of the PZT ceramics derived from nanocomposite precursors sintered at different temperatures: (a) 750ºC, (b) 800ºC, (c) 850ºC and (d) 950ºC for 2 h. (Reproduced with permission from [125] Zhu et al, Jpn J Appl Phys 2002; 41: 6969, Copyright The Physical Society of Japan 2002).

Figure 17. SEM images of thick film prepared from nanocomposite precursor: (a) and (b) cross-section and (c) surface. (Reproduced with permission from [126], Wang et al, Mater Sci Eng B 2003; 99: 56, Copyright Elsevier 2003).

Another significant progress is to use nanocomposite precursor to prepare thick films for piezoelectric transducers [129-132]. Thick films with thickness of up to 25 µm have been readily prepared via multilayer deposition with one layer being ~2 µm, which is nearly 50100 times the typical sol-gel solution process. An example is shown in Fig. 17 [126]. The film is very uniform in thickness, with a dense microstructure and narrow distribution of grain sizes (~100 nm). A well-developed P-E hysteresis loop is observed for the thick films, as shown in Fig. 18 [126]. Compared with the thick films prepared using a precursor made of

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microsized PZT powder, the nanocomposite film has a higher dielectric constant and a lower loss tangent, which has been attributed to the factor that the latter has a dense microstructure than the former. Such kind of composite thick films have been successfully combined with silicon technology to fabricate MEMS piezoelectric devices. An example is shown in Fig. 19.

Figure 18. Representative P-E hysteresis loop of a 10-µm PZT thick film. (Reproduced with permission from [126], Wang et al, Mater Sci Eng B 2003; 99: 56, Copyright Elsevier 2003).

Figure 19. Continued on next page.

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Figure 19. Structure of a membrane-type p-MUT fabricated on a silicon wafer: (a) schematic view, (b) top profile image of a fabricated element, and (c) detailed cross section image of the membrane. (Reproduced with permission from [130], Wang et al, Sens Actuat A 2006; 130-131: 485, Copyright Elsevier 2006).

5.4. Nano-microcomposites Similar approach has been also applied to barium strontium titanate (Ba0.6Sr0.4TiO3, or BST) [133]. BST thick films and dense ceramics were obtained from a sol-gel solution mixed with a combination of microsized and nanosized BST powders. Therefore, there is a slight difference between the two kinds of composites. The use of microsized powder is to compensate the shortage of nanosized components. A great reduction in annealing/sintering temperature has been achieved in such a process. Microsized BST powders were prepared via the conventional ceramic process, while nanosized powders (20-60 nm) were made from the microsized ones using a high-energy ball milling with tungsten carbide media. BST sol-gel solutions were prepared separately. Mixtures of the solutions and the microsized and nanosized powders were made by a normal milling. The mixture can be directly used to deposit thick films via spin-coating. By a simple calination at relatively low temperatures, the powders can be used to prepare films via screen-printing and bulk materials through the ceramic processing. It has been reported that the films derived from microsized powder has a large number of cracks, which is attributed to the presence of interstices among the BST particles [133]. With the use of nanosized BST powder, dense and homogeneous films are obtained. Using different electrodes (Ag or Pd-Ag), the annealing temperatures for the BST thick films can be varied from 750ºC to 1200ºC. Generally, the electrical properties of BST films will increase with increasing annealing temperature. Therefore, a tradeoff should be considered between performance and cost. Nevertheless, this technique provides us with flexibilities in fabrications of BST thick films. The electrical properties can be further improved by using screen-printing method to prepare the BST films. At the same time, BST ceramics have been obtained from the precursors at relatively lower temperatures than that required by the conventional ceramic processing.

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6. Summaries High-energy mechanochemical activation has been shown to be a versatile and reliable technique to synthesize various ferroelectric ceramic powders, especially those containing lead (Pb). Direct production of ferroelectric materials via the high-energy mechanochemical technique leads to nanosized powders, because the activation not only is able to trigger chemical reactions but also suppresses the grain growth. Various ferroelectric and antiferroelectric ceramics derived from the synthesized nanosized powders have relatively lower sintering temperatures and possess better or similar dielectric, ferroelectric, pyroelectric and piezoelectric properties. Ferroelectric ceramics, with promising microstructural and electrical properties, can be fabricated not only from the synthesized powders, but also from partially reacted or even unreacted mixtures activated by a high-energy ball milling. The activation of high-energy milling to synthesize relaxor ferroelectrics reveals some interesting aspects. For example, polycrystalline lead zinc niobate (PbZn1/3Zn2/3O3 or PZN), a low-temperature stable and high-temperature unstable phase, can never be synthesized via the conventional solid-state reaction process. However, by using a high-energy shaker mill, single phase PZN has been directly obtained from the oxide mixture, although PZN ceramics cannot be prepared from the synthesized powder. According to the phase diagram of PZT, it is confirmed that that the mechanochemical process is a low-temperature process. It is also found that single phase lead iron niobate (PbFe1/2Nb1/2O3 or PFN) can be formed from oxide mixture via a mechanical activation, but high performance PFN ceramics can only be derived from the columbite precursor of PbO and FeNbO4. Single phase lead iron tungstate (PbFe1/2W1/2O3 or PFW) can be synthesized neither from the oxide mixture of PbO, Fe2O3 and WO3, nor from the mixture of PbO and Fe2WO6. It is only possible to produce PFW from the combination of Pb3Fe2O6 and WO3. PFW is also available when 0.4-mol-PFW was used as seeds. Mechanochemical activation can be used to control the order-disorder status in some relaxors. Multilayer structured bismuth-containing ferroelectrics, belonging to the Arivillius family, have been either directly synthesized from oxide/carbonate precursors via high-energy milling techniques, or derived from amorphized precursor produced by post thermal annealing at relatively low temperatures, compared to their counterparts prepared via the conventional ceramic process. The mechanism that governs the phase formation of ferroelectric materials via highenergy mechanical milling process deserves further studies to clarify. The refinement of the precursor constituents, the creation of defects and imperfections, localized temperatures or pressures are all contributors the increase in reactivity of the starting materials. Significant progress has been made in the synthesis of ferroelectric materials using the high-energy mechanochemical process. Future studies should be focused on (i) systematically investigating the effect of various processing parameters on the ferroelectric phase formation, microstructure of the synthesized powders, as well as the microstructural and electrical properties of the final ceramic products, (ii) developing mechanisms/models to account for the phase evolution of ferroelectric compounds as a result of high-energy mechnaochemical milling. Examples have demonstrated that nanosized ferroelectric powders, produced by highenergy mechanochemical synthesis or activation, can find interesting applications in

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fabrications of thick films, nanocomposites and bulk ceramics. The hybrid nanocomposite processing technique has well combined the advantages of both nanosized materials and solgel process. Further works are necessary to explore more applications the materials synthesized via the new technique.

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[56] Kong L B, Ma J, Zhu W, Tan O K. Rapid formation of lead magnesium niobate-based ferroelectric ceramics via a high-energy ball milling process. Mater Res Bull 2002; 37: 459-65. [57] Algueró M, Alemany C, Jiménez B, Holc J, Kosec M, Pardo L. Piezoelectric PMN-PT ceramics from mechanochemically activated precursors. J Eur Ceram Soc 2004; 24: 937-40. [58] Algueró M, Moure A, Pardo L, Holc J, Kosec M. Processing by mechanosynthesis and properties of piezoelectric Pb(Mg1/3Nb2/3)O3-PbTiO3 with different compositions. Acta Mater 2006; 54: 501-11. [59] Kong L B, Ma J, Huang H, Zhang R F. (1-x)PZN-xBT ceramics derived from mechanochemically synthesized powders. Mater Res Bull 2002; 37: 1085-92. [60] Kong L B, Ma J, Huang H, Zhang R F. Lead zinc niobate (PZN)-barium titanate (BT) ceramics from mechanochemically synthesized powders. Mater Res Bull 2002; 37: 2491-8. [61] Tan Y L, Xue J M, Wang J. Stablization of perovskite phase and dielectric properties of 0.95PZN-0.05BT derived from mechanical activation. J Alloys Comp 2000; 297: 928. [62] Xue J M, Tan Y L, Wan D M, Wang J. Synthesizing 0.9PZN-0.1BT by mechanically activating mixed oxides. Solid State Ionics 1999; 120: 183-8. [63] Alguero M, Hungria T, Amorin H, Ricote J, Galy J, Castro A. Relaxor bahavior, polarization buildup, and switching in nanostrctured 0.92PbZn1/3Nb2/3O3-0.08PbTiO3 cermics. Nanostruct Ceram 2007; 3 (11): 1906-11. [64] Choudhary R N P, Pradhan D K, Tirado C M, Bonilla G E, Katiyar R S. Relaxor characteristics of Pb(Fe2/3W1/3)O3-BiFeO3 solid solution prepared by mechanosynthesis route. J Appl Phys 2006; 100: 084105-1-8. [65] Varshney D, Choudhary R N P, Katiyar R S. Low frequency dielectric response of mechanosynthesized (Pb0.90Ba0.1)(Fe0.50Nb0.50)O3 nanoceramics. Appl Phys Lett 2006; 89: 172901-1-3. [66] Shinohara S, Baek J G, Isobe T, Senna Mamoru. Synthesis of phase-pure Pb(ZnxMg1-x)1/3Nb2/3O3 up to x=0.7 from a single mixture via a soft-mechanochemical route. J Am Ceram Soc 2000; 83 (12): 3208-10. [67] Ang S K, Wang J, Xue J M. Phase stability and dielectric properties of (1-x) PFW+xPZN derived from mechanical activation. Solid State Ionics 2000; 127: 285-93. [68] Wan D M, Xue J M, Wang J. Synthesis of single phase 0.9Pb[(Zn0.6Mg0.4)1/3Nb2/3O3]0.1PbTiO3 by mechanically activating mixed oxides. Acta Mater 1999; 47 (7): 2283-91. [69] Wan D M, Xue J M, Wang J. Nanocrystalline 0.54PZN-0.36PMN-0.1PT of perovskite structure by mechanical activation. Mater Sci Eng A 2000; 286; 96-100. [70] Wan D M, Xue J M, Wang J. Mechanochemical synthesis of 0.9[0.6Pb(Zn1/3Nb2/3)O30.4Pb(Mg1/3Nb2/3)O3]-0.1PbTiO3. J Am Ceram Soc 2000; 83 (1) 53-9. [71] Ang S K, Wang J, Xue J M. Mechanical activation and dielectric properties of 0.48PFN-0.36PFW-0.16PZN from mixed oxides. J Alloys Comp 2000; 311: 181-7. [72] Gao X S, Xue J M, Wang J, Yu T, Shen Z X. B-site disording in Pb(Sc1/2Ta1/2)O3 by mechanical activation. Appl Phys Lett 2003; 82 (26): 4773-5.

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[73] Gao X S, Xue J M, Yu T, Shen Z X, Wang J. Mechanical activation-induced B site order-disorder transition in perovskite Pb(Mg1/3Nb2/3)O3-Pb(Mg1/2W1/2)O3. Mater Chem Phys 2002; 75: 211-5. [74] Gao X S, Xue J M, Yu T, Shen Z X, Wang J. B-site order-disorder transition in Pb(Mg1/3Nb2/3)O3-Pb(Mg1/2W1/2)O3 triggered by mechanical activation. J Am Ceram Soc 2002; 85 (4): 833-8. [75] Xue J M, Wang J, Wang D M. Nanosized barium titnate powder by mechanical activation. J Am Ceram Soc 2000; 83 (1) 232-4. [76] Abe O, Suzuki Y. Mechanochemically assisted preparation of BaTiO3 powder. Mater Sci Forum 1996; 225: 563-8. [77] Kong L B, Ma J, Huang H, Zhang R F, Que W X. Barium titanate derived from mechanochemically activated powders. J Alloys Comp 2002; 337: 226-30. [78] Berbenni V, Marini A, Bruni G. Effect of mechanical milling on solid state formation of BaTiO3 from BaCO3-TiO2 (rutile) mixtures. Thermochimica Acta 2001; 374: 151-8. [79] Van Hal H A M, Groen W A, Maassen S, Keur WC. Mechanochemical synthesis of BaTiO3, Bi0.5Na0.5TiO3 and Ba2NaNb5O15 dielectric ceramics. J Eur Ceram Soc 2001; 21: 1689-92. [80] Hotta Y, Tsunekawa K, Isobe T, Sato K, Watari K. Synthesis of BaTiO3 powders by a ball milling-assisted hydrothermal reaction. Mater Sci Eng A 2008; 475: 12-6. [81] Hotta Y, Tsunekawa K, Duran C, Sato Kimiysu, Nagaoka T, Watari K. Lowtemperature sintering of BaTiO3 powders prepared by a hydrothermal process with ball milling system. Mater Sci Eng A 2008; 475: 57-61. [82] Senna M. Smart mechanochemistry-charge transfer control for tailored solid-state reaction under external energy. J Alloys Comp 2007; 434-435: 768-72. [83] Hungria T, Castro A. Synthesis and structural characterization of the new solid solution Ba2-xSrxTiO4 (0 ≤ x ≤1): Effect of the method of synthesis on the polymorphic phase isolated. J Alloys Comp 2007; 436: 266-71. [84] Brankovic G, Vukotic V, Brankovic Z, Varela J A. Investigation on possibility of mechanochemical synthesis of CaTiO3 from different precursors. J Eur Ceram Soc 2007; 27: 729-32. [85] Mi G, Murakami Y, Shindo D, Saito F. Mechanochemical synthesis of CaTiO3 from a CaO-TiO2 mixture and its HR-TEM observation. Powder Tech 1999; 105: 162-6. [86] Kong L B, Ma J, Zhu W, Tan O K. Preparation of Bi4Ti3O12 ceramics via a highenergy ball milling process. Mater Lett 2001; 51: 108-14. [87] Alguero M, Ferrer P, Vila E, Iglesias J E, Castro A. Bi4Ti3O12 cermics from powders prepared by alternative routes: wet no-coprecipitation chemistry and mechanochemicl activation. J Am Ceram Soc 2006; 89 (11): 3340-7. [88] Ng S H, Xue J M, Wang J. Bismuth titanate from mechanical activation of a chemically coprecipitated precursor. J Am Ceram Soc 2002; 85 (11): 2660-5. [89] Zdujic M, Poleti D, Jovalekic C, Karanovic L. The evolution of structure induced by intensive milling in the system 2Bi2O3·3TiO2. J Non-Cryst Solids. 2006; 352: 3058-68. [90] Shantha K, Varma K B R. Preparation and characterization of nanocrystalline powders of bismuth vanadate. Mater Sci Eng B 1999; 56: 66-75. [91] Shantha K, Subbanna G N, Varma K B R. Mechanically activated synthesis of nanocrystalline powders of ferroelectric bismuth vanadate. J Solid State Chem 1999; 142: 41-7.

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[92] Shantha K, Varma K B R. Characterization of fine-grained bismuth vanadate ceramics obtained using nanosized powders. J Am Ceram Soc 2000; 83 (5): 1122-8. [93] Ricote J, Pardo L, Castro A, Millán P. Study of the process of mechanochemical activation to obtain Aurivillius oxides with n=1. J Solid State Chem 2001; 160: 54-61. [94] Castro A, Millán P, Ricote J, Pardo L. Room temperature stabilization of γ-Bi2VO5.5 and synthesis of the new fluorite phase f-Bi2VO5 by a mechanochemical activation. J Mater Chem 2000; 10: 767-71. [95] Castro A, Bégué P, Jiménez B, Ricote J, Jiménez R, Galy J. New Bi2Mo1-xWxO6 solid solution: mechanosynthesis, structural study, and ferroelectric properties of the x=0.75 member. Chem Mater 2003; 15: 3395-401. [96] Sim M H, Xue J M, Wang J. Layer structured calcium bismuth titanate by mechanical activation. Matter Lett 2004; 58: 2032-6. [97] Ng S H, Xue J M, Wang J. High-temperature piezoelectric strontium bismuth titanate from mechanical activation of mixed oxides. Mater Chem Phys 2002; 75: 131-5. [98] Ferrer P, Iglesias J E, Castro A. Synthesis of the Aurivillius phase SrBi4Ti4O15 by a mechanochemical activation route. Chem Mater 2004; 16: 1323-9. [99] Castro A, Millán P, Pardo L, Jiménez B. Synthesis and sintering improvement of Aurivillius type structure ferroelectric ceramics by mechanochemical activation. J Mater Chem 1999; 9: 1313-7. [100] Ricote J, Pardo L, Moure A, Castro A, Millán P, Chateigner D. Microcharacterization of grain-oriented ceramics based on Bi3TiNbO9 from mechanochemically activated precursors. J Eur Ceram Soc 2001; 21: 1403-7. [101] Moure A, Pardo L, Alemany C, Millán P, Castor A. Piezoelectric ceramics based on Bi3TiNbO9 from mechanochemically activated precursors. J Eur Ceram Soc 2001; 21: 1399-402. [102] Jiménez B, Castro A, Pardo L, Millán P, Jiménez R. Electric and ferro-piezoelectric properties of (SBN)1-x(BTN)x ceramics obtained from amorphous precursors. J Phys Chem Solids 2001; 62: 951-8. [103] Pardo L, Castro A, Millán P, Alemany C, Jiménez R, Jiménez B. (Bi3TiNbO9)x(SrBi2Nb2O9)1-x Aurimillius type structure piezoelectric ceramics obtained from mechanochemically activated oxides. Acta Mater 2000; 48: 2421-8. [104] Moure A, Castro A, Pardo L. Improvement by recrystallization of Aurivillius-type structure piezoceramics from mechanically activated precursors. Acta Mater 2004; 52: 945-57. [105] Moure A, Alemany C, Pardo L. Electromechanical properties of SBN/BTN Aurivilliustpye ceramics up to the transition temperature. J Eur Ceram Soc 2004; 24: 1687-91. [106] de Figueiredo R S, Messai A, Hernandes A C, Sombra A S B. Piezoelectric lithium niobate by mechanical alloying. J Mater Sci Lett 1998; 17: 449-51. [107] Castro A, Jiménez B, Hungria T, Moure a, Pardo L. Sodium niobate ceramics prepared by mechanical activation assisted methods. J Eur Ceram Soc 2004; 24: 941-5. [108] Hungria T, Pardo L, Moure A, Castro A. Effect of mechanochemical activation on the synthesis of NaNbO3 and precessing of environmentally friendly piezoceramics. J Alloys Comp 2005; 395: 166-73. [109] Rojac T, Kosec M, Malic B, Holc J. Mechanochemical synthesis of NaNbO3. Mater Res Bull 2005; 40 341-5.

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[110] Rojac T, Masson O, Guinebretiere, Kosec M, Malic B, Holc J. A study of the mechanochemical synthesis of NaNbO3. J Eur Ceram Soc 2007; 27: 2265-71. [111] Takacs L. Multiple combustion induced by ball milling. Appl Phys Lett 1996; 69: 4368. [112] Thadhani N N. Shock-induced and shock-assisted solid-state chemical reactions in powder mixtures. J Appl Phys 1994; 76: 2129-38. [113] Jang H M, Oh S H, Moon J H. Thermodynamic stability and mechanisms of Pb(Zn1/3Nb2/3)O3 formation and decomposition of perovskite prepared by the PbO flux method. J Am Ceram Soc 1992; 75: 82-8. [114] Wakiya N, Ishizawa N, Shinozake K, Mizutani N. Thermal stability of Pb(Zn1/3Nb2/3)O3 (PZN) and consideration of stabilization conditions of perovskite type compounds. Mater Res Bull 1995; 30: 1121-31. [115] Sundararaman D. Nanocrystalline state and solid state amorphization. Mater Sci Eng B 1995; 32: 307-13. [116] Fukunaga T, Ussumi M, Akatsuka H, Misawa M, Mizutani U. Structure of amporphous Se prepared by milling. J Non-Cryst Solids 1996; 205-207: 531-5. [117] Huang J Y, Yasuda H, Mori H. Deformation-induced amorphization in ball-milled silicon. Phil Mag Lett 1999; 79: 305-14. [118] Yang X Y, Wu Y K, Ye H Q. Localized amorphization in SiC induced by ball milling. J Mater Sci Lett 2001; 20: 1517-8. [119] Kong L B, Ma J, Zhu W, Tan O K. Highly enhanced sinterability of commercial PZT powders by high-energy ball milling. Mater Lett 2000; 46: 274-80. [120] Randall C A, Kim N, Kucera J P, Cao W, Shrout T R. Intrinsic and extrinsic size effects in fine-grained morphotropic-phase-boundary lead zirconate titanate ceramics, J Am Ceram Soc 1998; 81 (3): 677-88. [121] Pavlovic C V P, Nikolic M V, Pavlovic V B, Labus N, Zivkovic Lj, Stojanovic B D. Correlation between densification rate and microstructure evolution of mechanically activated BaTiO3. Ferroelectrics 2005; 319:75-85. [122] Tusseau-Nenez S, Ganne J P, Maglione M, Morell A, Niepce J C, Pate M. BST ceramics: effect of attrition milling on dielectric properties. J Eur Ceram Soc 2004; 24:3003-11. [123] Kosec M, Holc J, Duscer D, Drnovšek S. Pb(Mg1/3Nb2/3)O3-PbTiO3 thick films from mechanochemically synthesized powder. J Eur Ceram Soc 2007; 27: 3775-8. [124] Wang Z, Zhao C, Zhu W, Tan O K, Liu W, Yao X. Processing and characterization of Pb(Zr,Ti)O3 thick films on platinum-coated silicon substrate derived from sol-gel depositon. Mater Chem Phy 2002; 75: 71-5. [125] Zhu W, Wang Z, Zhao C, Tan O K, Hng H H. Low temperature processing of nanocrystalline lead irconate titanate (PZT) thick films and ceramics by a modified solgel route. Jpn J Appl Phys 2002; 41: 6969-75. [126] Wang Z, Zhu W, Zhao C, Tan O K. Dense PZT thick films derived from sol-gel based nanocomposite process. Mater Sci Eng B 2003; 99: 56-62. [127] Zhao C, Wang Z, Zhu W, Tan O K, Hng H H. Microstructure and properties of PZT53/47 thick films derived from sols with submicron-sized PZT particle. Ceram Int 2004; 30: 1925-7. [128] Wang Z, Zhu W, Chao C, Zhao C, Chen X F. Characterization of composite piezoelectric thick film for MEMS application. Surf Coat Tech 2005; 198: 384-8.

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In: Nanotechnology… Editors: C.J. Dixon and O.W. Curtines, pp. 371-407

ISBN: 978-1-60692-162-3 © 2010 Nova Science Publishers, Inc.

Chapter 11

PREPARATION AND CHARACTERIZATION OF MONOATOMIC CARBON CHAINS: UNRAVELING, FIELD ION MICROSCOPY, AND FIELD EMISSION Igor M. Mikhailovskij National Scientific Center “Kharkov Institute of Physics and Technology”, Academicheskaja, 1, 61108 Kharkov, Ukraine

Abstract Linear forms of carbon are important in a wide variety of application, ranging from highly conducting interconnects to field emission materials. By methods of field ion microscopy (FIM) and mass-spectrometry, it was revealed the presence of linear carbon chains at the surface of carbon fibers after high-voltage treatment. We present in this chapter a brief review of these research emphasizing recent developments. The carbon chains attached to the specimen tips can be produced in situ in a field ion microscope using low-temperature pulsed evaporation by electric fields of the order of 1011 V/m. Atomic C-chains are produced during the high-field unraveling of nanofibers. The experimental procedures used in FIM carbon chains studies are reviewed and the results in relation to the atomistics of unraveling processes are discussed. Molecular dynamics simulations and high resolution FIM experiments are performed to assess the evaporation of atomic chains under high-field conditions. Carbon exhibits a very rich dynamics of bond-breaking that allows transformation from graphenes to atomic chains. High-field experiments, theories leading to carbon chain formation, and methods to extract quantitative information on a variety of chain-surface interactions are described in detail. Isolated atomic carbon chains can be obtained at different temperatures, pulling speeds and forces. Current versus voltage field electron characteristics of monoatomic carbon wires were investigated. These results lend strong support to the conjecture of Smalley that linear carbon chains may provide the ultimate atomic-scale field emitters.

1. Introduction The continuing miniaturization of nanoelectronics raises the prospect of nanometre-scale devices with mechanical and electrical properties that are qualitatively different from those at larger dimensions. At the moment one of the major challenges of the research on molecular

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electronics is in connecting those nanodevices. Monoatomic chains and nanowires are viewed as having great potential as interconnects in molecular electronics applications. The fabrication of nanodevices using one-dimensional nanomaterials involves a new level of complexity that must be accompanied by an atomic-level understanding of involved processes. To develop realistic models of chain formation, it is necessary to determine quantitatively the topological and energetic peculiarities of chains. Most early studies have focused on metallic elements due to their capability to form nanowires by mechanical stretching. Interest in this field has been risen by electron microscopy observation of monoatomic gold chains with atomic resolution and by confirmation of the correlation between quantum conductance and their atomic structure [1]. Triggered by these studies on gold, theoretical investigations have been expanded to atomic chains of other elements, including nonmetals. Linear atomic chains of nonmetals are of great interest as either free-standing objects, or components of nanoelectronic device. Analysis of electronic structures of atomic chains reveals fundamentally and technologically interesting phenomena, such as structural instabilities [2]. Their topology and electronic structure may be controlled by charge doping gendered by the presence of an electrical current or the chemical enviroment [3,4]. Carbon atomic chains of various sizes have been registered in stellar atmospheres, in comet tails, and in the interstellar space. The carbon chain species, which had been discovered in the Universe, might be related to the carbon dust blown out from the shell of carbon-rich red giant stars [5,6]. As far as terrestrial studies of carbon chains, the major motives are due to the unusual electrical, optical and mechanical properties that result from the small sizes of these onedimensional materials. The understanding and functionalization of carbon atomic chains has become a topic of great interest to the nanomaterials science and technology communities. Monoatomic wires can sustain very large current densities, proving that the electron transport is ballistic, and that most of the power is dissipated in the supporting electrodes far away from the contact. This makes the atomic chains suitable for investigation as conductors in molecular electronic circuits and it is crucial for the design and fabrication of the nanodevices. The experimental studies and technological applications require advanced techniques of production of carbon atomic chains. Those include chemical synthesis [7], electric arc discharge between graphite electrodes submerged in a chemically inert liquid [8], electric-field-induced unfolding [9], and by mechanical pulling [10-12] of carbon nanotubes. In their paper, Rinzler et al. [9] showed that field electron emission from individual carbon nanotubes (CNTs) has been enhanced when the nanotube apexes are opened by laser evaporation. After considering several alternative explanations, authors concluded that the sharp emitting element is a linear chain of 10 to 100 carbon atoms pulled out from the open CNT by the high electric field in a manner that resembles unraveling the sleeve of a sweater. In these experiments, carbon chains were pulled out with no net decrease in effective bond order. Free-standing linear carbon chains pulled out from the open CNT are held taut under the action of the high electric field. The carbon chain at the point of attachment to the CNT is bonded to the delocalized π-orbitals of the CNT outer shell, this carbon system is both mechanically and electrically well coupled. Free-standing multi-wall nanotubes have been demonstrated to emit field electron current of up to several microamperes. Based on obtained results, Rinzler et al. conjecture that such a large current originate from “atomically sharp” structures and these structures may provide the ultimate atomic-scale field emitters. There was no information available about the detailed disintegration mechanism or about the

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threshold field strength needed for a CNT to unfold. The viability of the exotic unraveling mechanism was verified by Yacobson et al. [10] in molecular-dynamics simulations of CNT deformation, although used many-body interatomic potentials may not be fully reliable for one-dimensional structures. The plausibility of occasional formation of carbon atomic chains during plastic deformation of CNTs was revealed in these computer experiments. Despite its probabilistic character of formation, once a carbon chain is pulled it remains very stable. Various theoretical approaches used to the unraveling model were discussed, proceeding from simple pair-potential calculations to sophisticated, first-principles calculations. It was shown that that CNTs undergo a process of necking and some of the tubes deform into a single chain configuration, spanning the nanotube fragments before fracture [11,12]. Yuzvinsky at al., [13] have demonstrated a method to controllably shrink CNTs to any diameter. The method uses high-energy electrons from a transmission electron microscope to knock carbon atoms out of the nanotube while monitoring the changing CNT geometry in real time. The high temperatures attained through resistive heating assist a self-healing annealing and electromigration, resulting in a smaller CNT with perfect atomic structure. As the diameter of a multiwalled CNT decreases below 1 nm, the inner wall breaks and endcaps form on both ends of the CNT connected by a very thin bridge. The conductance of the nanotube device is determined mostly by the size of its narrowest cross-section and that, during the modification of CNT, all important structural and topographic transformations are localized to the neck region. The thinnest section of the nanobridge cannot be imaged in highresolution TEM but it is likely that the bridge had a carbon-chain-like structure. This conclusion was confirmed by a subtle conductance measurement. Authors found that the total conductance of the CNT device with a narrow neck is linearly proportional to its cross sectional area. During the last stage of the shrinking process, conductance stays in a limited range of values corresponding to one atom in cross-section. In this regime, the CNT device exhibits negative differential resistance at both negative and positive bias of a fraction of a volt. Theoretical investigations showed that such a behavior of the bridged nanotube junction at the end of the shrinking process is typical of the monoatomic chain regime [14]. The possibility of creating carbon chains in high electric fields was confirmed by the theoretical investigation using ab initio density functional formalism [15,16]. Linear carbon chains with more than ten atoms are thermodynamically less favorable than ring isomers, except for a C6 chain, which has degenerate linear and ring isomers in zero electric field. However, the linear chains are stabilized by the application of an external electric field relative to rings due to the preferential field-induced polarization for the linear chains [17]. Due to the polarization effect, the high external electric field can assist the extension of an atomic wire unraveled from the open edge of a graphene layer. Even though the unraveling mechanism is very appealing and promising, it has not been conformed by independent atomic-scale experiments till recently. Combined field ion and electron microscopy and mass-spectrometry have just now revealed the presence of linear carbon chains at the carbon surface treated by electric fields of the order of 1011 V/m [18,19]. The linear chains attached to the of sharpened carbon fibers consisting of more than ten atoms can be produced in situ in a field emission microscope during low-temperature field evaporation. The key findings of these works are the mechanical stability of the chains emanating from graphite in ultra-high electric fields and that linear carbon chains may provide the ultimate atomic-scale field emitters. This chapter ties together several aspects of

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recent research on carbon atomic chains and techniques used, focusing on a high-field technology, high-resolution microscopy, field emission, and molecular dynamics simulation.

2. Techniques and Instrumentations Although recent advances in experimental and theoretical methods yield detailed information on structure of nanomaterials, an essential part of our current atomic-scale understanding of atomic processes at surfaces of nanomaterials in a high electric field has come from investigations with the field ion microscope (FIM), an instrument invited in the early 1950’s. The unique attributes of this technique allow one to view an individual atom and track its motion as it moves across a well-defined specimen surface. A comprehensive review of the technique of field ion microscopy is beyond the scope of this book. For more information on the basic principles of field ion microscopy, the reader is referred to several review articles and books [20-22]. Only a brief overview of field ion microscopy is given here to provide a basis for the succeeding discussion of atomic chain studies. The FIM presents images of the protruding atoms at the apex of a sharp needle, chemically or electrochemically polished to a sharp point. The needle-shaped specimen is attached to a cryostat and a high voltage supply. FIM specimens are similar to the tips in scanning tunneling microscopy [20]. However, instead of using the tip to scan the specimen surface, the tip itself is studied in the FIM. An image of the atoms at the tip apex is obtained in the FIM after being evacuated to ultra-high vacuum by applying a high positive voltage in the presence of an image gas such as neon or helium. These gases were chosen because of their lack of chemical reactivity and their ability to be purified. A high voltage applied to the sharply pointed sample produces an extremely high and inhomogeneous electric field, of the order of several volts per Å. The voltage required is proportional to the end radius the tip and is typically 2-30 kV. Image gas atoms are polarized by the inhomogeneous electric field and draw towards the apex of the specimen. Gas atoms will hop around the specimen apex until they either escape to the free space or are field ionized a few Å above the surface and accelerates them toward an imaging detector in the radial direction of the tip hemisphere. Image contrast results from the field strength being higher at the position of protruding surface atoms. The magnification of the field ion microscope is dependent on the tip-todetector distance and typically is of the order of a million times. The resolution is a function of several factors including the sample temperature. At cryogenic temperatures (77 K or below) the resolution may approach of 2 Å and it is sufficient to identify individual atoms. The precision with which the relative positions of individual atoms can be measured is better then the nominal resolution and approaches 0.3 Å. A better resolution (~0.1 Å) may be achieved normal to the specimen surface using the indirect magnification phenomenon [20]. If the field strength is raised to a sufficiently high value, the most prominent surface atoms are themselves ionized and removed by the electric field. As a result, the clean uniform hemispherical end-form is produced. Such a cleaning process is unique to field ion microscopy and is known as field evaporation. In the region of the low-index planes, field evaporation removes atoms from the edge of the plane inwards, one atomic layer at the time. For stable operation of the FIM, it is necessary that the field strength required for evaporation of the specimen should exceed that required for field ionization.

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Specimen tips having an initial radius of curvature at the tip in the range 20-50 nm were prepared from carbon polyacrylonitrile-based fibers by the electrochemical etching in a 3N KOH solution. Carbon fibers, which have a cross-sectional size of 7-10 μm and an average strength of 3-5 GPa, are composed of nanofibrils, the average size of which varies between 50 and 200 nm. In addition to optical control the specimens during the course of electrochemical polishing, the tip topographies were examined using the transmission electron microscope (TEM) employing a goniometric tilt device and high-resolution scanning electron microscope SEM Germini 1540 (FIB) with the resolution of 0.5 and 0.7 nm, accordingly. Experiments were made in a field ion microscope (FIM) in which the samples were cooled to 4.2, 21 and 77 K [23,24]. The microscope was evacuated by cryogenic pumps to a residual gas pressure of about 10-7 Pa. Field ion and electron images were normally obtained in a voltage range 1-22 kV. Field ion images were formed by using helium gas under a pressure of 10-4–10-3 Pa. After sample mounting in the microscope, the surface was polished by low temperature field evaporation [21] until the formation of an atomically smooth hemispherical tip with a radius of curvature in the interval 30–70 nm. Field evaporation spectra of needle-shaped samples were determined by using the mass-spectroscopic add-on device. Evaporation of ions subjected to mass-analysis was implemented by a high-voltage pulse generator with the front steepness of 20 ns and pulse duration of 10 ns at a level of 0.8 the amplitude. The voltage pulse amplitude was varied from 0.5 to 7 kV. A total voltage of 122 kV was applied between the specimen tip and the entrance aperture of a straight drift tube 1.50 m long. The signals generated by individual ions or ionized clusters of atoms after the flight through the drift tube were detected by the chevron channel plate multiplier, and fed to the input of the timer. Field evaporation occurred under the action of the total voltage (a constant voltage required for the formation of the field ion image and a pulse voltage). Field-evaporation spectra in AP FIM are usually analyzed from the number of individual ion signals detected for different ion species. Since the time digitizer used with the chevron channel plate cannot discriminate signals according to the output pulse height, the ions with the same mass-to-charge ratio, which arrive simultaneously at the detector, are counted as only one signal. This problem is generally alleviated by operating the atom-probe at a very low counting rate. In this regime, the probability of pulse overlap becomes small, and the ion spectra derived should be accurate. But a preliminary study of the field evaporation of carbon specimens showed that this process is very irregular at low temperatures. A miss-counting of the simultaneously arriving ion signals cannot be eliminated even if the average evaporation rate of the carbon specimens is reduced below 0.01 atom/pulse. Usual time measuring devices were not convenient in this case, and the flight times of the ions were measured with an oscilloscope. A high voltage pulse removed the surface atoms in a form of singly or multiplycharged ion clusters and activated a time sweep oscillator. The signal voltage increased proportionally to the number of the ions or ionized clusters with the same mass-to-charge ratio, therefore the field evaporation spectra could be derived from the signal intensities measured for various ion species. The field F0 at the apex of a field ion specimen is reduced from the equivalent value at the surface of a free conducting sphere by the presence of the shank and can be expressed as

F0 =

V0 k f r0

(1)

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where V0 is the applied potential, r0 is the apex radius and kf (the geometric field factor) is a numerical constant which depends on the taper angle of specimen, but has an approximate value of 5. A usual for this method accuracy is ± 15%. More precise electric field calibration is based on the comparison with the best image field and voltage in the FIM operated with helium, hydrogen or nitrogen. The best image conditions are correspondent to the sharpest ion patterns with the best spatial resolution and contrast. The best image fields for He, H2 and N2 within ± 5% accuracy are 4.4·1010, 2.2·1010 and 1.7·1010 V/m, respectively [21]. Some control measurements of the field strength was made by the conventional method based on the validity of the Fowler-Nordheim (FN) equation for field electron emission with a usual for this method accuracy of ± 15%. Field ion microscopic images of carbon samples were analyzed by using the geometrical method of computer simulation of the images [20,21]. In the approximation of the model of thin envelope, it was assumed that the contribution to the formation of the ion-microscopic image comes from atoms located in a surface layer of a certain finite thickness.

3. Cluster Formation during Field Evaporation The potential technological benefits to be gained from a fundamental understanding of cluster formation have stimulated considerable research in recent years, both experimental and theoretical. There are plenty of subjects of scientific interest, such as binding energies of atoms in clusters, critical and magic numbers, ionization energies, topology of clusters, electronic energy levels, ets. It is usually guessed that carbon chains are produced during the vapor cooling process [25, 26]. The evaporation of individual atoms and their aggregation into clusters during low-temperature field evaporation are among the most fundamental atomic interactions. Several aspects in carbon cluster science can be investigated using lowtemperature pulsed-voltage field evaporation. In pulse field evaporation of carbon, clusters produced are ever charged and are accelerated at once in a high electric field and so after the cluster formation it no longer have any chance to interact with other clusters and atoms. Therefore, the abundance of cluster mass represents how carbon clusters are formed at the surface, the binding mechanism of clusters, and the absolute stability of the carbon species. During field evaporation of metals, ions formed are almost always monoatomic ions [21,22]. Carbon cluster ions were found in low-temperature [27] and pulsed-laser-stimulated [28] field evaporation. Mass spectra of graphite tips were obtained with a pulsed-laser timeof-flight atom probe with subnanosecond time resolution. When laser pulses were focused to the tip apex, a temperature rise of about 1000 K could be achieved. It was found that under an electric field strength of 20 - 30 V/nm and a temperature above 1000 K, carbon atoms were field evaporated generally as C++ and C+ ions but with about 10 – 20 % of cluster ions of various sizes and charge states. Cluster ions up to about 11 atoms were detected. The number of cluster ions decreases exponentially with increasing of cluster size. Clusters with an even number of atoms are less abundant than clusters with an odd number of atoms. Authors [28] established the existence of very large high-energy tails extending to about 350 V and it was supposed that some of carbon clusters were ejected out the specimen surface with a large initial kinetic energy. This result was explained by a strong photoexcitation effect in pulsedlaser stimulated field evaporation. Up to now there is neither a detailed explanation nor a comprehensive experimental study of this phenomenon. Analysis of the energy distribution in

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377

the time-of-flight spectra of low-temperature field evaporation of graphite [27] revealed that the registered charged clusters were not occurred right at the surface but were produced by dissociation within the acceleration region into several particles. The dissociation products are separately accelerated, until they reach the field-free region. As a result Cn clusters exhibit broadened energy distribution with both low and high energy tails. A mathematical treatment of obtained results showed that the carbon clusters were dissociated at a distance of 30 to 300 nm away from the specimen apex, which is reached within a travel time from 0.6 to 6 ps. The process of low-temperature field evaporation at the pulse voltage loading occurs sporadically with an anomaly large instant intensity of evaporation corresponding to explosive rate of about 1 – 10 m/s. Under an averaged electric field of about (7-8) ·109 V/nm carbon atoms are field evaporated mostly as Cn+ clusters with n = 2-6. A typical scan illustrating the cluster-wise field evaporation of the needle-shaped carbon nanofiber is shown in Figure 1. This time-of-flight mass spectrum is obtained at a low level of the signal voltage amplification, which precludes from the possibility to detect single ions. The height of steps at the oscillogram is proportional to the number of the clusters with the same mass-to-charge ratio. Taking into account this fact we showed that the maximum of carbon atoms registered by the detector is about of 103 atoms/pulse at the average evaporation rate of the carbon specimens of about 1 atom/pulse. Consequently, these results confirm our preliminary conclusion about significant irregularity of the field evaporation process for carbon specimens at low temperatures. This process was controlled by using the TEM with the resolution of 0.5 nm. It was shown that explosive field evaporation is accompanied by significant blunting and smoothing of the needle-shaped carbon specimens (Figure 2).

Figure 1. Field evaporation mass spectrum of carbon nanofibers at 77 K.

In the mass-spectroscopic mode, it was possible to determine the abundance of carbon clusters produced by field evaporation of graphite specimens at 77 K. Each set of massspectroscopic data contained about 104 ions which were collected from a probe hole of the area covering about ten atomic diameters at the specimen surface. Singly charged carbon clusters of up to 7 atoms are observed under pulse field evaporation of graphite fibers at cryogenic temperature (table 1). The number of clusters with odd carbon atoms is significantly larger than that with an even number of atoms. The intensity of C3+ peak in the mass spectra was prominently strong, indicating either a preferred formation or a high stability of the chains C3+. An anomaly abundance of the clusters C3 may be explained also by field-induced fragmentation of larger molecules. Really, for clusters with six or more atoms the dominant fragmentation channel is the loss of C3 to give Cn-3 clusters [29].

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Igor M. Mikhailovskij Table 1. Relative abundance of field evaporated carbon ions Ion Scaled abundance

C 0.84

C2 0.24

C3 1.00

C4 0.09

C5 0.58

C6 0.06

C7 0.08

An observed abrupt disappearance of the FIM image spots is readily understood as field evaporation of C3+ and other small charged radicals from the end part of the chain until this chain is so small that the electric field enhancement factor is no longer sufficient to produce field ionization of helium atoms. Experimental results and theoretical calculations ranging from sophisticated, firstprinciples approaches are now able to provide insight concerning the underlying physics involved in cluster nucleation and stability. The numbers of atoms in clusters produced by pulse field evaporation are much smaller as they are formed by vaporizing with a laser or an ion beam a graphite target [30]. The magic carbon cluster ions produced by desorption under a high electric field were also observed in field evaporation spectra of single-walled and multi-walled nanotubes in experiments of Hata et al. [31]. The spectral width is not due to a degradation of time-of-flight spectrum caused by neither in the initial kinetic energy of clusters nor by the space charge screening. The spacecharge effect is also negligible because the ion current density of our measurement amounts to less then 1% of the saturation limit caused by the space charge [32]. Since the wide spectral width is not caused by extrinsic effects, it has to come from disintegration of evaporated carbon clusters within the period of acceleration in the AP FIM.

4. Field Ion Microscopy of Carbon Chains As was shown in [18,19], graphite specimens are mechanically stable in the field-ion microscope and can withstand the evaporation field and imaging field strength of helium, which is about 4.5·109 V/nm. Field evaporation in the field ion microscope was accompanied by a dramatic blunting of the needle-shaped carbon specimens. SEM and TEM observations showed that the specimens having initial radii of curvature r0 in the range 20-50 nm were blunted up to r0 = 200-1000 nm in these conditions. Figures 2(a) and (b) show the electron shadowgraph and top-view SEM image of the carbon specimen formed by field evaporation

(a)

(b)

Figure 2. Electron microscope shadowgraph (a) and top-view SEM (b) images of a typical carbon tip after low-temperature field evaporation. (Reprinted with permission from [18]. © 2007, Institute of Physics.)

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Figure 3. FIM image of a carbon tip after low-temperature field evaporation at 6.60 kV.

during a steep (~ 1 V/s) increase of voltage V up to 6.60 kV. Examination of such micrographs shows that the apex part of specimens is more or less spherical without sharp edges and nanoprotrusions which would give rise to field enhancement and would show up by locally intensified field emission. Figure 3 shows the FIM image of the specimen formed by low-temperature field evaporation at 6.60 kV in ions of helium. As one can see in Figure 3 unlike metallic tips in high electric fields, hemispherical surface formed by field evaporation is microscopically rough. Its FIM image is the carbon characterized by an irregular distribution of image points. This type of FIM contrast is correspondent to the random character of the low-temperature field evaporation of carbon materials. During a field increase, the brightness of the image spots enhances. The FIM images reveal the process of long-range jumps of image spots at some stage in the voltage increase [19]. The main attribute of the FIM for the investigation of an atomic topography is its ability to reveal the location of atomic protrusions on a surface formed by field evaporation. Such protrusions are characterized by an enhancement in the local electric field. Due to the nature of the image formation process, these local field enhancements produce high-contrast image spots on a uniformly dark background. The observation of long jumps of bright FIM image spots during a high-field treatment points out that in this case a local field enhancement at bright image spots is not due to a sharp microprotrusion on the graphite tips. FIM images of carbon samples were analyzed by using Moore’s geometrical model of mathematical simulation [20,21]. Most features of FIM image contrast can be usually understood on the basis of this consideration. The main attribute of the FIM for the structure investigations is its ability to reveal the location of the isolated atom on a perfectly defined single-crystal terrace. Individual atoms typically protrude from the surface by a sufficient amount to cause an enhancement in the local electric field. Due to the nature of the image formation process, these local field enhancements produce high-contrast image spots on a uniformly dark background. In the approximation of the “thin shell” model, it was assumed that the atoms which give contribution to image spots are located in a surface layer of a certain finite thickness δM, typically 20-40 pm. The atoms lying at a greater depth will not field ionize enough gas atoms to form FIM image spots. It is important that the overall shape of the tip surface is usually assumed to be spherical and atomically smooth. In a spherical model, the shell has a constant, single value for the radius of curvature. The FIM image can

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be simulated by projecting the atoms in a thin shell onto a plane located at a specified distance from the specimen apex. A comparison of field ion micrographs of pure metals with shell model simulations usually illustrates a good matching in spite of fact that the model does not account for possible effects due to chemical bonds. However, in case of graphite the situation is entirely different. The process of regular field evaporation of graphite tip is illustrated in Figure 4 which shows computer simulated FIM images. The image plane is parallel to graphite layers. Parameters of the model in units of the interatomic separation a: the radius of curvature r0 = 500 a, the thickness of the imaged shell δM =0.04 a, the radius of the FIM image of 200 a. The decrease in diameter of the central atomic plane from Figure 4(a) to 4(b) indicates that atoms are removed from the surface during field evaporation.

(a)

(b)

Figure 4. Computer simulated FIM images of the atomically smooth graphite specimen formed by lowtemperature field evaporation: (a) immediately after removal of the previous atomic layer and (b) after formation of an atomic island at the center of the specimen (courtesy of V.A. Ksenofontov).

When metal tips are viewed in real time, a continuous series of collapsing rings are observed. Each time a ring collapses, one atomic layer is removed, leaving behind a surface smooth on an atomic scale. From comparison of real FIM images of carbon tips (Figure 3) with that obtained by computer simulations (Figure 4) it is clearly evident that FIM images are characterized by an irregularly distribution of image spots. It points out that lowtemperature field evaporation can not produce an atomically smooth surface of carbon tips. The mesoscopic field F at the apex of a field emission specimen is determined by Eqn. (1). For the specimen shown in figures 2 and 3 V = 6.60 kV and r0 = 297 nm, the averaged field is 4.44·109 V/m. This value is orders of magnitude lower than theoretical ones for lowtemperature field evaporation (10.3·1010 V/m) and the best image field for field ion microscopy (4.40·1010 V/m) [21]. However, it was revealed that this mesoscopic field required for low-temperature field evaporation of the carbon specimen exceeded that required to produce the helium field ion image (Figure 3). Taking into account ultra-high vacuum conditions and preliminarily cleaning the surface of the tips by field desorption and evaporation, we are confident that the localized emitting structures are not produced by

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381

chemisorption. One can arrive at a conclusion that there are the protrusions of a subnanometric thickness on the surface of the carbon tips invisible in the electron microscope. These ultrathin protrusions (nanowires) can provide areas for which the local field strength is much greater than that calculated for perfectly smooth needle-shaped specimens. The fieldenhancement factor β is about 10 or more. An analogous result was obtained by Rinzler et al. [9] in laser-beam studies on field emission from carbon nanotubes. Their results have been interpreted in terms of nanometric all-carbon atomic wires on an open-ended nanotube. The carbon atomic wires connected to the edges of graphene sheets are likely to experience very high electric field due to their high aspect ratio.

5. The Field Enhancement Factor for Linear Atomic Chains 5.1. The Linear Chain on the Tip Careful examination in TEM did not reveal protrusions, which could be characterized by such a high field-enhancement factor β (Figure 2). We suggested above that the field enhancement is caused by invisible in the scanning electron microscope objects with dimensions below the resolution limit of used TEM. The most available candidate for such a high field-enhancement factor is the normal to surface atomic C-chains with a high aspect ratio. We used the analytical model to calculate the field-enhancement factor for the single carbon chain on the needle-shaped electrode, described. It is supposed that the chain located normally on the parabolic electrode, having a cylindrical shape of height L and closed with a hemispherical cap with the radius ρ. The bond length structures of C-wires are essentially cumulenic with only a small dimerization with bond lengths between 1.27 and 1.29 Å [33]. So for appraisal the use of a constant interatomic distance d = 1.28 Å is satisfactorily accurate. The length of carbon chains Cn can be estimated as L = 1.28 ⋅10

−10

(n − 1) . The carbon atomic

chain can be considered as a conducting cylinder with the radius ρ = 1.2 Ǻ [34].

5.1.1. The Model for the Electric Field Field-induced charges are exiting near the end of the atomic carbon chain. Therefore, the hemisphere cap and the cylinder under the cap can be replaced by two charges and dipoles taking into account the image forces at the conductor surface [35]. The contribution of the needle-shaped carbon specimen can be estimated using a model of the field ion emitter approximation by equipotential surfaces for a point charge situated at the end of a uniformly charged filament [36]. The potential around the hemispherical end of the chain on the surface of a needle-shaped specimen with the radius of the apex equal to r0 can be expressed at an arbitrary point (x,y) as ⎧⎪

ϕ (x, y ) = F0 ρ ⎨

[

L

⎪⎩ x + ( y + 2 L ) 2

]

2 1/ 2

L ρ ⎞ ⎛ ⎟− 2 ⎜1 + ⎝ 2L ⎠ x + y 2

(

[

)

1/ 2

ρ ⎞ ρ2y ⎛ ⎟− 2 ⎜1 + ⎝ 2L ⎠ x + y 2

(

]

⎧⎪ x 2 + ( y + r + L )2 1 / 2 ⎫⎪ 0 + F0 r0 Ln⎨ ⎬+C , r ⎪⎭ ⎪⎩ 0

)

3/ 2

−

[x

ρ 2 ( y + 2L) 2

+ ( y + 2L )

]

2 3/ 2

⎫⎪ ⎬ ⎪⎭

(2)

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Igor M. Mikhailovskij

Where F0 is the external field at the needle-shaped specimen, C is the constant potential and the carbon chain is oriented along the x-axis. The “surface” of the chain and carbon tip constitute an equipotential surface φ(x,y)= 0. Figure 5 shows the calculated potential distribution in the vicinity of nanowire on the parabolic specimen formed by field evaporation at +15.0 kV.

15

23 20

10

z, C

17 5

14 11

0

8 4

-5

0

-15

-10

-5

0

5

10

15

x, C

Figure 5. Equipotential lines for a terminating C-wire on the paraboloidal electrode. The atomic wire is at 0 V.

The equipotential lines are bent around the carbon nanowire producing, as expected, an enhancement of the applied electric field. Figure 6 gives our calculated field-enhancement factor β as a function of aspect ratio L/ρ. We have performed such calculations for several radii r0 of the paraboloidal electrode and obtained nearly the same results. The only quantity that matter in this problem is the ratio L/ρ. 50 40

β

30 20 10 0

10

20

30

40

L/ρ

Figure 6. Field-enhancement factor for atomic chain on the tip as a function of aspect ratio L/ρ.

Preparation and Characterization of Monoatomic Carbon Chains

383

In fact at r0 >> ρ we obtained by fitting within an error of 2 % the following empirical formula for the field-enhancement factor at the end of carbon chain on the needle-shaped electrode:

β ≈ 3.4 + L / ρ .

(3)

By using this formula, we showed that observed values of a local field strength can be obtained at the open end of Cn-chains with n = 10-12.

5.1.2. The Local Magnification of FIM Images The number of atoms in the linear chain can be obtained by an independent comparison of experimental and theoretical values of local magnifications [21] for field-ion images. Determination of the ion and electron trajectories is an essential aspect when interpreting results from experiments with the field ion microscopy and field electron emitters. When a nanoprotrusion is on the top of a conventional needle-shaped FIM specimen, the half-width angle of the divergent beam is reduced by factor of ten. This effect is a result of a selffocusing process caused by the changed configuration of the electric field in the close vicinity of the sample. Conducting line chains distort and compress the equipotentials in their vicinity. This gives rise to local field enhancement and increased magnification. The resultant magnification will exceed that of the supporting tip but will not correspond to the radius of the chain because of a strong compression factor. The behaviour of ion trajectories in the FIM has been treated analytically, for a general field which is symmetric about the main x-axis of the microscope [21]. The ion paths are not only independent of the mass and charge of the ion, but also of applied voltage, being wholly a function of the field distribution, which is controlled by specimen and counter electrode geometry. The locations of the image spots corresponding to individual surface atoms are identical for hydrogen, helium or any other image gas. The image of a given specimen is also completely unaffected by raising or lowering the applied voltage, provided that the geometry of the specimen end-form remains unchanged. The calculations show that the effect of the specimen shank is to introduce some degree of curvature into the trajectories, compressing them towards the axis of the microscope. The comparison of the electric lines and the ion trajectories in a region close to the specimen apex showed that the ion trajectories do not bend as sharply as the electric field lines since the ions acquire kinetic energy near the specimen surface. Field evaporated ions are formed right at the surface and imaging gas atoms are ionized at the critical distance away from the surface. Therefore, the trajectory of the field ionized gas atoms is not the same as that of field evaporated carbon ions. In the idealized case where the specimen and screen of the FIM consisted of two concentric spheres, the trajectories of ions would be straight lines normal to the specimen surface and the magnification of the image M would be given simply by the ratio of the specimen-to-screen distance R to the apex radius of the specimen r0. Most of the difficulties in determining a local magnification arise from the necessity of an adequate model for the electric field in the vicinity of the wire and for the overall shape of the needle-shaped tip, which, in general, does not coincide with analytical model. Since the image size should be independent of length for a long wire this effect must arise from the

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Igor M. Mikhailovskij

influence of the needle-shaped substrate on the potential distribution near the tip of wire. It can be seen that an image compression must result from the finite length of the atomic wire. This problem was solved by analytical methods with minor simplifications only for the particular case of a hemisphere on the emitting tip [21,37]. It was shown that resolution and magnification for small protrusions on the surface of tip are much greater then that calculated for a perfectly smooth tip. An approximate solution was developed by Rose [37] for the special case of a small hemispherical protrusion of radius rp, at the apex of a paraboloidal specimen of end-radius r0. The overall magnification of the protrusion M', is given by

M '≈ 1.1

r0 M, rp

(4)

where M is the conventional magnification of the image (without microprotrusion). One consequence of this local magnification effect is that the apparent size of, for example, a precipitate particle in an FIM image may differ from its true size. We calculated the magnification for ions, originating from a point at the end of a nanowire (figure 5), based on numerical calculations of the equation of motion in an electric field determined by Eqn. (2) from a starting position, where the ion is assumed to be at rest. For the case of the single carbon chain on the needle-shaped electrode (figure 5), we obtained the semiempirical expression for the local magnification M´ :

M ′ ≈ 1.6

L H ⋅ , r0 ρ

(5)

where r0 is the radius of the paraboloidal tip, H is the interelectrode distance. One consequence of this local magnification effect is that the apparent size of a nanosized object in an FIM image may differ from its true size. This expression was used for an independent determination of the average wire length and was shown again that in our experiments n was of order of 10.

5.2. Multistage Nanostructures In the special case (often realized) of multistage emitters [38], the field enhancement factor can be defined as β = F/FM ,

(6)

where F is a local field and FM is the macroscopic field defined as FM = V/D. Here D is the interelectrode distance in geometrical configuration resembling a parallel-plate capacitor. In our experiments a voltage corresponding to the threshold of low temperature field evaporation ranged from 2.5 to 22 kV was applied on the electrodes with D of about 5 mm. As was shown above, low temperature field evaporation of carbon takes place at electric field strength F of about 70-80 V/nm. Consequently, one can obtain the calculated value of the field-

Preparation and Characterization of Monoatomic Carbon Chains

385

enhancement factor in the range of 1.59·104 < β < 1.33·105. The extremely large fieldenhancement factors have been reported previously from field electron emission measurements made over large areas and individual tips [38-40] Huang et al. [38] showed that a giant field-enhancement factor (up to 1.88·104) is a result of a multistage structure of the nanoemitters. Their structure is characterized by an order of magnitude smaller features (nanotubes or nanowires), growing at the tips of larger features. A multistage structure, with each stage much smaller than the previous one is characterized by the total field-enhancement factor at the tip of the smallest feature given by m

β tot = ∏ β i ,

(7)

i =1

where m is the total number of stages, i is the stage index. The experimental value of the total field-enhancement factor for one of typical specimens formed by evaporation at electric field Fev = 75±4 V/nm can be found from

β tot =

Fev D . V

(8)

By using the value of D equal to 5±0.5 mm, we obtain the experimental value βtot = 6.47·104. The value of β1 is correspondent to the mesoscopic field-enhancement factor at the apex of carbon needle-shaped emitters. The electric field and potential distribution between such a tip and a counter-electrode can be considered in terms of the hyperboloidal model [21]. In this approximation, the tip and counter-electrode are represented by two confocal hyperboloids of revolution. A local field F at the tip surface is a logarithmic function of the radius of curvature and D. Taking r0 to be the radius of curvature at the specimen apex, simple transforms of the equation for F yielded the following expression for the mesoscopic fieldenhancement factor in the hyperboloidal approximation:

β1 =

2D . ⎛ 4D ⎞ ⎟⎟ r0 Ln⎜⎜ ⎝ r0 ⎠

(9)

Thus, from equation (9) for the specimen under consideration, the mesoscopic fieldenhancement factor is equal to 3.54·103 and β1 << βtot. Such a discrepancy shows that the carbon specimen formed by field evaporation has a multistage structure and a local electric field is determined to a considerable degree by smaller topographic features. A certain contribution to βtot is introduced by an additional enhancement of the local field in the vicinity of the nanoscale undulations. But it is intuitively clear that observed nanoprotrusions with low aspect ratios can not responsible for a high discrepancy of the experimental and theoretical field-enhancement factors. To our knowledge the fieldenhancement factors were calculated only for protrusions with high ratios of the total

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Igor M. Mikhailovskij

protrusion length to the radius (ν = L/ρ ≥ 1). In this range of the ratio ν two analytical models give satisfactory results: the “hemisphere on the plane” model and the “hemisphere on a post” model [35 and references therein]. In terms of the hemisphere on the plane model the potential at the sphere surface is given by a superposition of the field of a dipole of strength p = 4πε 0 ρ FM located at the centre of 3

a hemisphere of radius ρ and a planar charge sheet of density ε0FM where ε0 is the electric constant:

ϕ=−

p cosθ + rFM Cosθ . 4πε o r 2

(10)

There (r, θ) denote the spherical coordinates with the origin in the sphere center. The total potential at any point on the hemisphere surface is zero: the emitter plane and the hemisphere boss surface constitute an equipotential with φ = 0. Figure 7 shows the calculated potential distribution in the vicinity of the hemisphere on a plane. The equipotential lines are bent around the hemisphere producing an enhancement of the applied electric field. This data can be used for a determination of the field-enhancement factor for smooth protrusions with ν < 1. To pursue this object, we calculated a family of equipotentials such that one coincides with the protrusion surface and one with the counter-electrode. Such a representation for rather short nanoprotrusions was given in Cartesian coordinates (x, y). 4

y

3 5

2

3

1

1 0

0 -1 -6

-4

-2

0

2

4

6

x Figure 7. Equipotential lines generated by “the hemisphere on a plane”. The numbers near the equipotentials denote the potential in volts.

It was found that the geometry of a typical nanoprotrusion could be fitted accurately by one equipotential surface from the set surrounding a charged body which consisted of an isolated hemisphere on the plane. Figure 8 illustrates this possibility for an approximation of one of rather short nanoprotrusions on the specimen with ν = 0.24. The surface of the nanoprotrusion coincides with an equipotential surface of value φ(x, y) = 4 V.

Preparation and Characterization of Monoatomic Carbon Chains

387

5 nm a

b Figure 8. Electron microscope image (a) of the nanoprotrusion on the carbon tip formed by field evaporation and its electrostatic approximation (b). For the matching equipotential φ(x, y) = 4 V. (Reprinted with permission from [19]. © 2007, Institute of Physics.)

Figure 9 (curve 1) gives the nanoscopic field-enhancement factor β2 obtained from equation (10) as a function of aspect ratio for ν < 1. For ν < 1, the values of ρ were determined at the apex of the matching equipotentials. The field enhancement was obtained from equation (10) using the gradient of φ at the apex equipotentials. It demonstrates the wellknown exact result that the local field enhancement factor at the hemisphere apex (ν → 0) is 3. In the range 0.01 ≤ ν ≤ 1 the factor is represented, to within 4 %, by

β 2 = 1.19 + 2.30ν − 0.49ν 2 .

(11)

In the vicinity of ν = 1 our results are very close to that for ν > 1 (curve 2 in figure 9) represented by the approximation [35]

β P = 1.2(ν + 2.15)0.90 ,

(12)

obtained for the “hemisphere on a post” model.

5 2

β

4

3 1 2

1

0

1

2

3

L/ρ

Figure 9. The field enhancement factor as function of the aspect ratio ν < 1 (curve 1). The dotted line 2 corresponds to equation (12) for ν > 1.

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Igor M. Mikhailovskij

For the average value of the aspect ratio ν = 0.30 for nanoprotrusions on the specimen under consideration the equation (11) gives the field-enhancement factor β2 of 1.84 and β1β2 = 6.51·103. The magnitude β1β2 is much smaller than the experimental total fieldenhancement factor (βtot = 6.47·104). To interpret this result, we can conclude that the carbon emitters with nanoprotrusions formed by field evaporation are not atomically smooth electrodes but have invisible in TEM atomic-scale features possessing high aspect ratios. An atomic-scale β3 is equal to βtot / β1β2. The specimen shown in Figure 2 has the atomic-scale field-enhancement factor β3 = 9.94. This value is corresponding to atomic-scale protrusions at the apex of nanoprotrusions. For flat parts of the needle shaped carbon specimen β2 ≈ 1 and therefore, β3 = 15.3. The most available candidate for such a high field-enhancement factor β3 is the normal to surface atomic carbon wires with a rather high aspect ratio. To determine the length of C-nanowire on the needle-shaped electrode we used the model [34], in which the nanowire stands normally on, having the length L and radius ρ = 1.2 Ǻ. Taking into account that ρ << r0 the mesoscopic electrode (hyperboloid) can be replaced by plane. In this case the equation (7) is valid and gives the aspect ratios of 8.34 and 14.79 and the lengths of the nanowires of L = 11.9 and 18.4 Å for β3 = 9.94 and 15.3, accordingly. The numbers of carbon atoms in a linear chain estimated as n ≈ L/d + 1 were equal to 10 and 15 for these two field-enhancement factors on the surface of a typical carbon tip, where d = 1.28 Å (§5.1). The size of clusters (n = 2-7) produced by pulse field evaporation is smaller then those above calculated for linear chains on the needle-shaped electrode. As was shown [27] that small carbon clusters could be the result of the decomposition of field evaporated heavier cluster ions within the period of acceleration in the spectrometer. The spectral width and the low energy tails of mass marks are correspondent to the flight time of the primary clusters before dissociation of the order of 10-12 s. Due to the low-energy tail of the C5+ mass mark and the very small voltage signal of the C6+ cluster, it is difficult to determine the abundance of the C6 isomer.

6. Atomic Resolution of One-Dimensional Nanomaterials The imaging process in field ion and electron microscopes has continuously been a topic of debate [20,21]. However, there are not convincing arguments which explain the shape of the image and other intrinsic peculiarities of the formation of field ion and electron images of 1D nanomaterials. In a general case, the image magnification is given by

M′=

R , f k r0

(13)

where fk is an image compression factor which takes into account the fact that field ions do not follow exactly the radial direction of the nearly hemispherical tip surface because of the tip shank and other supporting leads. For simple field emitter geometries, fk can be calculated by determining the ion trajectories. It can always be found from field ion images where angular separations among different crystallographic poles are well known. In normal tip geometry the magnification is reduced by an image compression factor fk equal to 1.5-1.8

Preparation and Characterization of Monoatomic Carbon Chains

389

[21]. The image compression factor fc for ions, starting from the end of the linear chain (Figure 5), was determined by numerical calculations of the equation of motion in an electric field described by Eqn. (2). Figure 10 shows the dependence of the image compression factor on the number atoms in carbon chains on the needle-shaped electrodes of the radius 100 (1), 200 (2), and 500 (3) nm.

30 3

25

fc

20 2

15 10

1

5 0

1

10

100

n

Figure 10. The FIM image compression factor for the case of the single carbon chain on the needleshaped electrodes of the radius 100 (1), 200 (2), and 500 (3) nm.

The FIM image resolution is primarily determined by the ion beam broadens on its way to the screen by the slightly compressed radial projection. Besides the broadening due to the radial projection of the ions, the tangential component of the momentum of the particle will further broaden the spot size. De Castilho and Kingham [41] expressed the resolution of the FIM image in terms of the following expression: 1/ 2

⎛ 2 2 f c hρk f 1 / 2 ⎛ fc 2 ρ 2 ⎞ ⎞ ⎜ ⎟ε ⎟ + 16k f ⎜ δ = δ0 + 1/ 2 ⎜ eV ⎟ T ⎟ ⎜ ( ) 2 π meV ⎝ ⎠ ⎠ ⎝

,

(14)

where e is the elementary charge, h is Planck’s constant, m is the mass of the image gas ion, εT is the transverse thermal energy associated with the imaging gas atom when it is ionized and the δ0 is the diameter of the ionization zone above a given atom. The value of δ02 in this equation relates to the ionization area and depends on the field strength as well as some other parameters. Lacking detailed calculations, one may assume that δ0 is of the order of the radius of the image gas atom (0.12 nm). The second term is correspondent to the broadening of the ion image to the uncertainty principle. The third term takes into account the spreading of the ion beam due to the thermal energy of the image ions. Atoms of the imaging gas, striking the surface of the nanotube or atomic chain, have a high kinetic energy εT of about 0.1 eV, acquired when the polarized atoms are drawn into the region of high nonuniform electric field. This energy is partially diminished in the collision of the gas atom with the surface of the cooled sample. Upon repeated collisions the gas atoms lose kinetic energy, leading to thermal accommodation down to the temperature of the

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sample. In a theoretical determination of the resolution of the FIM it is assumed that the atoms of the imaging gas have accommodated to the sample temperature and, accordingly, the accommodation factor is equal to unity [20]. Figure 11 shows a calculated dependence of the FIM image resolution on radius ρ for helium atoms accommodated to various temperatures. 0.32

δ (nm)

0.28 77K

0.24 0.20

4.2K

0.16 0.12 0

10

20

30

40

50

ρ (nm)

Figure 11. Theoretical FIM image resolution for helium atoms accommodated to cryogenic temperatures. (Reprinted with permission from [42]. © 2007, American Institute of Physics.)

The magnitude of ρ in Eqn. (14) is correspondent to the radius of the closed carbon nanotubes or the effective radius of the end of the carbon chain. The compression factor for CNT is assumed to be ~10. As follows from the data in Fig. 11, at nanotube radii of ~1 nm the resolution of the microscopic is equal to the theoretical limit (0.13 nm). It is certainly so for the case of carbon atomic chains characterized by ρ = 0.12 nm. As a result, the end atom of a carbon chain is perfectly resolved in the field ion microscope (Figure 3). Due to an enhanced FIM resolution of 1D structures, a direct imaging of the atomic structure of the cap of closed carbon nanotubes was achieved for the first time by the methods of high-resolution field ion microscopy with sample cooling to liquid helium temperature [42]. We used in the study the products of gas-phase catalytic pyrolysis of hydrocarbons in the form of graphitized fibers containing carbon nanotubes. A fiber is a graphitized carcass containing a large number of “ropes” of cylindrical and conical multiwall carbon nanotubes which are oriented approximately along the axis of the fiber and do not contain catalytic particles. The surface of the samples was subjected to controlled field evaporation to clean it of adsorbed gases, remove the microscopic asperities formed during the electrochemical etching, and form an apex of hemispherical shape. As was shown above in section 3, unlike metallic samples in strong electric fields, carbon materials evaporate not atom-by-atom but in the form of atomic clusters. Because of this, the hemispherical surface formed by field evaporation is microscopically rough. FIM images of needle-like carbon specimens formed by field evaporation are usually characterized by an irregular distribution of image spots as that in case of the graphite specimens (Figures 3 and 12(a)). This type of FIM contrast is correspondent to the random character of the low-temperature cluster field evaporation and

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surface defects, such as process-induced microroughness, of a nanostructured carbon material. In the process of field evaporation the interior region of the needle-shaped samples was exposed, irregularly revealing nanotubes characterized by another form of ion images. The brightness of the nanotube cap with a diameter of about 0.90 nm is significantly greater than the brightness of the image of the graphite matrix. On the surface of the cap one observes several regularly spaced dark and bright spots. Analogous contrast was observed previously on field ion microscope images of the cap fragments of multiwall carbon nanotubes of relatively large diameters [43,44].

(a)

(b)

(c)

Figure 12. Field ion microscope images of a graphite fiber (a) and carbon nanotube of diameter 0.90 (b). Diagram (c) illustrates the predominance of hexagonal rings in the structure of the cap of the carbon nanotube. (Reprinted with permission from [42]. © 2007, American Institute of Physics.)

The value of fk was determined empirically as the coefficient of angular image compression, equal to the ratio of the angle of field ion emission to the angle between lines drawn from the tip to diametrically opposed peripheral parts of the FIM image of the nanotubes. Measurements showed that for the tip-nanotube configurations used in the present study the local compression factor was about 13, and the field factor χ was equal to 300. The full resolution of the structure of the fullerene caps of the carbon nanotubes achieved in these experiments is apparently due to deep cooling of the samples. On cooling to 4.2 K a monolayer of helium atoms can form on the surface of the tip, which is found in a high electric field [20], and this improves substantially the conditions of accommodation of the polarized helium atoms.

7. Field Emission Characteristics The development of field electron sources based on the phenomenon of emission in high electric field is a promising direction of application of new nanostructural carbon-based materials. Advanced field electron sources often are regarded as the most important shortterm application of carbon nanotubes and nanowires. Multi-emitter carbon cathodes are characterized by high local electric field enhancement factors and, accordingly, relatively low values of working voltages. The same advantages can be used for the development of new field-emission sources of inert gas ions [45,46] for various fields of physical metallurgy,

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accelerator technology, supersonic molecular flux detectors, etc. Difficulties encountered in the creation of highly effective ion sources are mostly related to the high field strengths necessary for the field ion emission and to the irregular character of the process of emitter surface shaping by means of field evaporation. The phenomenon of cluster evaporation of carbon materials in an electric field with a strength exceeding 5·1010 V/m (see section 3) substantially limits the use of high-field surface finishing nanotechnology. 7.1. Field ion emission We demonstrated the possibility to create a wide-area field-emission source of inert gas ions with based on a nanostructural carbon, which can operate at a relatively low applied voltage [36]. The proposed field ion emitter represented a tungsten point with a ball of aluminum-filled epoxy resin composite formed on the point tip. The surface of the ball with a diameter of about 10 μm was covered with a disperse layer of the catalytic carbon vapor deposition products obtained by gas-phase pyrolysis of hydrocarbons. An FIM and TEM analysis of the deposited carbon material showed the presence of multi-wall nanotubes in the graphite matrix. The ion source was tested in a FIM with a sample holder cooled by liquid helium. The residual gas pressure in the working chamber was 10-8 Pa, and the partial pressure of helium could be varied from 10-3 to 10-2 Pa. The field strength at the spherical surface of the emitter, which is necessary for the field ionization, was created by applying a positive potential variable within 1–10 kV. The distance D from the ion source apex to the counter-electrode was 3–10 mm.

(a)

(b)

(c)

Figure 13. FIM images of the surface of a multi-emitter field ion source obtained by a high-field shaping at a voltage of (a) 5.80 (b) 6.45, and (c) 5.82 kV.

As shown in Figure 13(a), at first the main contribution to the ion current was due to separate microprotrusions present on the hemispherical surface. The surface density of emitting centers was in the range (5–20)·109 m-2. In order to increase the homogeneity of ion emission, the applied voltage V was increased at a rate of about 100 V/min. The field evaporation led to the removal of sharpest and highest microprotrusions, which were characterized by the maximum local electric field enhancement factors. As can be seen from the series of FIM images in Figure 13 (a-c), the surface density of emitting centers was gradually increased by almost two orders of magnitude to reach a level of 2·1012 m-2. It is difficult to determine atomic structures of nanoprotrusions. Upon completion of the high-field surface shaping process, the ion beams emitted from the neighboring centers exhibited

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overlap, thus creating a nearly homogeneous FIM pattern. The electric field strength at the emitting element of the field ion source can be expressed as F = αβ V / D , where α and β are the meso- and nanoscale electric field enhancement factors, accordingly. The value β is determined by the configuration of the field ion source and is less dependent on the pulling electrode shape. The value of α can be estimated using a model of the field ion source approximation by equipotential surfaces for a point charge situated at the end of a uniformly charged filament. These equipotential surfaces are characterized by a dimensionless form factor ω that varies from 0 (for a parabolic source) to infinity (for a spherical source). Within the framework of this model, the mesoscale electric field amplification factor can be expressed as

α=

D (ω + 1)

.

(15)

⎛ ⎛ D ⎞⎞ ωr r0 ⎜⎜ ω − 0 + ln ⎜ ⎟ ⎟⎟ D ⎝ r0 ⎠ ⎠ ⎝

There r0 is the curvature radius of the field ion source apex. For the configurations of field ion sources used in that study, the estimates of α varied from 250 (ω = 1) to 650 (ω = 10). The calculated values of the nanoscale electric field amplification factors fell within 70–150. The total field amplification factor varied within (2.5-6.5)·104 in the course of the ion source surface shaping by field evaporation. These estimates are close to the maximum electric field amplification factors reported for the apexes of carbon nanotubes [38]. Increasing the positive voltage on the carbon tip with the cap resulted in a growth of field ion emission up to 2 nA (Figure 14). Such a behavior of field ion images is characteristic of very small aggregates of atoms acting like a lengthy conducting protrusion on the tip surface. 0.5 0.0

LgIi

-0.5 -1.0 -1.5 -2.0 3.8

3.9

4.0

4.1

LgV Figure 14. Current–voltage characteristics of the helium field ion source at 4.2 K.

Figure 14 shows an experimental current–voltage characteristic of the helium ion source plotted on the logarithmic scale. In the low-field region the ion current increases extremely rapidly with voltage, amounting to 8-th power of the field. In this case the ion current is

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measured in the units nA. The slope for relatively high-field strengths (above 4 ·1010 V/m) is correspondent to about the 3-d power of the field strengths. This character of the I–V curves of the multi-emitter field ion source shaped by field evaporation is close to that of the characteristics of metal field emitters [20,21], but the slope of the low-voltage part is about four times smaller than that typical of the metal emitters. In the region of low field strengths, the total current is limited by the rate of ionization for helium atoms colliding with the emitter surface. For field strengths above that of the best image, the field ionization rate is sufficiently high to provide for the ionization of virtually all helium atoms trapped by the polarization forces. A lower rate of current growth increase in this region is explained by a small increase of the volume of the region of trapping related to the attraction of inert gas atoms (polarized in the inhomogeneous electric field) to the region of maximum field strength at the surface of a multi-emitter field ion source. A relatively low field sensitivity of the ionization rate in the region of high fields indicates that the total ion current is only limited by the supply of helium atoms from the gas phase. This is an additional evidence of a high efficacy of the field strength leveling over microprotrusions in the course of the emitter surface shaping by field evaporation [47].

7.2. Field Electron Emission Rinzler et al. [9] for the first time conjectured that atomic carbon chains may provide the vital point field electron cathodes. The delocalized, cylindrically symmetrical π-bonding along the linear carbon chain is responsible for nearly metallic screening of the electric field. Ab initio calculations of carbon chains in an applied uniform electric field showed these chains to screen the electric field as efficiently as metal rods of the same configurations [48] and revealed that unlike bulk carbon, monoatomic C-chains are metallic independent of doping [49]. The character of the bonds in the linear chain was found to be of cumulene-type, i.e., with all nearly equivalent bond lengths. The electronic structure determines the physical properties and the equilibrium geometry of nanoobjects. On the contrary, geometrical structure determines the local electric field confining the electrons. This close interdependence underlies the complex behavior of monoatomic chains in electric fields. The high aspect ratio and small dimensions of the carbon chains are responsible for the field concentration. The maximum induced field strength outside the end atom of the linear carbon chain was more than 100 V/nm. Field electron emission properties of atomic carbon chains were studied in detail using of sophisticated theoretical methods [33,34], and was shown that a carbon chain is really a good candidate as a nanomaterial for fabricating field emitters. To investigate the field emission characteristics of linear carbon nanowires, their current versus voltage characteristics were measured. The question arises whether the FowlerNordheim (F-N) theory of field emission [20] is still valid for carbon atomic wires, as it was elaborated for a surface that appears flat for “imaged” electrons, which might not be the case for nanowires on account of their ultrahigh curvature at the wire apex. Figure 15 shows a typical result as a F-N plot for the multiwire field emitter formed by field evaporation of the carbon tip. In this case the electron emission pattern was formed by about twenty nanoprotrusions. The data follows a straight line in the F-N coordinates, which is an apparent indication of a field electron process and a possibility to qualify linear carbon chains as metallic atomic wires. As can be seen in figure 15, significant current instabilities took place

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at 77 K and some long-term current drift occurred. The typical current fluctuation is of the order of 1 μA. The exact origin of the current fluctuations is not clear at the moment. These instabilities of the field emission current may be due to the residual gas ion bombardment and the process adsorption – desorption [50]. -26.5 -27.0

2 ln(I/V )

-27.5 -28.0 -28.5 -29.0 0.17

0.18

0.19

0.20

0.21

0.22

3 -1 10 /V (V )

Figure 15. Fowler-Nordheim plot of the multiwire field electron emitter at 77 K.

The field electron emission patterns usually consist of doublets and/or singlets. Figure 16 shows field electron images of the needle-shaped specimen formed by field evaporation at +14.50 kV. The images obtained at negative voltages 1.21 (a) and 1.38 kV (b).

(a)

(b)

Figure 16. Field electron images of the carbon tip after pulse field evaporation obtained at the bias voltages 1.21 (a) and 1.38 kV (b).

In Figure 16 (a), the doublet consisted of two elliptical spots are clearly resolved. This microphotograph was taken at current. An increase of voltage was accompanied by the stepwise change from a doublet to a singlet shown in Figure 16(b). A further increasing the negative voltage on the tip resulted in a growth of field electron emission from this image spot up to about 1 μA. Such a behavior of field electron images is characteristic of very small aggregates of atoms acting like a lengthy conducting protrusion on the tip surface [50,51]. The images in figure 16 are selected out of about 103 microphotographs which revealed the great variation in size, shape and intensity of the duplets and singlets. The variation in the emission patterns can be ascribed to variation in the length of linear carbon chains connected to the edges of graphite layers.

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The measurements of the current-vs-voltage characteristics were performed also at 21K in a vacuum system with the residual pas pressure of 5·10-7 Pa. Figure 17 shows the typical experimental current-vs-voltage plot (curve 1) for the multiwire field emitter formed by field evaporation of the carbon tip. In these cases the electron emission pattern was formed by about twenty nanoprotrusions. An observed in figure 17 suppression of current instabilities at a deep chilling up to 21 K can be explained by decreasing the ion bombardment and the rate of surface migration of adsorbed atoms. Because of the lack of reliable experimental data for emission properties of the linear carbon chains our results can only be compared to theoretically determined data. In theoretical studying the field electron emission from linear carbon chains, it was found [34] that these chains can be viewed as atomic-scale metallic wires with a diameter of 2.4 Å. In keeping with Refs. [33,34], for one-dimensional systems like the linear atomic chains, quantum size effects are of considerable significance, and the total field electron current is the summed contribution of single-particle quantum levels. The equivalent to the FowlerNordheim law for an atomic carbon chain is the modified equation with a prefactor linear in the electric field

⎛ −b EHOMO I = aF exp ⎜ ⎜ F ⎝

3/ 2

⎞ ⎟ ⎟ ⎠

(16)

where a and b are modified FN constants, EHOMO is the energy of the highest occupied molecular orbits. 12 10

1

ln(I/F)

8 6

3

4

2

2 0

1.0

1.2

1.4

1.6

1.8

2.0

2.2

2.4

2.6

1/F (Å/V)

Figure 17. Modified Fowler-Nordheim plot for the multiwire field electron emitter at 21 K (1) and the theoretical data [15] for C10 (2) and C11 (3) monoatomic carbon chains. (Reprinted with permission from [19]. © 2007, Institute of Physics.)

Figure 17 shows a modified F-N plot for the current-vs-voltage characteristic of the multiwire carbon emitter formed by field evaporation at 21 K (1) and the theoretical data [34] for two typical even and odd Cn-chains with n = 10 (2) and 11 (3) at 0 K. In this case the current and the field are measured in the units nA and V/Ǻ, correspondingly. The data follow

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the best fit straight lines in the modified FN coordinates and the plot slopes are 8.13 ± 0.10, 6.68, and 8.63 V/Ǻ for the carbon emitter, C10 and C11 monoatomic carbon chains, correspondingly. Taking into account that the validity of the FN theory has been established experimentally to within about ± 15% [20], the plot slope for the carbon emitter under consideration and that for theoretically studied C11-chains can be considered as equals. We can conclude that even linear clusters with lower EHOMO are marginals in the morphology of multiwire carbon emitters. These emitters formed by field evaporation at 21 K predominately consist of odd carbon atomic chains. This conclusion is in accordance with the results of comparative studies of thermodynamic stabilities of odd and even atomic carbon chains [33]. The length of the carbon nanowires in our cathodes could not be controlled. The natural fluctuations in the length of self-assembled nanowires contributed to the non-uniformity of field emission in the device. Despite the fact that uniformity and stability the field emission from the carbon multiwire cathode is not yet sufficient these results lend strong support to the conjecture of Rinzler et al. [9] that linear carbon chains may provide the ultimate atomic-scale field emitters.

8. Atomic Simulation of Unraveling and Disintegration 8.1. High-Field Induced Unraveling of Graphene It was shown in Ref. [9] that monoatomic chains can be formed by unraveling suggested in field electron emission experiments with carbon nanotubes, where the electric force unravels the multiwalled tube like the sweater. An analogical conclusion was made as a result of computer experiments on mechanical pulling of carbon nanotubes [10]. It was revealed in these molecular-dynamics simulations of CNT deformation that CNTs undergo a process of necking up to formation of a single chain configuration, spanning the nanotube fragments. As was revealed in [52], it is possible to use the electron bombardment in a TEM to elaborate a CNT synthesis method from an amorphous carbon film. In earlier irradiation research [53] was suggested that the end stage of radiation induced CNT reconstruction should involve the creation of an atomic carbon chain. Atomic chain formation was verified for short lengths of time (seconds) both experimentally and through computer simulations [11,54]. Authors [11] find that there is a significant probability for formation of a stable bridge with the smallest cross-section of one atom. It was concluded that all the evidence combines to show that a freely suspended atomic chain was formed. Chain lifetimes of up to several minutes prior to fracture were registered for the lowest current density conditions. In computer experiments [11], it was observed the formation of a single or a double wire, or more complicated wire structures, between the tube ends. Authors obtained linear carbon chains, which had a remarkable length up to 14 atoms. The high rate plastic deformation resulted in numerous collisions between the wires, and at last only one wire survived. An additional increase of the distance between the tube fragments does not break this wire, which elongates by increasing the number of atoms that burst forth out from both ends into the chain. Linear carbon chains were also revealed in samples of double-wall CNTs annealed at temperatures between 1000 and 2000°C [55]. From the comparison between the first-principles calculations and the resonance Raman experimental results authors concluded that some unusual Raman features are related to short C5 and C7 linear carbon chains, interconnecting the CNT surfaces.

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Detailed calculations of the graphene unraveling were carried out to examine the mechanism and trend for atomic chain formation in high electric fields [18,19]. The numerical simulations were performed using the classical molecular dynamics (MD) method, employing the short-range Tersoff-Brenner bond order potential, which reproduces lattice constants, binding energies, and the mechanical properties of diamond, graphite, and carbon nanotubes. The electric force producing an axial tension is localized at the top of the chain. In our molecular dynamics modeling the electric force was 3.8-5.6 nN. Figure 18 shows a linear carbon chain extending out from the graphene layer under the action of electric field at the beginning (a) and after 3.2·10-14 s (b), 6.4·10-14 s (c) and 9.6·10-14 s (d). The nanowire held taut in this case by the electric force of 5.1 nN.

(a)

(b)

(c)

(d)

Figure 18. Forming of the carbon nanowire during the unraveling of a graphene sheet in a high electric field. Labels (a)–(d) denote the various unraveling stages.

At first, the dangling bond at the graphene edge orients along the electric force. After that the weakest link has proved to be the nearest bond aligned along the applied electric force. This bond stretches and finally breaks up first under a local force produced by a high electric field (Figure 18(b)). After that the atomic wire turned longer and straightened under the action of an electric field. The resulting effect is to increase the chain length by two atoms without any change in bond order of the open edge of the graphene sheet. The atomic chain extending out from the graphene edge held straight and taut by the electric field. Further pulling of the atomic chain by the electric field repeats the described sequence of events,

Preparation and Characterization of Monoatomic Carbon Chains

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unraveling the linear carbon. The second bond normal to surface the graphene sheet also starts to stretch and breaks up (Figure 18(c)). The orientation of electric force favors breaking the bonds normal to the open edge of the graphene layer. It is also essentially that the failure force of the linear carbon chain is slightly higher than those of the graphene bonds. Therefore, the disintegration process is successive and clearly localized, with sudden complete drop in the local failure force due to the bond breaking. The energy per bond in general depends in a complicated way on the geometry. However, as was shown by Abel [56] and Tersoff [57], for covalently bonded systems the most important variable appears to be the coordination number. The bond strength of carbon atoms is a monotonically decreasing function of the number of nearest-neighbour atoms close enough to form bonds. If the bond energy decreases rapidly with the coordination number, then the linear atomic chains along with the diatomic molecule will be characterized by a high value of the bond strength. This effect for carbon interatomic bonds quantitatively described by the Brenner version [58] of the Abel-Tersoff bond-order formalism used in our simulation.

(a)

(b)

Figure 19. Linear carbon chains on a graphene sheet without an electric field at step (a) and kink (b) sites.

Figure 19 shows two carbon chains formed as a result of the high-field unraveling of the graphene sheet at 0 K. Orientations of the chains without an electric field are determined by the equilibrium angles between atomic bonds of the nearest neighbors indicating a high extent of orientational correspondence of the liner chains and the perfectly defined graphene edges.

8.2. Dissociation and Explosive Heating of Carbon Atomic Chains Snapshots of the dissociation dynamics of a single C-chain extending out from the graphene sheet is in Figure 20. One can observe an eight atom chain that detached from the graphene edge after the high-field unraveling (Figures 20(a) and (b)). Carbon species are evaporated from a graphite surface initially as linear clusters which decompose mostly into smaller atomic clusters (Figures 20(c) and (d)). The number of C3 clusters in the field evaporation spectra was prominently large, indicating their preferred formation during fragmentation of the chains. These results of simulation confirm given above conclusion that the carbon clusters experimentally registered in the field evaporation spectrum are not occurred right at the specimen surface but are produced in a high vacuum by dissociation into

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several particles (see section 3 and the review [59]). It is also become obvious that the numbers of atoms in registered carbon clusters are much smaller as they are formed by the graphite unraveling. The carbon clusters were dissociated at a distance of about 1 nm away from the specimen apex, which is reached within a travel time of about 0.1 ps. The time of cluster disintegration found in MD simulations is an order of magnitude less then that obtained by analysis of the energy distribution in the time-of-flight spectra of lowtemperature field evaporation of graphite discussed in section 3. Such a discrepancy is away from possible errors in measurements and needs a special consideration.

(a)

(b)

(c)

(d)

Figure 20. Snapshots of the dissociation dynamics of a single C-chain extending out from the graphene sheet: (a) at the beginning; (b) after 3.2·10-14 s, (c) after 6.4·10-14 s, and (d) 7.1·10-14 s.

Heating of linear carbon chains produced by unraveling of CNTs was observed in the molecular dynamics experiments [11]. Remarkably, the kinetics of chain breaking varies with strain rate and temperature. The drop in potential energy makes the temperature of the chain rise very sharply. The excess kinetic energy in the region just unraveled then propagates to the graphene, until the whole system again reaches the equilibrium state. The potential energy rises steadily with the tensile deformation until the CNT breaks. The succeeding contraction of the carbon-carbon atomic bonds leads to a sudden drop of potential energy and to a sharp

Preparation and Characterization of Monoatomic Carbon Chains

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increase of the temperature to about 10000 K. Despite the very high temperature, the formation of some short-lived wires was still possible to observe. The fact that the computer simulations demonstrate CNT breaking modes with the presence of carbon atomic wires emphasizes two aspects: that CNT breaking proceeds in a high-temperature region and that the carbon structure is unable of evacuating the heat fast at high rates of local deformation. We note that under the simulation conditions the unraveling graphite layers discussed in section 8.1 is quasiadiabatic and does not occur at constant temperature [60]. At the beginning of the experiment until the bond breaking starts, the graphene sheets are at about 0 K. Figure 21 shows that potential energy of the neighbor atom of the chain (curve 1) and the third nearest atom (curve 2) in the graphene sheet rises gradually with the bond stretching until the first bond breaks up. The time unit is 3.526·10-14 s and the time step is 7.052·10-16 s. The consequent contraction of the atomic bonds causes an abrupt decrease of potential energy and atomic oscillations with the period of 0.028 ps. The average amplitude of oscillations within 0.1 ps is correspondent to temperature of about 10000 K. It indicates that the graphene based carbon structure is incapable of fast evacuating the heat. Therefore, the unraveling proceeds in an ultrahigh-temperature surface region of carbon fibers. The possibility of explosive local overheating above the critical point of carbon was demonstrated earlier in molecular dynamic simulations of the CNT breaking [11]. One can conclude that ions are field evaporated from a graphite surface initially in linear cluster forms, which decompose mostly into smaller atomic clusters and individual ions because of the ultrahigh-temperature excitation during unraveling.

-4 -6

2

Q, eV

-8 -10

1 -12 -14 -16 0

20

40

60

80 100 120 140 160 180

t (time steps)

Figure 21. Potential energy per atom versus time for the MD simulation of the unraveled monoatomic C-chain. Curve 1 is the neighbor atom of the chain, and 2 is the third nearest atom in the graphene sheet.

8.3. Carbon Chains with Macro Length Preparation of long atomic carbon wires is of immense importance for both technological and scientific research. The results of mathematical simulations with the Brenner potential

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showed that free atomic carbon wires with macro length could be obtained by pulling the edge atoms of a graphene sheet with speeds lower than 30 m/s at about 300 K [61]. A monoatomic wire was formed with the carbon atoms that burst forth out of the graphene sheet in a process that resembles unraveling a knitted scarf. The computer experiments indicate that the entire graphite atomic layer is unraveled at about 1000 K to form macroscopically long monoatomic wires. On increasing pulling speed from 70 to 300 m/s the calculated atomic block tended to form shorter chains. More sophisticated tight-binding MD calculations have confirmed these classical simulations based on the bond-order potential. The bond breaking process in graphite with release of atoms to joint the chain was described in details. In the bond breaking process at constant temperature, failure of the chain itself is hardly probable because the energy of the sp bond in a monoatomic wire is larger than that of the sp2 in graphene and the joining bond that connect the carbon chain and the graphene sheet. Wang et al. [61] also performed mathematical simulations without the thermostat. In these experiments, transverse vibrations of a monoatomic wire grow to be more and more intensive during unraveling. This process is accompanied by rising temperature up to about 3000 K and by subsequent failure of the wire. The rupture in the constant pulling speed regime is observed in the middle part of the wire but not at the junction of the wire and the graphene sheet, in contrast to that in the constant force regime (Figure 20). The mechanical loading regimes used in Ref. [61] and that described in section 8.1 are reciprocal. The constant pulling rate can be experimentally realized with ideally stiff testing machines which have two heads, one of which is driven to change the distance between them and thus to impose a definite rate of elongation. A stiff system required little deflection for adjustment of the load and can therefore follow the drop in load required for elongation in the unstable region of unraveling. During the unraveling process, carbon atoms come off the graphene sheet occasionally and join the linear chain. As was shown by Wang et al. the sheet shrinking takes place after an avalanche event of atom bond breaking with a number of atoms joining the carbon chain within a short time. The loading of linear atomic chains by the high electric field in the FIM and our simulations (figures 19 and 21) is the other extreme. The high-field loading discussed in section 8.1 is correspondent to the soft method based on dead-weight loading through a lever. In this case the load cannot fall and remains above the load required for plastic deformation or unraveling of specimens. Moreover mechanical testing in the FIM alone is an ideally soft testing method ensured the absolutely constant pulling force regime. In this regime the unraveling is accompanied by an ultrahigh-temperature surface heating of carbon tips and atomic chains. We can see that molecular dynamic simulations describe the main features of the highfield chain formation. However, further investigation is required to determine how initial atomic-sized faults, which could serve as nanowire nucleus at the graphene edges, emerge. Local oxidation is one of the possible processes of graphite exfoliation to produce graphitic nanoflakes. The details of how oxygen attacks carbon bonds to break up the atomic structure of the graphite sheet were recently understood by Li et al. [62]. They connected oxidation chemistry to the morphology of the graphene sheets and showed how the mechanical stress generated by oxygen-containing groups led to unfolding of the graphite lattice. Epoxy and hydroxyl groups make bridges across carbon atoms at the graphite layer. The oxygen acts mechanistically like a minuscule wedge, pushing apart the bridge’s carbon atoms and stretching the carbon-carbon bonds. As a result, epoxy and hydroxyl groups together induce enough tension in underlying lattice to break the carbon links. These groups mainly attach at

Preparation and Characterization of Monoatomic Carbon Chains

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defects or edge-atom sites in the graphite lattice, where the carbon atoms are not fully bounded. The stress generated by epoxy bridges and structural relaxation around ridges can lead to unraveling of the lattice near the edge sites and to the formation of nanowire nucleus at the graphene edges.

8.4. Young’s Modulus and Tensile Strength of Atomic Chains One can use the interatomic potential to find the ideal tensile strength and the value of resistance to axial stretching, determining nanowire energetics at small strains. The second derivative of energy with respect to axial strain divided by the cross section of a carbon chain S corresponds to Young’s modulus Y. In the Brenner interatomic potential, the cutoff function introduces an artificial increase in the restoring force. As was shown Belytschko et al [63], the extremely large computed values of the CNT failure stress in early theoretical studies were an anomaly due to the cutoff problem typical for Brenner-like potentials. Using the modified Morse potential [63] that closely matches the Brenner potential in the 0–15 % strain range, we calculated Young’s modulus of carbon atomic chains. The special treatment for calculating the chain diameter or the load-bearing cross-sectional area must be carried out for determining the breaking strength and Young’s modulus values. The result Y = 3.68 TPa was obtained with the carbon chain diameter of 2.0 Å [64] equal to the cutoff distance in TersoffBrenner bond order potential. The diameter of carbon atomic chains is not a well-defined parameter but this value seems reasonable. In fact, according to recent first-principles calculations of the electronic structure of molecular wires this value corresponds to the spatial extent of the highest occupied π-states [33]. We have also calculated the critical strain and tensile stress of carbon nanowires.

( ) , where F

The stress σ is defined by σ = 4 Fa / πd

2

a

is the axial force, d is the diameter *

of carbon atomic chains. The strain is given by ε = (r – r )/r*, where r and r* are the current and initial interatomic lengths per atom, respectively. The maximum force is 7.916 nN at a strain of 0.19 and the ideal tensile strength is equal to 252.1 GPa. The ratio σ/E for monoatomic carbon wires is of 0.069. To our knowledge there are no experimental data on mechanical properties of linear carbon chains. Real strengths of solids are orders of magnitude lower than theoretical ones. It is usually observed that the strength of materials depends on size of the sample. The typical strength for microfiber is 1-5 GPa, the strength of nanofibers can reach 12 GPa [65] and graphite whiskers exhibit strength as high as 20 GPa. Carbon nanotubes have extremely high tensile strength approaching 60 GPa and axial Young’s modulus of about 1 TPa [66,67]. This behavior is due to the fact that the number density of defects is significantly reduced in whiskers and nanosized objects. The experimentally measured strength of CNTs and the theoretical results [63] are not in quantitative agreement but are of the same order of magnitude. The nanotube fracture behavior depends on topological defects and vacancies which are always present in the nanotubes used for mechanical testing. Vacancies and vacancy-related defects have a strong effect on the tensile strength and critical strain of CNTs. The tensile strength will significantly drop if single vacancy is present owing to the quasione-dimensional structure of CNT–the strength of low-dimensional objects is governed by the weakest part. Vacancies and other structure defects cannot be present in linear atomic

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chains. Thereupon, we can expect that the real and ideal tensile strength of carbon chains under consideration basically coincide. The values proved to be higher than most results of similar simulations and tensile-loading experiments for carbon nanotubes. Studying the mechanical properties of this unusual form of matter should enable rigorous tests of our understanding of atomic interactions.

9. Conclusion Linear carbon chains have drawn tremendous interest from fields ranging from condensed-matter physics to astronomy because of their unique properties such as nanometric dimensions, quantized conductance, high aspect ratio and modulus of elasticity, unusual optical features and electromagnetic response. Monoatomic chains show transport properties of fundamental and technological interest since they represent the ultimate size limit of functional electronic devices. This interest is also stimulated to a large extent by a desire to understand the atomic-scale mechanisms underlying the preparation of carbon chains. Experiments with the FIM have revealed new mechanisms for low-temperature field evaporation of graphite with important implications for device fabrication. This overview ties together several aspects of recent research on carbon atomic chains and the techniques used, focusing on a high-field technology, field-ion microscopy, field emission, and molecular dynamics simulation. It was shown that carbon chains are perfectly resolved in the field ion microscope. An analysis of the cluster images and determination of the field enhancement factors strongly indicate that the field produced clusters are linear chains of one carbon atom in diameter. The process of field evaporation of graphite at the pulse voltage loading is sporadical with an anomaly large instant rate of evaporation corresponding to explosive moving off about 1011 atomic layers per second. MD simulations demonstrate that electric field pulling a single carbon atom on the graphene edge does not preferentially breaks the carbon-carbon chemical bond. Instead, it is found that this process leads to the formation of a monoatomic carbon chain, followed by breaking a carbon-carbon bond at the graphene surface with a high rupture force. A detailed experimental comparison with theory has not been carried out, although such a comparison is critical for advancing our understanding of these nanoobjects. FIM experiments and MD calculations indicate that carbon exhibit pronounced many-body interactions with strong bonding in low coordinated systems. A universal interpretation of the force-induced unbinding of graphite layers is established. At the moment the stability of carbon chains is one of the main problems to deal with. We are witnessing novel properties of linear carbon chains such as the explosive heating during the high-field unraveling of graphene. Further theoretical and experimental research is still required so that novel technologies will turn into a reality in the near future.

Acknowledgments The author is grateful to A.S. Bakai and N. Wanderka for helpful comments. The author would like to thank V.A. Ksenofontov, T.I. Mazilova, V.A. Sadanov, and O.A. Velicodnaja for their help in preparation of this article, numerous fruitful discussions and collaboration in some of the work reviewed here. The original research of the author discussed in this article

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was supported, in part, by the NATO International Program No. SA (PST.CLG.976376) and the Deutsche Forschungsgemeinschaft (grant 436 UKR 17/26/06).

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[54] Troiani, H.; Miki-Yoshida, M.; Camacho-Bragado, G.A.; Marques, M.A.L.; Rubio, A.; Ascencio, J.A.; Jose-Yacaman, M. Nano Lett. 2003, 3, 751-755. [55] Fantini, C.; Cruz, E.; Jorio, A.; Terrones, M.; Terrones, H.; Van Lier, G,; Charlier, J-C.; Dresselhaus, M.S.; Saito R.; Kim, Y.A.; Hayashi, T.; Muramatsu, H.; Endo, M. and Pimenta, M.A. Phys. Rev. B 2006, 73, 193408-1-193408-5. [56] Abel, G.C. Phys. Rev. B 1985, 31, 6184-6196. [57] Tersoff, J. Phys. Rev. Lett. 1988, 61, 2879-2882. [58] Brenner, D.W. Phys. Rev. B 1990, 42, 9458-9471. [59] Nishikawa, O.; Ohtani, Y.; Maeda, K.; Watanabe, M.; Tanaka, K. Mater. Characterization 2000, 44, 29-57; Nishikawa, O.; Taniguchi, M. Chinese J. Phys. 2005, 43, 111-123. [60] Ksenofontov, V.A.; Mazilova, T.I.; Mikhailovskij, I.M.; Sadanov, V.A.; Velikodnaja, O.A. Kharkov Nanotechnology Assembly, Kharkov, April 2007, 2, 78-83. [61] Wang, Y.; Ning, X.-J.; Lin, Z.-Z.; Li, P.; Zhuang, J. Phys. Rev. B 2007, 76, 165423-1165423-4. [62] Li, J.-L.; Kudin, K.N.; McAllister, M.J.; Prud’homme, R.K.; Aksay, I.A.; Car, R. Phys. Rev. Lett. 2006, 96, 176101-1-176101-5. [63] Belytchko, T.; Xiao, S.P.; Schatz, G.C.; Ruoff, R.S. Phys. Rev. B. 2002, 65, 235430/1235430/8. [64] Wu, G.; Dong, J. Phys. Rev. B 2005, 71, 115410/1-115410/7. [65] Mordkovich, V.Z. Theoretical Foundations of Chemical Engineering 2003, 37, 429438. [66] Shenderova, O.A.; Zhirnov, V.V., Brenner, D.W. Critical Reviews in Solid State and Materials Sciences 2002, 27, 227-356. [67] Terrones, M. Annu. Rev. Mater. Res. 2003, 33, 419-501.

Reviewed by Professor A.S. Bakai, Institute of Theoretical Physics of National Scientific Center “Kharkov Institute of Physics and Technology”.

In: Nanotechnology… Editors: C.J. Dixon and O.W. Curtines, pp. 409-433

ISBN: 978-1-60692-162-3 © 2010 Nova Science Publishers, Inc.

Chapter 12

SEQUENTIAL NUCLEATION AND GROWTH OF COMPLEX NANOSTRUCTURES BY A TWO-STEP STRATEGY Li Yang1,*, Paul W. May1 and Lei Yin2 1

School of Chemistry, University of Bristol, Cantock’s Close, Bristol BS8 1TS, United Kingdom. 2 Department of Aerospace Engineering, University of Bristol, Queen’s Building, University Walk, Bristol BS8 1TR, United Kingdom.

Abstract Self-assembled nanostructures are new forms of materials which are interesting from a fundamental scientific perspective, as well as having many potential technological applications. It is believed that the ability of nanostructures to self-assemble with controlled crystalline orientation, size, complexity and crystal morphology, provide potential applications in data storage, functional devices, communications and technology. Recently, a two-step strategy was successfully developed in our lab to produce two-dimensional or threedimensional carbon nitride well-defined hierarchical complex structures. This strategy is a combination of a novel laser-induced deposition technique followed by self-assembly. In the first step, a suspension of carbon nitride nanoparticles was prepared by liquid-phase pulsed laser ablation (LP-PLA). In the second stage, this suspension was deposited onto a silicon substrate to act as a ‘seed’ layer. Via controlling the rate of evaporation of the liquid phase part of the seed suspension, and the size and the quantity of nanocrystals within the droplet, it was possible to create a range of nanoscale structures, including dense nanospheres, highlysymmetric flowers, hollow core-shell and uniform grass-like structures. The growth of such complex structures is governed by an evaporation-driven self-assembly process. As the droplet dries, small building blocks, such as nanoparticles (NPs) or nanorods (NRs) nucleate upon the existing crystals and template, sharing the same edges, to form a close-packed arrangement. By varying the design of the building blocks, materials combination, interfacial chemistry, and confining dimensions, it is expected to extend this synthetic approach to a range of new structured materials with useful functional properties.

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1. Introduction Self-assembly is an incredibly powerful concept in modern molecular science. The ability of carefully designed building blocks to spontaneously assemble into complex nanostructures underpins developments in a wide range of technologies — from materials science through to molecular biology. Due to the multi-disciplinary nature of this new subject, the definition of a self-assembling process may be different among researchers from different fields. Generally speaking, self-assembly is a process in which components, either separately or linked, spontaneously form ordered aggregates [1]. The interactions involved usually are non-covalent, such as electrostatic interactions, hydrogen bonds, van der Waals’ forces, coordination interactions and hydrophobic effects [2]. In self-assembled structures, these intermolecular forces connect the molecular building blocks in a reversible, controllable and specific way. Of particular value are the possibilities offered by self-assembly to generate nanoscale complexity with relatively little synthetic input. Furthermore, the ability of assembled superstructures to behave as more than the sum of their individual parts, and exhibit completely new types of behaviour, is of special interest [3]. Due to the emergence of a new generation of high-technology materials, two-dimensional (2D) and three-dimensional (3D) self-assembled, aligned nanomaterials have begun to be widely investigated over the last decade [4,5]. It is believed that the ability of nanostructures to self-assemble with controlled crystalline orientation, size, complexity and crystal morphology provide potential applications in data storage, functional devices, communications and technology. As a result, a rapidly growing field of science has emerged to understand and control these self-organized architectures, involving the dedicated efforts of chemists, physicists, material scientists and biologists. In general, physical methods (including chemical vapour deposition [6], vapour phase transport [7], and pulsed laser ablation in vacuum [8]) and chemical methods (including hydrothermal methods [9], softtemplate [10] and use of various surfactants [11,12]) have been developed to fabricate nanoto microscale materials with a range of morphologies, such as compact hexagonal networks [13], rings [14], dandelion-shaped hollow structures [15], strips [16], tubes [17], and flowerlike structures [18]. For example, ZnO forms micron- and submicron-scale ‘dandelion’-like structures, which are comprised of single-crystalline building units (either nanorods (NRs) or nanoparticles (NPs)) [19], and which can be constructed via a modified Kirkendall process in solution, where the pre-formed oxide layer serves as a shell template for the initial nucleation and growth. Also, uniform Sb2S3 bundles coalesced from numerous NRs have been synthesized on a large scale using a hydrothermal technique at a temperature of 180°C for 20 h [20]. However, it is well known that these physical techniques require relatively high temperature, vacuum conditions, and expensive equipment, or sometimes complicated processes, which limit them to smaller scale fabrication. Also, conventional chemical methods usually involve the use of catalysts, surfactants and possibly complex chemical reactions, which often produce a significant amount of unwanted byproducts requiring further purification. Therefore, a technique which combines the merits of both physical and chemical methods, while giving high yield at low cost, would be desirable, if novel self-assembled materials are to be produced on an industrial scale. One such technique, liquid-phase pulsed laser ablation (LP-PLA), has only relatively recently been applied to produce self-organized nanomaterials [21,22,23]. Details about the

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technical aspects of LP-PLA can be seen in a recent review by Yang [24] who gives comprehensive details of the nucleation thermodynamics, the phase transition, and the growth kinetics of nanocrystals by laser ablation of liquids. Briefly, LP-PLA involves focusing a high power laser beam onto the surface of a solid target, which is submerged beneath a liquid. The interaction of the laser with the target causes the surface to vaporise in the form of an ablation plume, which contains species such as atoms, ions, and clusters, travelling with high kinetic energy. The species in the plume collide and react with molecules of the surrounding liquid, producing new compounds containing atoms from both the original target and the liquid. Due the intensity of the laser and the nanosecond timescales, the instantaneous temperatures and pressures within the reaction volume can be extreme (many thousands of K at tens of GPa) [25]. Such high temperature, high pressure, and high density conditions provide a ‘brute force’ method of synthesising novel materials that have hitherto been inaccessible using milder, more conventional techniques. Extensive progress has been made towards the production of many types of inorganic nanoparticles, including metals, metal oxides, and other semiconductors [26]. However, to date, research on the understanding of the self-assembly processes of Group IV-V compounds, such as carbon nitride, is still in its infancy, and synthesis of complex nanostructures has only just begun. Carbon nitride has been the subject of numerous publications since the prediction by Liu and Cohen [27] in 1989 that crystalline C3N4 should have superhard properties. However, a successful synthesis of bulk amounts of this material still remains a challenge. The synthesis difficulties are due to its low thermodynamic stability and complex bonding environment. Recently Li and co-workers [28] demonstrated a range of self-assembled carbon nitride morphologies (including nanotube bundles, aligned nanoribbons and microspheres) could be prepared by a solvothermal technique. Also, our recent findings [29,30] indicated that the instantaneous high temperature, high pressure and high density conditions that arise when a high-intensity focused laser beam impinges upon a graphite target confined by a thin layer of liquid ammonia can promote growth of crystalline carbon nitride NPs. In this chapter, we will further demonstrate a successful synthesis of carbon nitride hierarchical nanostructures via a two-step strategy, whereby the nanocrystals self-assemble into complex 2D or 3D superstructures. Fabrication of these carbon nitride structures from small building blocks via self-organization suggests that the chemical and physical properties of these superstructures are intrinsic to the self-assembly induced by the close vicinity of the NPs or NRs. Finally, a summary and expectation are discussed with regard to the application of LP-PLA as a synthesis route for highly desirable complex architectures.

2. Processing Hierarchically Structured Nanomaterials Production of zero-dimensional (0D) NPs and one-dimensional (1D) nanocrystals by laser ablation in liquid media has been extensively studied [31,32,33]. Since the 0D and 1D nanocrystals can serve as building blocks in forming 2D or 3D complex architectures with long-term periodic structures, it should be expected that the LP-PLA approach would have great potential as a means to grow large arrays of hierarchical, complex, oriented and ordered superstructures. Extending this concept, we performed a two-step synthesis to produce oriented carbon nitride nanostructures based on previous results [Error! Bookmark not

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defined.,Error! Bookmark not defined.]. In a typical synthesis process, the carbon nitride seed solution was initially prepared by LP-PLA. Briefly, a solid graphite target (Testbourne Ltd., 99.99%) was ablated at room temperature while submerged in a 5 ml solution of 25-35% ammonia solution (Fisher Scientific, used as received without further purification) inside a sealed stainless steel cell. In order to reduce the effect of target aging, the cell was rotated at 700 r.p.m. during ablation using a standard magnetic stirrer. A Q-switched Nd:YAG laser (532 nm, pulse duration 15 ns, frequency 10 Hz) was directed by a prism and then focused onto the graphite surface using a 25 mm-focal-length lens. The intense laser light passed through a quartz window in the top of the cell, then through ~5 mm layer of the liquid covering the graphite, to form a ~0.5 mm-diameter spot on the target surface. The ablation was typically carried out at laser fluences of 25-125mJ / pulse for reaction times t = 0.5-24 h. After ablation, a pale yellow colloidal suspension was obtained, which contained a mixture of unreacted graphite and ablation product, both in the form of NPs. The suspension was stable, with no precipitate being observed for months or even longer. The graphite sediments were filtered and removed as much as possible by boiling with 70% perchloric acid, before further characterization. For the second step, a few drops of this seed suspension was deposited onto a 1 cm×1 cm silicon p-(100) substrate and the liquid allowed to dry. Many more complex nanostructures than the simple NPs and NRs can be produced in a controlled fashion by simply altering the drying time and drying method of the suspension of ablated product. Four different drying processes were used in the present study: (1) dry naturally in air; (2) dry in a sealed tube; (3) dry rapidly in an oven or on a hotplate; (4) dry in a critical point dryer (CPD). Those procedures allowed the time taken to evaporate the liquid to be controlled. For materials analysis, a drop of the suspension was deposited onto a silicon p-type (100) substrate or transmission electron microscopy (TEM) grid, and then allowed to dry. The product was characterized using X-ray diffraction (XRD; Bruker-AXS D8 powder diffractometer, Cu Kα radiation), field emission scanning electron microscopy (FE-SEM; JEOL 6300 F), TEM (JEOL 1200 EX, 120 kV), high-resolution (HR) TEM (JEOL 3000F, 300 kV) and energy-dispersive X-ray analysis (EDX). X-ray photoelectron spectroscopy (XPS; Thermo VG Scientific), laser Raman spectroscopy (Renishaw 2000, excitation wavelength 325 nm) and Fourier-transform infrared spectroscopy (FTIR; using a KBr disc) were performed to search for evidence of C–N, C=C and C≡N bonds.

3. Sequential Nucleation and Growth of Complex Nanostructures 3.1 General Structure of the Nanocrystal Self-assembly All reported samples in the present work were first examined with XRD, EDX, Raman and XPS techniques to obtain their crystallographic structure and chemical composition. Experimental details of the procedures can be seen in our previous reports [Error! Bookmark not defined.,Error! Bookmark not defined.,Error! Bookmark not defined.]. Studies performed at various conditions showed that factors such as ablation time, laser energy, ammonia concentration and the drying speed were important in order to obtain organized assemblies of NPs. By depositing the carbon nitride seed solution onto a silicon substrate and controlling the drying process under different conditions, four main classes of structure were identified in the ablation product, categorised as Types I-IV based upon their

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shape and size (Figure 1(a)). The first class of structure (Type I) had the shape of thin plates with rounded edges. Since they were the components of the larger ‘flower-like’ structures (described later), they have been termed ‘nano-petals’. The quantity of these nano-petals and their location with respect to the larger structures (see later) were dependent upon the deposition and drying conditions. As shown in Figure 1(b-c), this indicated that these nanopetals were 2D crystallites of carbon nitride which preferentially aligned themselves perpendicularly to the surface of the Si substrate. The number and length of these nano-petals increased with increasing laser ablation time from 0.5 to 2 h for the same laser fluence. X-ray diffraction (XRD) analysis (spectrum not shown here) of these nano-petals showed that they were crystalline, and all the diffraction peaks were consistent with (h00) preferential orientation [Error! Bookmark not defined.]. The crystallographic information was indexed to hexagonal β-C3N4 (P63/m (176)) with lattice constants a0 = 6.4017 Å and c0 = 2.4041 Å [34]. Interestingly, it is also possible that these nano-petals began to aggregate and selfassemble (Figure 1(d)). When the concentration

Figure 1. (a) Schematic illustration of the growth process leading to the observed four main classes of hierarchical structures, labelled I to IV. (b)-(d) Scanning electron microscopy (SEM) images of carbon nitride ‘nano-petals’ following ablation times of: (b) 0.5 h, (c) 2 h, and (d) 3 h. (e-f) Overall ‘flowerlike’ structure following 5 h laser irradiation (synthesis conditions: laser power 125 mJ, 35% ammonia solution, drying in air). (g-h) 2D flattened flower, sample conditions identical to (e-f) except it was dried on a hotplate at 200ºC (g) and an oven 80ºC (h). (Reproduced from [Error! Bookmark not

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defined.] by the permission of Royal Society of Chemistry (RSC) and [Error! Bookmark not defined.] by the permission of Nova publishers).

of nano-petals in the suspension increased, they tried to minimize their interfacial energy upon subsequent drying of the liquid by preferential tilting with respect to each other. This produced the second class of nanostructure (Type II), which has a ‘grass-like’ shape and exhibits several different morphologies (see later discussion). However, all are produced in large quantities and cover the whole substrate. By carefully controlling the drying process, ‘flower-like’ spiked, crystalline superstructures were formed (Figure 1(e-f)). This third class of structure (Type III), now fully 3-dimensional, with sizes 1-20 µm, were created when many nano-petal structures coalesced at a common centre with multi-fold symmetry. One possible explanation is that the presence of the solid substrate physically hinders growth in that direction; so many branches are tilted away from the substrate, towards the solution. When the evaporation speed of the liquid was rapid (for example, drying in an oven or hotplate), a fourth class of structure (Type IV) was observed (Figure 1(g-h)). Instead of 3D flowers, the carbon nitride now formed 2D ‘star-like’ or flattened flower-like structures. New dendrites emanated from the core and acted as nucleation centres, eventually allowing the structure to expand into 2D horizontal flowers (Figure 1(h)). It is suggested that the higher water evaporation rate increased the interparticle capillary forces [35]. As the continuous flux of particles filled up the spaces on the substrate, successive layers would be formed rather than 3D complex shapes. Although these four types of structures have different densities and morphologies, they all exhibit high surface-to-volume ratios and so might have potential in semiconductor devices, anticorrosion protective coatings and new applications. By controlling the solution evaporation rate on a carbon-coated TEM copper grid (as our substrate), different final complex architectures could be also achieved on various length scales. If the droplet dried in air (less than 1 h), the structure had the shape of nearly monodispersed spheres with rounded edges (Figure 2(a)). Those spheres have very high density and are close-packed. When increasing the laser fluence and ablation time, the nanostructure resembles ‘flower-like’ spiked crystallites (Figure 2(b)) for the same drying speed, where the NRs have coalesced at a common centre with multi-fold symmetry. It seems that these flower-like structures are simply less dense versions of the spheres (Figure 2(c)), with numerous voids between the NR framework comprising the flowers. Similar interior space would be eventually generated within the solids, as the smaller NPs are undergoing mass transport throughout the prolonging drying times (Figure 2(d)). The solid carbon nitride flower (similar to that in Figure 2(b)) has become hollow (named ‘hollow-flowers’), which divided the pristine solid sphere into two discrete regions and formed a core-shell structure. In these structures, the NPs remain loosely attached to the outside, forming the interconnected void space, although the interior vacant volume might vary in each individual hollow-flower (for instance the interior void space size is in order: 1st >2nd >3rd >4th, as shown colour contrast for two discrete regions of individual hollow-flower in Figure 2(d)). It is worth pointing out that by simply altering the drying time and drying method of the suspension, the final morphology can exhibit semi-hollow, core-shell, or even full-hollow structures. These observations are similar to the report by Liu and Zeng [36], who demonstrated the fabrication of ZnS nanomaterials with hollow interiors. Yang and Zeng also reported a simple approach

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to prepare hollow TiO2 nanospheres with highly organized crystallites in the shell structure and surface regions [Error! Bookmark not defined.]. Compared to the flower-like structures, the hollow-flowers are less common. A longer drying time in an above-ambient temperature aids the small crystallites to move freely in the solution enabling oriented aggregation. The relative ratios the nanosphere, nanoflower and hollow-flower structures are approximately 30%, 60%, and 10%, respectively (based on a total of 800 TEM images), noting that the statistical ratios vary for different experimental conditions. In general, a higher synthetic energy (i.e. high laser power, long ablation time, high ammonia concentration and long drying time) leads to better crystallinity, but larger size for the products.

Figure 2. (a) Nanospheres with rounded edges (synthesis conditions: laser power 50 mJ, ablation time 2 h, 35% ammonia solution, drying in air). (b-c) Nanoflowers with numerous protruding spiky surfaces (synthesis conditions: laser power 100 mJ, ablation time 8 h, 35% ammonia solution, drying in air). (d) Hollow-flowers with tunnels connecting from the centre to the outward shell (synthesis conditions: laser power 100 mJ, ablation time 10 h, 35% ammonia solution, drying in a sealed tube). The numbers denote the individual hollow-flower. (e) HR-TEM image of a single NR, the inset shows the atomic arrangement and scale bar 1 nm. (f) The Fourier-Filtered Transform (FFT) pattern corresponding to the

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region shown in (e). K, J, H, G, P and P' label the positions of the six domains (see text for details). (Reproduced from [Error! Bookmark not defined.] by the permission of Institute of Physics (IOP)).

In principle, assembly is energetically favoured because the formation of larger crystals can greatly reduce the interfacial energy of isolated NPs or NRs. Therefore, most of our final products are either well-defined nanosphere structures or self-assembled flower-like architectures if the TEM grid dried naturally in air. Only occasionally were monodispersed NPs or isolated NRs seen by TEM. The mechanism for how these lower dimensional building blocks construct to higher dimensional arrangements, rather than random clumps, is still unclear. However, the above evidence indicates that ‘oriented attachment’ [37] was observed among the crystallites, in which a larger crystal structure is formed from smaller ones by direct joining of suitable crystal planes. In particular, in the HRTEM image of Figure 2(e) taken from a single NR, the periodic lattices show the atomic arrangement (Figure 2(e), inset) with very few defects, and reflect the relationship between the orientation of the NPs and the crystallography of the ordered NR array. The corresponding Fourier-Filtered Transform (FFT) pattern illustrates that the nanocrystal consists of six domains with sixfold twins. The split spots K and J, H and G in the FFT pattern is due to the small angle between the twin boundaries. The reflections P and P' in Figure 2(f) are not split, and represent the coherent positions of the twin boundaries.

Figure 3. FTIR spectra of three different carbon nitride structures: (a) nanosphere, (b) nanoflower, (c) hollow-flower. (Reproduced from [Error! Bookmark not defined.] by the permission of Institute of Physics (IOP)).

Information regarding the chemical bonding structure was obtained from Fourier transform infrared spectroscopy (FTIR). Figure 3 shows the FTIR spectra of the abovementioned different morphology of carbon nitride crystals, which exhibits several peaks related to the chemical bonding between carbon and nitrogen [38]. The region 1000-1500 cm1 corresponds to C-N single bonding, while the regions 1500-1750 cm-1 and 2150-2300 cm-1 are related to C=N and C≡N bonding, respectively. Similar spectra were also obtained by

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Zimmerman and co-workers [39]. In our case, for the nanosphere structures (Figure 3(a)) the two peaks at 1040 cm-1 and 1406 cm-1 correspond to the C-N stretching mode. The absorption band at 1639 cm-1 is assigned to the stretching vibrational modes of C=N. Moreover, a small peak at 2208 cm-1 can probably be attributed to C≡N bonds, although it is much weaker compared with the other stretching modes. A broad band centred at 3464 cm-1 is due to NH group vibrations [40]. The 558 cm-1 peak can be assigned to the out-of-plane bending mode for graphite-like sp2 domains, which become IR active due to nitrogen incorporation into the bonding network. The obvious differences between these samples are that the intensities of all the features are weaker and broader for the nanoflower and hollow-flower shape nanostructures. We attributed this to the voids in the nanoflowers and the internal nanospaces existing in the hollow-flowers.

3.2. Control of the Quality of Self-assembly 3.2.1. Diffusion of the Building Blocks Usually, ordered organization was obtained following evaporation of a drop of a nanocrystal solution that had been deposited onto a TEM grid. Figure 4(a) shows a typical symmetric carbon nitride ‘flower’ together with its nano-petal building blocks (which look like flattened rods in the TEM). Figure 4(b) shows that these nano-petals appear fused together and ‘interwoven’ to form a lattice-like framework of the flower-like superstructure. The figure also shows the NPs that surround each nano-petal, and which fill in the holes within the framework to produce a dense, solid structure. EDX analysis confirms that carbon and nitrogen are present in all these structures, and micro-diffraction pattern (MDP, Figure 4(c)) was also consistent with crystalline hexagonal β-phase carbon nitride oriented along the [001] zone-axis. Several [001] patterns in Figure 4(c) can be identified at the same time, indicating that the nano-petals consist of several domains, with different rotational orientations contributing to the diffraction pattern. The HR-TEM image in Figure 4(d) shows that the nano-petals at the very edge of the flower contain very few defects and are single crystalline, as was anticipated from the MDP. Again, the lattice fringes (d200 = 0.28 nm, d14 0 = 0.15 nm) and their angles (106°) are in good agreement with the calculated values for hexagonal β-C3N4 [41]. The smaller NPs, which lie in and around the nano-petal-framework comprising the flowers, appeared to be mobile with respect to this framework, and diffused outward from the centre of the flower with longer drying times. The results of this diffusion can be seen in Figure 4(e), where the solid carbon nitride flower (similar to that in Figure 4(a)) has become hollow. The NPs have diffused from the centre but remain loosely attached to the outside, making the outer shell of the flower appear fuzzy. The thickness of the fuzzy shell was ~ 140 nm and that of the hollow core was ~ 200 nm (shown as a lighter colour in the image). When the suspension was placed onto a hot-plate at 200ºC for 0.5 h, the hollowing process was accelerated to form a semi-core-shell structure. The radial distribution of the NRs formed channels leading from the centre to the shell (Figure 4(f)). Another type of hollow structure was observed when the core-shell structures were essentially separated by a hollow tunnel, without linkage to the sphere (Figure 4(g)).

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Figure 4. TEM image of the flower-like structures produced by LP-PLA of a graphite target in 35% ammonia solution (laser fluence 125 mJ/pulse) for 5 h. (b) A higher magnification image of the framed region in (a), showing a high density of NPs surrounding the nano-petal framework. (c) [001] zone-axis MDP from the tips of the nano-petals in (b), which corresponds closely to the calculated interlayer dspacing of β-C3N4. Arrows point to different sets of [001] reflections (see text). (d) HR-TEM image recorded from the edge of the flower nanostructure that is oriented along [001]. (e) TEM image of a hollow flower formed after 8 h LP-PLA and prolonged drying. (f) Semi core-shell structure. (g) coreshell structure with a hollow tunnel. (Reproduced from [Error! Bookmark not defined.] by the permission of Royal Society of Chemistry (RSC)).

§ 3.2.2. Interconnection between the Big Structures Self-assembly normally occurs when nanoscale objects interact with one another through a balance of attractive and repulsive interactions. If the attractive force is dominant, the components may interconnect and form larger aggregates. In contrast, the exterior appearance of the interlinked nanostructures did not change appreciably when the interaction was weak. It was found that the interconnections among the structures were different. Figure 5 shows two examples of these cluster-like combinations, for nanospheres and nanoflowers. Figure 5(a) shows a single, typical nanosphere with a dense surface, similar to those mentioned previously in Figure 2(a). Figure 5(b-c) shows that two neighbouring spheres can be fused together, or just joined loosely via their outside edges, with no change of internal structure. However, for the nanoflowers, the situation was different. Figure 5(e) shows that when two nanoflowers fuse, the inner spaces between the NRs were integrated throughout the entire cluster structure. One important note is that the fusing together of pristine nanoflowers alters the crystalline size slightly (~800 nm in diameter) but the morphology of nanoflower remained (Figure 5(f)).

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Figure 5. Examples of coalescence (eventually leading to clustering) of dense carbon nitride nanospheres and porous nanoflowers. (a) A single nanosphere, (b) 2 nanospheres touching, (c) 2 nanospheres fused together. (d) A single nanoflower, (e) 2 nanoflowers partially fused together, (f) 2 nanoflowers completely fused.

Figure 6. (a-b) nanospheres (carbon nitride seed solution is identical to that used for Figure 2(a)). The fused nanospheres are marked by arrows. (c-d) nanoflowers (carbon nitride seed solution is identical to Figure 2(b-c)). The drying process was in air. The arrow in (d) shows possible fused nanoflowers. Note that image (b) and (d) are recorded in higher magnification from (a) and (c), respectively.

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If the TEM grid was replaced by a silicon substrate, and the previously prepared carbon nitride seed suspension was applied and dried naturally in air, similar morphologies to those shown in Figure 5 were obtained. However, the nanospheres were now formed with a more uniform size distribution, as shown in Figure 6(a-b). Some nanospheres (size about 500 nm1 µm) were attached with a boundary visible in between (highlighted by arrows for clarity). The interaction between such interconnected nanospheres is weak since they could be broken up with a few minutes’ of sonication. In contrast, the nanoflowers were dispersed individually (Figure 6(c)) with a larger size distribution (~1-15 µm). It is noted that the size of the nanoflowers found by SEM on the Si substrate appeared much larger than those under TEM observation (size about 500 nm-800 nm). This is probably due to some of the nanoflowers fusing together completely to produce larger, interlinked nanoflowers (see Figure 6(d), highlighted by the arrow). Such nanoflowers are stable and sustainable without any change even after a few hours’ of sonication. The reason the nanoflowers fused more completely was to do with the different interaction between the particles within the seed suspension and the two types of substrate. The TEM grid was normally placed onto a filter paper and a drop of carbon nitride aqueous suspension was pipetted onto the grid. Due to the presence of the filter paper, water was quickly removed. For the SEM sample on a Si substrate, the droplet of suspension took at least 2 h to dry in air. This allowed sufficient time for the nanoflowers to aggregate. Further investigation needs to be carried out.

Figure 7. Fused-flower (carbon nitride seed solution is identical to that used in Figure 2(b-c)). Drying process is in a sealed tube (~ 8 h evaporation). See text for discussion. Arrows in (c) and (d) are marked for clarity.

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Control experiments provide some evidence to support this proposed mechanism. Under identical conditions, a droplet of carbon nitride suspension was deposited onto a Si substrate and placed into a sealed tube. The drying process was now estimated to be around 8 h. As expected, larger flowers constructed from hundreds of thin plates were formed (Figure 7(a)). Careful observation found that these larger flowers were actually an aggregate of more than one nanoflower. The boundary between the nanoflowers was perfectly fused and aligned. The enlarged region in Figure 7(a) marked by black box shows that numerous nanoplates are bunched up and cross-linked with recognizable boundaries or voids between the component subunits (Figure 7(b)), which are still maintaining their close proximity. Such results can be further seen in Figure 7(c-d). Occasionally, flowers were not fully developed, since the rate of aggregation among the nanoflowers varies depending upon local conditions. But the arrows marked in Figure 7(c-d) clearly indicate the fusion of the adjacent nanoplates. Given sufficient time (e.g. by prolonging the drying time), the growth of the flower will fully complete into three dimensions.

Figure 8. (a-b) 3D interconnected flowers. (c-d) 2D flattened flowers (synthesis conditions: 35% ammonia solution, laser power 100 mJ, 3 h ablation and dried in a sealed tube, (c-d) seed solution was diluted by using deionised water). Numbers in (a-d) denote the individual flower.

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A few more examples of interconnection between larger complex structures are given in Figure 8. As shown in Figure 8(a), the neighboring ‘long-petalled’ crystals can further expand themselves and eventually merge into an interconnected structure [42]. At least 10 petals are needed to form the secondary flower-like architectures. Some of them sprouted out to three dimensions but some were lying down along the substrate (see 3rd flower in Figure 8(b)). However, when using diluted carbon nitride seed suspension, the morphology changes (Figure 8(c-d)) in the mesoscopic structure may be correlated to interactions between nanocrystals. The interconnection between those flowers was also observed, although the forces seemed weaker (larger voids exist, see the numbered regions in Figure 8(c-d) for demonstration) compared with those shown in Figure 7. The disappearance of the 3D structural dimension and the occurrence of 2D flattened flowers which cover the substrate are believed to be a result of a significant decrease in the interaction between the dispersed nanocrystals in the droplets. In a slow evaporation, the attractive force between NPs is low and they are likely to diffuse close to the substrate leading to formation of 2D flattened structures.

3.3. Dynamic Study of Self-assembly Formation 3.3.1. Influence of the Drying Time Morphology development of the high-order grass-like architectures at different growth stages was monitored by SEM (Figure 9). The evaporation rate was controlled by either the substrate temperature or the solvent saturation degree of the surrounding atmosphere. The seed suspension used was spherical NPs with an average particle size of 15-20 nm under TEM observation (see Figure 9(a) inset). When a droplet of this suspension was deposited onto a silicon substrate and placed on a hot-plate, the solution dried quickly (about 30 min). In this case, the NP morphology remained, and small islands of NP aggregates dispersed on the surface sparsely (Figure 9(a)). In contrast, if the droplet on the Si substrate dried naturally in air (~2 h), 1D NR nuclei started branching on the surface and gradually formed 2D ‘roots’ (~200 nm in size), see Figure 9(b). If instead, the droplet dried inside a sealed tube (~8 h), the number of nanopetals increased and started interconnecting or aggregating (see the left side of Figure 9(c)). The 2D primary nanopetals took about 12 h to coalesce into grass-like structures (Figure 9(d)). Upon further increasing the drying time to 24 h, the NPs adjusted their position and continued to assemble, stem-like, and eventually expanded into fully-developed 3D architectures (Figure 9(e)). In the enlarged image of Figure 9(f), it can be seen that nanopetals on each side stem were nearly parallel to one another. Moreover, NR bundles can be clearly observed in each nanopetal (see the highlighted area in Figure 9(f)). It should be noted that such heterogeneous nucleation and growth in solution were not observed by means of TEM. Similar experiments were also performed using the carbon nitride seed suspension produced by 35% ammonia solution, laser power 50 mJ, 12 h ablation. This produced nanoflowers which exhibited surfaces composed of NRs, where the NRs radiated outward from the centre, as shown in Figure 10(a) inset. The starting seed morphology was more complex than the NPs used in the previous time evolution study. Since it was observed that nanoflowers exhibited a surface composed of NRs, we expect that a similar morphology

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Figure 9. Time-dependent evolution of grass-like crystal morphology at different growth stages for (a) 30 min, (b) 2 h, (c) 8 h, (d) 12 h, and (e-f) 24 h, respectively. The box in (f) highlights the nanopetals arrangement. (Synthesis conditions: 35% ammonia solution, laser power 75 mJ, 10 min ablation, spherical NPs with an average particle size of 15-20 nm were used as seeds, as shown in Figure 9(a) inset. See text for discussion).

should appear on the Si substrate. However, no regular shaped objects were detected by SEM when the droplet was dried in air. Only some low contrast flower-shape patterns (Figure 10(a)) were seen which appeared uniformly dispersed. The enlarged image in Figure 10(b) shows that such patterns are actually formed by a number of particles with preferential arrangements, which were presumably the later nucleation sites for the growth of the flower. With controlled drying time (~8 h) in a sealed tube, large numbers of flower-like structures

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(Figure 10(c)) about 500 nm in size were distributed on the silicon surface. The enlarged image Figure 10(d) from the area of Figure 10(c) shows these nano-objects possess external bonding capacity or ‘adhesiveness’ for self-assembly and self-alignment. For example, they have extending ‘arms’ (NRs or nanopetals) ideal for external connectivity (highlighted by arrows in Figure 10(d)). Fully complex geometric flowers took more than one day to develop (Figure 10(e)). Such structures comprise numerous 1D NRs or 2D nanopetals with their longaxis pointing toward the centre of each flower. Those components were arranged side-by-side and some were tightly bonded each other (highlighted by arrows in Figure 10(f)).

Figure 10. Time-dependent evolution of the flower-like crystal morphology at different growth stages for (a-b) 2 h, (c-d) 8 h and (e-f) 24 h, respectively. Arrows in (d) and (f) highlight the bonding components. Note that images (b), (d) and (f) are recorded at higher magnification than (a), (c) and (e), respectively. (Synthesis conditions: 35% ammonia solution, laser power 50 mJ, 12 h ablation, a flower exhibited a surface composed of NRs, where the NRs radiate outward from the centre as shown in Figure 10(a) inset. See text for discussion).

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From the above discussion, it seems that the morphology of the initial stage of complex architectures (grass-like or flower like) was a small aggregate of compacted 15-20 nm-sized dense particles, which were approximately spherical in shape. Although the nucleation and growth of grass-like or flower-like carbon nitride superstructures might accompany NP fusion and self-alignment to bigger building blocks (such as NRs or nanopetals), no direct evidence was observed for this direct aggregation. Fusion and self-organization may also be associated with thermal motion of the droplets, for example, the evaporation speed, which is determined by the temperature and the flow rate of air above the surface. However, what is the driving force that controls the final hierarchical complex? The mechanism is still not fully understood. One possibility is that growth continues within the local fluid environment as long as there is substantial mobility of the seed particles for further exchange. If the process is sufficiently slow, this will continue until all the building blocks adjust to their desired (lowest energy) position. Evidence for this is that the fully developed grass-like or flower-like complexity was only observed on decreasing the evaporating rate, i.e. for drying times above 12 h in the later growth stage (Figure 9(e-f) and Figure 10(e-f)).

3.3.2. Influence of the Starting Seed Solution

Figure 11. TEM images obtained by LP-PLA in 25% ammonia solution: (a) isolated carbon nitride NRs (50 mJ/ pulse, t = 1 h) (b) Branched NRs (50 mJ/ pulse, t = 3 h) (c) highly branched flowerlike architectures (100 mJ/ pulse, t = 12 h). (d) Rod-like structures showing straight, long and sharp tips. (e) Enlarged image of the region at the top of the NRs indicated by the open box in (d). (f) HRTEM image of a single NR, the inset shows the atomic arrangement, scale bar 1 nm. (Reproduced from [Error! Bookmark not defined.] with permission of Nova publishers).

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When 35% ammonia was replaced by 25% ammonia solution, carbon nitride was still formed; however, the morphology was different under TEM observation. For low laser power (50 mJ) and short ablation time (1 h), the product contained mostly a sparse collection of isolated NRs (Figure 11(a)). With longer ablation times (3 h), the NRs started to aggregate into branched structures (Figure 11(b)), serving as the starting points (or nucleation seeds) for the subsequent growth. With increasing laser power (100 mJ) and ablation time (12 h), the concentration of NRs increased, and highly-branched flower-like architectures completely composed of NRs can be observed (Figure 11(c)). The NRs at the edge of the flower structures appeared to be protruding outward by ~10 nm. As shown in Figure 11(d-e), these NRs are themselves composed of a large number of smaller NPs that have packed together in an ordered arrangement to form the rod-like shapes. In particular, in the HRTEM image in Figure 11(f) taken from a single NR, the periodic lattices clearly show the atomic arrangement (Figure 11(f) inset) with no visible boundaries, and reflect the well-aligned the orientation between the NPs and the ordered NR array.

Figure 12. SEM images of different rod-like patterns obtained by drying the carbon nitride colloidal solution (a) in air (synthesis conditions: 25% ammonia solution, 50 mJ/ pulse, t = 1 h, TEM image shows the isolated carbon nitride NRs (see Figure 11(a)). (b) in air and (c) in sealed tube (synthesis conditions: 25% ammonia solution, 50 mJ/ pulse, t = 3 h, TEM image shows branched NRs (Figure 11(b)). Image (d) is recorded at a different magnification to (c).

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The starting seed suspension can strongly influence the 2D and 3D arrangements of the nanostructures. Using the carbon nitride seed suspension prepared from ablating 25% ammonia solution, and varying the concentration and liquid evaporation rate, rod nanocrystal formation was been observed with different morphologies (Figure 12). These include irregular disordered rods (Figure 12(a)), rods with branches (Figure 12(b)), and celery-like structure (Figure 12(c-d)). One thing these morphologies have in common is that they are all originally made from rod-shaped seed particles.

Figure 13. SEM images of carbon nitride ‘grass-like’ structures following ablation times (a) 0.5 h, (bd) 2 h (synthesis conditions: laser power 125 mJ, 35% ammonia solution, and drying process: (a-b) drying in air and (c-d) drying in a sealed tube). Note that image (d) was recorded at high magnification than (c). Arrows in (a) point out the vacancies and cracks within the film.

We found that the size of the NPs in the initial starting suspension is also very important in determining the morphology of the carbon nitride hierarchical nanostructures. If the starting seeds were short rods (Figure 13(a)), this favoured the sticking of the NRs onto the substrate. The vacancies or cracks observed in Figure 13(a) were probably due to the fact that not enough nanocrystals were deposited on the substrate. However, if the seed solution contained sufficient number of nanocrystals with longer (about 500 nm) length, rod-rod interaction dominated and they tended to form a denser and rougher grass-like surface (Figure 13(b)). The tips of those rods were very blurred, which implies that the solvent-substrate

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interaction became unstable. With longer drying time, an identical seed suspension coated on the substrate formed a homogenous grass-like film (Figure 13(c)) with plenty of radial NRs (50-200 nm in width, 1-5 μm in length). It seems that in a slow evaporation process, increasing the size and the number of the NRs is beneficial to improve the crystallinity of the individual NRs and the uniformity of the film (Figure 13(d)).

Figure 14. (a-b) TEM images of NR building blocks that form the ‘grass-like’ structures shown in Figure 13. (Synthesis conditions: laser power 125 mJ, 35% ammonia solution, 2 h ablation times, and drying process: (a) drying in air and (b) drying in a sealed tube). (c-f) HRTEM images recorded from the rods. (c) and (e) correspond to the region in (a), while (d) and (f) correspond to the region in (b). Arrows in (a-d) point to the growth direction of the NRs. (g) Microdiffraction pattern of corresponding (f) HRTEM image.

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The crystallinity improvement related to the atomic and nanocrystal ordering was further observed by TEM and HRTEM. Under identical synthesis conditions, but using different drying methods, Figure 14(a) and (b) show apparent differences between those NRs subunits corresponding to Figure 13(b) and (c), respectively. When the solution is immediately dried in air (with the help of filter paper), the NRs were parallel-aligned and were 50-80 nm width and had a short length. The outline of the NRs was very blurred (Figure 14(a)). However, when drying was very slow, the NRs were very sharp, had longer length and larger width (~80-200 nm). Some NRs were overlapped and tilted away, similar to those shown in Figure 13(c-d). These results support the previous SEM observations. From the HRTEM images shown in Figure 14(c) and (d) (which correspond to the rod area in Figure 14(a) and (b)), it was found that the crystallinity within the domain was different. However, the growth direction for those rods perfectly matched the orientation that appeared in the TEM images (highlighted by arrows in Figure 14(a) and (b)). The enlarged images of Figure 14(e) and (f) show the individual atoms and their arrangements in this small region. Figure 14(e) consists of two slightly distorted layers, which are superimposed with respect to each other, and joined in the [101] direction. Consequently, the C3N4 groups deviate slightly from a planar arrangement. In contrast, Figure 14(f) is more close to the idealized β-C3N4 structure. The atoms are linked with one edge parallel to, and one edge perpendicular to, the (001) plane, with regular ordering [43]. The micro-diffraction pattern (Figure 14(g)) from the self-aligned NPs (Figure 14(f)) shows the single-crystal type [001] zone-axis pattern. Thus, the orientation order among the NPs in the assembly was further confirmed.

4. General Discussion about Self-assembly Mechanism Our results clearly indicate that the formation of 2D or 3D hierarchical complex architectures is evaporation-driven self-assembly [44]. Evaporation-driven self-assembly is one of the most promising techniques for practical use [45], because it is inexpensive, has a high throughput, and it is a suitable technique for both low-dimensional assemblies and longrange-ordered complex structures. Various factors in our system, such as the rate of evaporation, the starting seed suspension, and the size and the quantity of nanocrystals within the droplet, are very important to determine a well-defined self-assembly. The formation of fully 3D carbon nitride structures was a slow process. During the evaporation process, capillary forces arise. At the edge of the meniscus, the thickness of a liquid film becomes small. The meniscus between particles has an unstable form in a thin liquid film (Figure 15). There are strong attractive interactions among the colloidal particles because of this instability. The attractive capillary force is called the ‘lateral capillary force’. When the liquid evaporation rate is high, the liquid from the bulk suspension flows to the edge of the meniscus, and other particles in the suspension are driven toward this nucleus by the resulting convective transport (Figure 16). As in our system, the droplet was composed of carbon nitride solid objects floating at the silicon substrate interface, which interact by lateral capillary forces. Such forces might direct the patterning of the wettability of the surfaces via self-assembly minimization of the interfacial free energy of the liquid-liquid interface [46]. Clearly, the morphology transformation process of NPs to hierarchical architectures requires a significant degree of mobility in the local environment [47], which could explain why the

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well-defined flower-like or grass-like structures are only formed during a slow drying process. Similarly, the diffusion of the building blocks and the interconnection between the larger structures are also related to such motion of the liquid flux. A high concentration of starting seed suspension containing more components within the droplet may induce more interaction between the nanocrystals and thereby increase the chance for the self-organization of aggregates in the later stages.

Figure 15. Schematic diagram of the lateral capillary force caused by an unstable meniscus. (Reproduced from [48] with permission of Wiley interscience).

Figure 16. Schematic diagram of evaporation-driven self-assembly.

In self-assembly, the molecular structure determines the structure of the assembly [49]. Therefore, it is believed that the self-assembly into functional structures can be fabricated through design of the various components. Our results in section 3.3.2 have shown that the different shape of the starting components leads to the formation of different structurally defined aggregates. For example, the rod-like seeds generated rod-branches (Figure 12) or grass-like structures (Figure 13) depending on the conditions, where the original shape of the components remained intact. The more complicated the shape, the more difficult it is to predict the final form of any aggregate. Further in-depth understanding of such phenomenon is required in future research.

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5. Conclusions The formation and morphological evolution of 2D or 3D carbon nitride hierarchical complex structures was studied in this chapter. A two-step strategy was developed to control the final well-defined superstructures. In the first step, carbon nitride seed suspensions were prepared by liquid-phase pulsed laser ablation. In the second stage the chosen seed suspension was deposited onto a silicon substrate. Via controlling the rate of evaporation, the starting seed suspension, and the size and the quantity of nanocrystals within the droplet, it was possible to create dense nanospheres, highly-symmetric flowers, and uniform grass-like structures, respectively. Because of the mobility in the local environment, the smaller building blocks (NPs) appeared to be mobile with respect to the larger structures (flowers), and diffused outward from the centre with longer drying times. Such diffusion created different degrees of hollow structure, including core-shell, semi-hollow, and hollow organization, which may find applications in new technological areas. In addition, interconnection between the larger structures was also observed for slow evaporation processes. For dense nanospheres, the fusing of the pristine nanospheres was weak and they were easily separated. However, the nanoflowers with many protruding NRs had sufficient capacity for the component NRs to occupy the spaces within the nanoflowers and fuse the flowers together more strongly. From studies of the dynamics of self-assembly formation, it seems that the morphology of the initial stage of the complex architectures (grass-like or flower-like) was small aggregates of compacted 15-20 nm-sized dense spherical particles. The self-organization in the later stage may be associated with thermal motion of the droplets, for example, the evaporation speed. Although the mechanism is still not fully understood yet, our experiments which compared drying time and starting seed suspension indicated that self-assembly into functional structures can be achieved through the design of the various components. A slow drying process will favour an increase in the structural complexity. The formation of 2D or 3D hierarchical complex architectures in this work seems to be an evaporation-driven self-assembly process. The beauty of this approach is its simplicity and efficiency. If combined readily with micrometre-scale film patterning strategies, such as microcontact printing [50], direct writing, or lithography, it may provide a convenient pathway to the formation of functional hierarchical devices. By varying the design of the building blocks, materials combination, interfacial chemistry, and confining dimensions, we should expect to discover new materials properties.

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Li Yang, Paul W. May and Lei Yin Kempa, K.; Kimball, B.; Rybczynski, J.; Huang, Z. P.; Wu, P. F.; Steeves, D.; Sennett, M.; Giersig, M.; Rao, D. V. G. L. N.; Carnahan, D. L.; Wang, D. Z.; Lao, J. Y.; Li, W. Z.; Ren, Z. F. Nano Lett. 2003, 3, 13-18. Duan, X. F.; Lieber, C. M. Adv. Mater., 2000, 12, 298-302. Sun, Y.; Fuge, G. M.; Fox, N. A.; Riley, D. J.; Ashfold, M. N. R. Adv. Mater., 2005, 17, 2477-2481. Yang, H. G.; Zeng, H. C. J. Phys. Chem. B, 2004, 108, 3492-3495. Mirkin, C. A.; Letsinger, R. L.; Mucic, R. C.; Storhoff, J. J. Nature, 1996, 382, 607609. Sun, Y. G.; Gates, B.; Mayers, B.; Xia, Y. N. Nano Lett., 2002, 2, 165-168. Li, Y. D.; Li, X. L.; Deng, Z. X.; Zhou, B. C.; Fan, S. S.; Wang, J. W.; Sun, X. M. Angew. Chem. Int. Ed., 2002, 41, 333-335. Motte, L; Billoudet, F.; Pileni, M. P. J. Phys. Chem. 1995, 99, 16425-16429. Maillard, M.; Motte, L.; Pileni, M. P. Adv. Mater. 2001, 13, 200-204. Liu, B.; Zeng, H. C. J. Am. Chem. Soc., 2004, 126, 8124-8125. Petit, C.; Legrand, J.; Russier, V.; Pileni, M. P. J. Appl. Phys., 2002, 91, 1502-1508. Ngo, A. T.; Pileni, M. P. Adv. Mater., 2000, 12, 276-279. Liu, J. P.; Huang, X. T.; Li, Y. Y.; Sulieman, K. M.; He, X.; Sun, F. L. J. Mater. Chem., 2006, 16, 4427-4434. Liu, B.;Zeng, H. C. J. Am. Chem. Soc., 2004, 126, 16744-16746. Lu, Q. F.; Zeng, H. B.; Wang, Z. Y.; Cao, X. L.; Zhang, L. D. Nanotechnology, 2006, 17, 2098-2104. Yang, L.; May, P. W.; Yin, L.; Brown, R.; Scott, T. B. Chem. Mater., 2006, 18, 50585064. Yang, L.; May, P. W.; Yin, L.; Scott, T. B.; Smith, J. A.; Rosser, K. N. Nanotechnology, 2006, 17, 5798-5804. Yang, L.; May, P. W.; Yin, L.; Huang, Y. Z.; Smith, J. A.; Scott, T. B. Nanotechnology, 2007, 18, 335605 (5pp). Yang, G. W. J Prog. Mater. Sci., 2007, 52, 648-698. Wang, J. B.; Zhang, C. Y.; Zhong, X. L.;Yang, G. W. Chem. Phys. Lett., 2002, 361, 86-90. Sounart, T. L.; Liu, J.; Voigt, J. A.; Hsu, J. W. P.; Spoerke, E. D.; Tian, Z.; Jiang, Y. B. Adv. Funct. Mater., 2006, 16, 335-344. Liu, A. Y.; Cohen, M. L. Science 1989, 245, 841-842. Li, J.; Cao, C. B.; Hao, J. W.; Qiu, H. L.; Xu, Y. J.; Zhu, H. S. Diamond Relat. Mater., 2006, 15, 1593-1600. Yang, L.; May, P. W.; Huang, Y. Z.; Yin, L. J. Mater. Chem., 2007, 17, 1255-1257. Yang, L.; May, P. W.; Yin, L. Nanotechnology Research Developments, Nova Publishers, Hauppauge, NY, 2007. Sylvestre, J. P.; Kabashin, A. V.; Sacher, E.; Meunier, M. Appl. Phys. A, 2005, 80, 753758. Pyatenko, A.; Shimokawa, K.; Yamaguchi, M.; Nishimura, O.; Suzuki, M. Appl. Phys. A, 2004, 79, 803-806. Rakshit, R. K.; Budhani, R. C. J. Phys. D, 2006, 39, 1743-1748. Wang, J. B.; Lei, J. L.; Wang, R. H. Phys. Rev. B, 1998, 58, 11890-11895. Brinker, C. J.; Lu, Y. F.; Sellinger, A.; Fan, H. Y. Adv. Mater., 1999, 11, 579-585.

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[36] Liu, B.; Zeng, H. C. Small, 2005, 1, 566-571. [37] Penn, R. L.; Banfield, J. F. Science, 1998, 281, 969-971. [38] Lu, T. R.; Kuo, C. T.; Yang, J. R.; Chen, L. C.; Chen, K. H.; Chen, T. M. Surf. Coat. Tech., 1999, 115, 116-122. [39] Zimmerman, J. L.; Williams, R.; Khabashesku, V. N.; Margrave, J. L. Nano Lett., 2001, 1, 731-734. [40] Likhacheva, A. Y.; Paukshtis, E. A.; Seryotkin, Y. V.; Shulgenko, S. G. Phys. Chem. Miner., 2002, 29, 617-623. [41] Liu, A. Y.; Cohen, M. L. Phys. Rev. B 1990, 41, 10727-10734. [42] Zhang, Z. P.; Shao, X. Q.; Yu, H. D.; Wang, Y. B.; Han, M. Y. Chem. Mater., 2005, 17, 332-336. [43] Wang, C. M.; Pan, X. Q.; Rühle, M.; Riley, F. L.; Mitomo, M. J. Mater. Sci., 1996, 31, 5281-5298. [44] Maenosono, S.; Okubo, T.; Yamaguchi, Y. J Nanoparticle Res., 2003, 5, 5-15. [45] Salamanca, J. M.; Ciampi, E.; Faux, D. A.; Glover, P. M.; McDonald, P. J.; Routh, A. F.; Peters, A. C. I. A.; Satguru, R.; Keddie, J. L. Langmuir, 2001, 17, 3202-3207. [46] Bowden, N.; Terfort, A.; Carbeck, J.; Whitesides, G. M. Science, 1997, 276, 233-235. [47] Li, M. PhD thesis, University of Bristol, 2000, pp. 132. [48] Matsushita, S.; Onoue, S. Y. Three-dimensional self-assemblies of nanoparticles, Nanocrystals Forming Mesoscopic Structures, Pileni M P (ed.), Wiley-VCH Verlag GmbH & Co. KGaA, 2005, Weinheim, Germany, p.140. [49] Whitesides, G. M.; Simanek, E. E.; Gorman, C. B. Nano Advanced Study Institute on Chemical Synthesis: Gnosis to Prognosis, eds. Chatgilialoglu C, Snieckus V, 1996, Kluwer, Dordrecht, Netherlands, pp. 565-588. [50] Jeon, N. L.; Finnie, K.; Branshaw, K.; Nuzzo, R. G. Langmuir, 1997, 13, 3382-3391.

In: Nanotechnology… Editors: C.J. Dixon and O.W. Curtines, pp. 435-457

ISBN: 978-1-60692-162-3 © 2010 Nova Science Publishers, Inc.

Chapter 13

PROGRESS OF SELF-STANDING DIAMOND FILM FABRICATED BY DC ARC JET PLASMA CVD G.C. Chen1,*, F.X. Lu1, B. Li2, C.M. Li1, W.Z. Tang1, J.H. Song1, L.F. Hei1 and Y.M. Tong1 1

School of Materials Science and Engineering, University of Science and Technology Beijing, 100083, Beijing, P. R. China 2 School of Metallurgy and Ecology, University of Science and Technology Beijing, 100083, Beijing, P. R. China

Abstract Self-standing diamond films were fabricated by a 30 kW DC Arcjet CVD system. The novel progresses, including layer-structured film (nano-/micro-crystalline layer) fabrication, high orientated film deposition with high growth rate at very high ratio of CH4/H2, crack-free thick and large area films growth, and single crystal fabrication, were reported. Layer-structured self-standing films, 2- and 4-layered ones, were fabricated by fluctuating the ratio of methane to hydrogen with deposition time. Results of scan electronic microscopy (SEM) and Raman spectra showed that the layered films were constructed by the micro-crystalline grains layer / nano-crystalline grains layer. The residual stress within the films were balanced, and even diminished in the certain layer. The layer containing nanocrystalline grains due to a plenty of secondary nucleation could weakly inherit the columnar growth feature of the overlaid layer containing micro-crystalline grains. The grain size and growth orientation of the layer containing micro-crystalline grains could be adjusted by introduction a mid-layer containing nano-crystalline grains. Growth rate was over 10μm/hr in layered film fabrication. The effect of very high concentration of CH4 in H2, 10%≤CH4/H2≤ 25%, was studied on the film morphology and orientation. Diamond films with morphology containing nice faceted micro-sized grains were obtained with CH4/H2 up to 17%. The film composition change was found by Raman spectra. High (111)-oriented films were deposited under the condition of CH4/H2=15% at the maximum growth rate about 50μm/h. Deposition temperature could influence both the morphology and orientation of the diamond films. The higher deposition temperature, the higher CH4/H2 could be allowed to deposit micro-sized grain-containing *

E-mail address: [email protected]. (corresponding author)

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G.C. Chen, F.X. Lu, B. Li et al. films. However, high deposition temperature would spoil (111)-orientation. As a consequence, (220) and (311) would be enhanced. Crack patterns occurring in self-standing films were classified as network shape, river shape and circle shape. The distribution and style of dominating crystalline surface was found to influence the strength of self-standing film. The films with 60-120mm of diameter and 2mm of thickness were successfully deposited by controlling of dominating crystalline surface in the films. A new approach to single diamond crystal fabrication by arc jet was proposed and discussed. This method was named as “stable-tip” method which was applied to overcome the morphology instability. Single crystal, 1×1×0.6mm3 in size, was successfully fabricated by this method. The synchrotron radiation topography was adopted to characterize this single crystal diamond.

Introduction Self-standing diamond film, also as free-standing one [1], is the film that exists without adhering to any substrates, and its thickness is usually 500-1000 micrometers. Sometimes, self-standing diamond film (not discussed in this paper) also means that whose substrate is got rid of by chemical solution, and the thickness even less than 15 micrometers. The property of self-standing diamond film is very similar to that of nature diamond, and its area has reached to 175mm in diameter [2]. Therefore, self-standing diamond film is very promising material for application in window/dome of missile [3], thermal sink for electronic power device [4], substrates of microelectrical engineering mechanical systems (MEMS) [5], surface acoustic wave (SAW) devices [6] as well as cutting tool. In light of the film quality, self-standing diamond film is classified as tool-grade, optical-grade and device-grade [7,8]. Tool-grade film is usually opaque due to relatively high concentration of non-diamond composition. Optical-grade film is transparent with less or without non-diamond composition. The device-grade one is applied into device fabrication. No matter which grade film it is, mechanical strength of the film is expected as high as the nature diamond. However, self-standing film has to face the fact that the fracture strength is one-order-magnitude lower than that of nature single crystal. Gray and Windischmann [9] regard that the high residual stress in the film results in the low strength. They give an expression of the stress in the film:

σf=σth+σin where σth is thermal stress, and σin is instinct stress. σth is due to the different expansion property between diamond and the substrate. σin is related to the structure of the film. For self-standing diamond film, the first term in the right side of the expression is equal to zero due to the release of thermal stress. The high stress in the film is, therefore, ascribed to the instinct stress. This is not to say that there is no any effect of thermal stress on the film strength. The fact is that the thermal stress has exerted on the film when the film peels off from the substrate after the deposition finishes. As the stress is very high, micro- even macrocracks may occur. These micro- or macro-cracks result in low strength of the film. Among various diamond synthesis techniques, DC Arc Jet Plasma enhanced CVD is known as high growth rate diamond fabrication technique, and it is the first technique that

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achieves at the fabrication of self-standing diamond film [10]. Kurihara et al first reported diamond fabrication by this technique in 1988 [11], and then Ohtake and Yoshikawa achieved at the highest growth rate, ~930μm/hr, in 1990 [12]. The very high growth rate is commonly regarded due to high enthalpy arcjet flows in this technique. Cappelli [10] regards that “Arcjet” is from the fact that the energy, via ohmic heating in the electric arc discharge, is dissipated into a flowing, high-pressured gas, resulting in the increase of the gas enthalpy and kinetic. The temperature in the core of the arc can reach 40,000K, and high enough to form the plasma. The plasma’s conductivity increases with increasing temperature. The high gas temperature results in vigorous decomposition of reactant gas and relatively high velocities (~1-10km/sec) in an arcjet flow. Therefore, more efficient delivery of the atomic hydrogen to a substrate occurs because of productive atomic hydrogen and reduced boundary layer thickness. He gives an expression for the fraction of the plasma jet atomic hydrogen that is delivered to a substrate surface in terms of the expansion pressure (in torr), pe, the plasma jet velocity (in cm/sec), u0, the substrate diameter (in cm), ds, and the reaction probability of atomic hydrogen on the substrate, γ: [H]s [H]0 ≅

1 2

(1)

{1+1100γ peds/u0 }

Here, [H]o and [H]s are the mole fractions of atomic hydrogen in the incident plasma jet and substrate respectively. It is apparent from this equation, that the maximum delivery of atomic hydrogen to a substrate from an arcjet is favored by reduced pressure and increased jet velocities. The high growth rate, on the other hand, contributes some reverse effect on the film quality. The most obvious phenomenon is morphology instability [13]. Palmer-Gordon described the course of CVD by the expression as: ω= −

Ω2υDsγ 4 γc0eqΩD0 3 ΩD0(n0-c0eq) kBT κ − (1+sκ)kBT κ + (1+sκ)(L+s) κ

(2)

Where ω is the stability parameter, Ω is the atomic volume in the bulk, υ is the concentration of surface particles, γ is the surface tension, Ds is the surface diffusion coefficient, D0 is the gas phase diffusion, c0eq is the equilibrance concentration of the active species in contract with a flat interface, s is the sticking parameter, L is the mean curvature of the interface, κ is a constant, and T and kB are the absolute temperature and Boltzmann’s constant [14]. Palmer-Gordon model is based on the assumption that: there are three procedures for deposit species arrive at the interface. It is either (a) incorporated into the surface; (b) it can move over the surface by surface diffusion; or (c) it can be re-evaporated into the gas phase. Therefore, the first term in right side of the equation is related to the surface diffusion, the second is due to deposition or re-evaporation of surface particles, and the third term represents the net flux of particles to the surface and is governed by the gas phase diffusion coefficient. The first two terms in the equation work in such a manner as to stabilize and smooth out the interface, whereas, the third term is destabilize and causes morphological

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instability. However, this model neglects the effect of a temperature gradient in the gas and in the growing film. Mahalingam and Dandy [15] modify the Palmer-Gordon model. They assume that the growth of diamond is due to both surface reaction and surface migration since re-evaporation or desorption is expected to be negligible in the case of diamond deposition. They mentions that diamond phase of carbon is very stable and gasification of diamond will occur very slowly. The etching rates of diamond during diamond synthesis ambient are orders of magnitude smaller than the deposition rates. Based on this consideration, the kinematic equation for the growth rate of the deposited diamond film in the direction normal to the surface (i.e., z-direction) is illustrated as: DsςΩ2Γ d2 r dCsCH3 Vz=ν(-DCH3 dz ) + ( k T ) dz2 B s

(3)

where: Vz is the linear growth velocity along z-direction; ν is the molar volume of diamond; Ds is the surface diffusivity of C atoms on diamond; ς is the interfacial energy of diamond per unit area; Ω is the molecular volume of diamond; Γ is the number of atoms per unit area; kB is the Boltzmann constant; and r is the total curvature at a point on the interface. By this model, Mahalingam and Dandy analyze the effect of deposition temperature and the species concentration. They point out that as the growth rate increases, the morphological instabilities are enhanced in the growing film, and these instabilities are manifested as a finger-like morphology. As the deposition is concerned, the increase of deposition time can enhance the morphological instability. Besides the morphology instability, the film quality also suffers from the rapid growth rate in DC Arcjet method. The films usually contain relatively high compressive stress, and much more cavities. General speaking, paralleled with low growth rate techniques, such as microwave and hot filaments, the DC arc method possesses the fastest diamond nucleation and crystal growth rates. The higher deposition rate tends to form smaller and poorly developed crystals. These tiny imperfect crystals could trap more non-diamond carbons in diamond crystals [16]. Cappelli, Goodwin and Harris created several methods to study the plasma property, including electron temperature and electrode processes and arc jet enthalpy balance. The detail can be seen in reference [17-20]. In summary, self-standing diamond film faces several problems, such as low strength, morphological instability and much more defects. The high growth rate enhances these problems by DC Arc jet method. How to increase the strength and whether the morphological instability can be controlled, therefore, is the challenge for diamond fabrication by DC Arc jet method. Here, we report our work in novel progresses of fabricating self-standing film.

Experiments A 30 kW DC Arc Plasma Jet CVD system was applied to fabricate self-standing films [21]. The fed gasses were Ar, H2, and CH4. The pressure of deposition chamber was usually between 4-8kPa. The substrate was multi-crystalline molybdenum block with 50mm in diameter and 30mm in height. In order to enhance the nucleation density, the molybdenum

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substrate was cleaned in an ultrasonic bath with acetone and then rubbed with 20, 10, 5 and 0.5μm diamond powder in turn, at last drying after rinsed by ethanol. Deposition temperature was from 900 to 11000C measured by IR pyrometer. During the experiments, we usually changed the heat-conducting mediums between substrate and its holder, which allowed us to control substrate temperature independently of deposition parameters. As-deposited films were characterized by scanning electronic microscope (SEM, LEO 1450 or XL30 S-FEG), X-ray diffraction (XRD, D/MAX-RB) and Raman spectroscopy (JY－T64000). The gas phase species in the reaction region were diagnosed by optical emission spectra (OES) during diamond films growth. The emission light from the reaction region was transferred to the entrance slit of a monochrometer (WDP500-2A) through the quartz window and a lens system with an optical fiber. Scan rate was 200nm/min., and collection time was 1sec./step controlled by a computer. The detection position was focused 0.8mm above the substrate. The typical optical emission spectrum from the CH4-H2-Ar mixture in the deposition system is illustrated in Fig.1. The major optical emission lines from atomic H, CH and C2 are observed. No Ar related emission lines, like 714 nm, 750 nm, 763 nm, are found due to the limit of narrow scanning band of the monochrometer (between 300 and 700 nm). 4200

Hα

Intensity (a.u.)

3600 3000 2400 1800

Hβ

1200 600

CH

Hγ

C2

H2

0 400

450

500

550

600

650

Wavelength (nm)

Figure 1. Typical optical emission spectrum from the CH4-H2-Ar mixture in the high power DC arc plasma jet CVD system.

Results and Discussion 1. Fabrication of Nano-/Micro Multi-layered Self-standing Diamond Film As stated above in part of introduction, low strength is regarded due to high residual stress exerting on the film. Novel film structure fabrication may be useful to reduce or redistribute the stress so as to raise the film strength. Catledge et al proposed nano/micro multiplayer diamond film in 2000 [22], and Takeuchi et al found that the fracture strength of this multiplayer film rose up to 30% compared with the ordinary micro-crystalline diamond film in 2001 [23]. Jiang et al [24] believed that the interface between nano/micro should effectively prevent crack from propagation. That may be the reason why the fracture strength became higher. Not only strength was expected to be raised, smooth surface was also

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obtained by this multiplayer film based on the works of Jiang et al [24] and Xin et al [25]. It shall be pointed out that these reported layered films are not self-standing ones, and all fabricated by microwave or hot filaments CVD methods. For DC Arcjet method, the deposition procedure is difficult to control. Meanwhile, the change of film structure may result in the film peeling off during deposition. Thus, it is a challenge to fabricate nano/micro multiplayer self-standing diamond film by DC Arc Plasma Jet CVD. We tried deposition of self-standing films with this layered structure. During deposition, the feed gasses flow rates were of H2 6 slm, Ar 2 slm, and CH4 60-900 sccm, respectively. The deposition chamber pressure was maintained at 4kPa. The substrate temperature was kept as 910±100C. By fluctuating the flow rate of CH4 with deposition time, layered film structures would be fabricated. The 2- and 4-layered film procedures were specified as following: The 2-layered film was fabricated as that: 5% of CH4/H2 was applied in the first beginning 15 minutes to enhance diamond nucleation. Then, 1% was used to deposit microcrystalline diamond grains. 2 hrs later, 15% was used for 2 hrs to enhance the secondary nucleation and expect to obtain the following deposition of nano-crystalline diamond grains. The total deposition time was 4.25hrs. The 4-layered film was fabricated as that: the first beginning 4.25 hrs deposition was as the same as that of procedure 1. Then, 1% was used for a 2 hrs deposition. After then, the ratio was increased to 15% for another 2 hrs deposition to deposit nano-crystalline grains. The total deposition time was 8.25 hrs. The deposition procedures are marked in Fig. 2.

Figure 2. Ratio of CH4 /H2 during deposition process.

Figure 3 is the cross-section and growth surface SEM results of as-fabricated diamond films. From Fig. 3a and 3c, the layered structures are clearly distinguished as 2-layered and 4layered films though the boundary between each layer is not very straight. The layer fabricated under low ratio of CH4 consists of micrometer-sized columnar crystals. However, the layer fabricated under high ratio of CH4 does not, and the columnar growth feature is weakly inherited in these high ratio of CH4 fabricating layers. The growth surfaces are terminated by cauliflower-like morphology. Our former work showed that there were nanosized diamond crystalline grains in the film with “cauliflower” morphology [26]. Results in Fig.3 and in reference [27] indicate that (1) the grain size in the layer with micrometer-sized columnar crystals can be adjusted by introducing a different thickness mid-layer; (2) the layer deposited by secondary nucleation under the condition of high CH4/H2 ratio can inherit the

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growth feature of the overlaid layer. The inherited growth feature becomes strong or weak depending on the deposition time. This point can also be seen in other author’s results [25]. Growth rate can be calculated by the thickness divided by deposition time. It is about 12mm/hr for 2-layered film and 9mm/hr for 4-layered film. This means that the growth rate decreases with the increase of the layer number. It is also found that the deposition rate of layer with columnar micro-sized crystal is smaller than that of the layer deposited under high CH4/H2 ratio condition. This phenomenon has also been mentioned by Sternschulte et al [28]. The reason is due to more amounts of nutritional radical for diamond growth. Although the growth rate is different from each other, it is all about 10mm/hr. This value is much higher than that in case of fabricating non-self-standing film by HFCVD and MWCVD. This shows the attractive application of DC Arc Plasma Jet CVD in fabrication of layered structure diamond film.

a

b

c

d

Figure 3. SEM results of as-grown layered films. a.cross-section of 2-layered film; b. growth surface of 2-layered film; c. cross-section of 4-layered film; d. growth surface of 4-layered film.

Raman spectra detection was done on both layered films’ nucleation surface and growth surface. The results are shown in Fig.4. The peak ascribed to sp3 band appears near 1332cm-1 in both the nucleation surface and the growth surface. There are another two peaks of which one appears at 1128.1cm-1, and the other one is the broadening peak between 1410 cm-1 and 1620cm-1, in the growth surface. Nemanich et al believed that the peak at 1128.1cm-1 was ascribed to nano-sized diamond [29]. However, Ferrari et al regarded it due to the scattering

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of TPA [30]. The broadening peak can be discomposed to two peaks of which one at 1548cmascribed to G-band of sp2 [31], and the other one at 1470.9cm-1 is believed to be relevant with the nanometer diamond structure [32,33]. This means that the layer grown under high ratio of CH4/H2 condition consists of nano-crystalline diamond. These results are consistent with the SEM results. The peak position of sp3 is different from nucleation surface and growth surface, which indicates that the stress style is different from each other. It is compress stress in nucleation surface, and tensile stress in growth surface. Thus, there exists an interlayer without stress in the film. Base on our work in reference [27], micro-crystalline layer can be fabricated by introducing a interlayer with proper thickness. 1

a

b

Figure 4. Raman results of layered films. a. 2-layered film; b. 4-layered film.

The growth surfaces and the nucleation surfaces of the films shown in Fig.3 were surveyed by XRD. The results are shown in Fig.5. For 2-layered film and 4-layered film, (111)-crystalline plan (at 2θ=430) possess absolutely strong diffraction intensity while (220)crystalline plan (at 2θ=740), (331) - crystalline plan at 2θ=930 and (400) - crystalline plan at 2θ=1180 possess weak intensity in both the nucleation surface and the growth surface. Thus, the growth orientation is <111> in both the nucleation surface and the growth surface. We found the phenomena of transition of orientation by introducing a proper mid-layer between two micro-crystalline layers [27].

a

b

Figure 5. XRD results of as-grown films. a. 2-layered film; b. 4-layered film.

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The fracture strength is measured by three-point-bend method [34]. The values of each film are illustrated in Table 1. As reference value, the fracture strength of micro-layered film is also measured under the condition of the thickness as same as the multiplayer film. One can find that the value is 20% higher than that of mono-layer film. Table 1. The fracture strength of layered film Film structure Mono-layered microcrystalline film 2-layered film 4-layered film

Thickness (μm) 170±5 170±3 174±4

Fracture strength (MPa) 334±20 392±12 408±32

In summary, layered diamond films were fabricated by fluctuating the ratio of methane to hydrogen in high power DC Arc Plasma Jet CVD. The film structure consists of the layers containing nano- and micro-crystalline grains. The layer containing nano-crystalline grains formed by a plenty of secondary nucleation can weakly inherit the columnar growth feature of the overlaid layer containing micro-crystalline grains. The residual stress in film can be balanced by this layered structure, and even can be diminished by a mid-layer containing nano-crystalline grains with proper thickness. The grain size and growth orientation of the layer containing micro-crystalline grains can be adjusted by introduction a mid-layer containing nano-crystalline grains. The fracture strength is 20% higher than mono-layer micro-diamond film. The growth rate was about 10mm/hr in layered film fabrication.

2. High Oriented Film Fabrication with Very High Ratio of CH4/H2 Micro-crystalline diamond film is usually obtained by using mixture of hydrocarbon gas (commonly methane) and hydrogen in which the concentration of methane in hydrogen, CH4/H2, is always less than 2% [35] since Spitsyn [36] and Matsumoto [37] succeeded in deposition of diamond under low pressure [38]. Nano-crystalline diamond grains will appear by using methane rich fed gasses, or even without hydrogen [39]. The reason is well known that atomic hydrogen plays an important role on suppressing the formation of unsaturated carbon nuclei, and etching graphitic nuclei [38]. The more H fraction in reactive gas phase, the easier faceted grains are obtained. DC Arc Plasma Jet is believed to possess higher number density of atomic or ionic hydrogen compared to “cool plasmas” (such as microwave plasma). Therefore, much higher concentration of methane can be used as 5-15% [40], and very high growth rate (near to 1mm/h) has been achieved [12]. So far, however, the morphology obtained at high concentration of methane (over 10%) is un-faceted. Therefore, we report our experimental results of diamond deposition by DC Arc Plasma Jet CVD under the condition of very high concentration of methane in hydrogen. The high concentration of CH4/H2, up to 25%, was used to deposit diamond film. The asdeposited films’ morphologies of growth surfaces are shown in Fig.6. The morphology with nice faceted grains can be observed in the films deposited at the concentration of 10% and 15%, even 17%, seen in Fig. 6a, b and c. The size of grains in these four films is in the micron order-of-magnitude. Cauliflower-like morphologies occur in the films deposited at the concentration of 25%. In fact, a lot of secondary nucleation can be found in the film deposited

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at CH4/H2=17% though nice faceted grains still remain. It seems that the morphology of the film deposited at CH4/H2=17% is the middle-state between cauliflower-like morphology and nice faceted grain-containing morphology. In other words, the secondary nucleation is the important factor to result in the morphology changing from the nice faceted grain to cauliflower.

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Figure 6. SEM results of as-grown films under different ratio of CH4/H2. a. 10%; b. 15%; c. 17%; d. 25%.

This morphology change has also been proven by Raman Spectra survey (seen in Fig.7). It can be seen that the composition in the film changes with the increase of CH4/H2. It is pure diamond component first, and then nano-sized crystalline diamond existence, finally graphitic composition occurring. Considering the SEM results in Fig.6, it can be deduced that composition change is ascribed to the secondary nucleation, and results in the morphology change. Morphology change was also found when we varied the deposition temperature. Fig.8 is the growth surface and cross-section SEM results of the films deposited at different deposition temperature with CH4/H2=15%. Cauliflower-like morphology is obtained at 9000C (seen in Fig.8a), which is similar to Fig.3b and 3d. Nice faceted micro-sized grains appear when the deposition temperature higher than 9000C, seen in Fig.7b, c and d. There are cogrown grains due to secondary nucleation on these faceted micro-sized grains, and the amount of these co-grains increases with the increase of the deposition temperature. Growth rate of each film in Fig.8 is calculated as about 40mm/h, 50mm/h, 30mm/h and 40mm/h according to the film thickness and deposition time. It means that film with micro-sized grain can be grown at the rate over 30mm/h, and the maximum is about 50mm/h.

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Figure 7. Raman spectra results of the films shown in Fig.6. a. CH4/H2=10%; b. CH4/H2=15%; c. CH4/H2=17%; d. CH4/H2=25%.

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Figure 8. SEM results of as-grown films under different deposition temperature. a. 9000C; b. 9700C; c. 10000C; d. 10500C.

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XRD was used to analyze the orientation in the films shown in Fig.8. The results are illustrated in Fig.9.The intensity of (111) crystalline surface is the most strongest one, and the ratio of intensity of (220) to that of (111), I(220)/I(111), is less than 25% in all the films’ spectra. This means that (111) crystalline surface is the dominant surface in the films’ growth surfaces no matter whether the film’s morphology is cauliflower-like or faceted grain-containing. In the film with the faceted grain-containing morphology, the intensity of (220) increases with the deposition temperature increasing. The same case occurs for the intensity of (311). The phenomenon of high intensity of (311) has been reported by Ohtake [12]. Generally speaking, high (111)-oriented film can be obtained under the condition of CH4/H2=15%, and the highest degree of orientation, I(220)/I(111), can arrive at 16 occurring at 9700C of deposition temperature. The degree of orientation becomes weak with the increase of deposition temperature.

Figure 9. XRD results of as-grwon film in Fig. 8.a. 9000C; b. 9700C; c. 10000C; d. 10500C.

In summary, very high concentration of methane in hydrogen, from 10% to 25%, was applied to deposit self-standing diamond films. The effect of concentration of methane in hydrogen and deposition temperature on the films’ morphologies and orientations was studied. Morphology containing nice faceted micro-sized grains were obtained with CH4/H2 up to 17%, and high (111)-oriented films were deposited under the condition of CH4/H2=15% at the maximum growth rate about 50mm/h. Columnar growth feature was weak in the film with cauliflower-like morphology, and became weak with the increase of deposition temperature in the films with faceted micro-sized grain-containing morphology. Micro-sized grain-containing films are prone to be grown at high deposition temperature and high CH4/H2. However, high deposition temperature would spoil (111)-orientation. As a consequence, (220) and (311) would be enhanced. The change of morphology and orientation was ascribed to the change of the composition in the film due to the occurrence of secondary nucleation.

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3. Large Area Film with High Fracture Strength Self-standing diamond film is currently fabricated on the substrates, such as Si, W, Mo, WC-Co. These substrates possess the larger thermal expension coefficence than diamond (4144). Thus, the thermal stress is not able to eliminate at the time of plasma shut-off if these substrates are still employed. Sussman (45), Field (46), Windischmann (9) and Lu (47) have already noted that the film structure have a special effect on the fracture strength of selfstanding film. They all find that the fracture strength decreases with the increase of crystal size. This indicates that the fine crystalline grain is pursued. Their results seem more useful not for optical grade self-standing diamond films, but tool grade ones. Jeong (48) found that the crack-free film can be obtained if the fracture strength is larger than the shear stress at the film edge. Kamiya (49) found that the crack path is along and/or around the large crystal grain boundary. These two results indicate that the crystalline structure design is necessary for self-standing crack-free diamond film to obtain high fracture strength. In the passed two decade, lots of works (13,35,50,51) were done on crystalline structure development. The effect of growth parameters on crystalline structure have been studied well. However, there is little work on the relationship between crystalline structure and strength.

Figure 10. Relationship between growth temperature and ratio of X-ray diffraction intensity.

Fig. 10 is the relationship between growth temperature and ratio of X-ray diffraction intensity, I(111)/I(220). The ratio decreases with the increase of growth temperature on both nucleation side and growth side of the film body. It is larger than 4 below 900°C, and less than 4 over 900°C on nucleation side. On growth side, the “boundary” of growth temperature becomes 800°C, i.e., the ratio is larger than 4 below 800°C, and less than 4 over 800°C. This means that (111)- and (220)-crystalline surface alternately becomes dominant with the increase of growth temperature, and transfer temperature is higher on nucleation side than that on growth side. By reducing the size of X-ray beam to 400μm×4mm, the fractured sections are also detected. The variety of ratio, I(111)/I(220), is unlike on neither nucleation nor growth side. It is less than 4 only in a very narrow temperature range that is around 900°C. Beyond this range, the ratio is much larger than 4. This means that (111)-crystalline surface keeps on being dominant surface under most of employed growth temperature in fractured sections. (220)-crystalline surface becomes dominant only in a very narrow temperature window. In

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this research, (311)- and (400)-crystalline surface are not considered because the intensity of these two surfaces are very small, and little variety occurs. Three kinds of crack patterns are found in the film body. They are network-shape pattern, river-shape pattern and circle-shape pattern, depending on the different growth temperature. The typical patterns are shown in Fig. 11. The network-shape patterns usually occur at the growth temperature below 800°C. The cracks always pass through the whole film body in thickness direction, and are exposed on both growth and nucleation side. The samples are, therefore, very brittle, and their shapes are irregular. The circle-shape patterns take place at the growth temperature over 900°C. The curvature of the crack is, somehow, related to the curvature of the substrate. The samples show a regular shape like ribbon with certain curvature, and they are not easy broken. The river-shape patterns are never found below 800°C, and tend to combine with circle-shape patterns over 900°C. Unlike network-shape patterns, the cracks don’t always pass through the film body in thickness direction. They are only exposed on growth or nucleation side sometimes. The condition is not clear for the cracks to pass through the film body yet. The brittleness of samples is in the middle rank comparing with the former two patterns. The fractured sections are detected by AFM. The very smooth surface is found in the samples with network-shape patterns. The typical roughness is about 30nm, seeing in Fig. 12a. The terraces are found in the samples with circle-shape patterns, seen in Fig. 12b. The roughness is also small, but it is higher than the former. Both smooth surface and terraces are found in the samples with river-shape patterns. The roughness is the highest one among the patterns. The typical result in Fig. 12c shows three adjacent broken grains and their boundaries.

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Figure 11. Different crack pattern. a. network-shape pattern; b. river-shape pattern; c. circle-shape pattern.

The full width at half maximum (FWHM) of sp3 peak along the crack in river-shape pattern was measured by micro-Raman. In this procedure, we made the laser beam very near to the crack path in which the distance is less than 2μm localized by optical microscope attached to micro-Raman. The measured points are shown in Fig. 10. It is found that the FWHM of Raman peak becomes the minimum value at the tip of the crack no matter on the growth surface or the nucleation side. The values are almost same besides the crack path. The results are showing in the Table 2.

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Figure 12. SEM and AFM results of different crack path. a. and b. network-shape pattern; c. and d. river-shape pattern; e. and f. circle-shape pattern.

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nucleation side growth side nucleation side growth side

side of the crack path A-point B-point -1 9.31cm 9.22 cm-1 9.02 cm-1 8.90 cm-1 7.25 cm-1 7.25 cm-1 6.07 cm-1 5.63 cm-1

Tip of the crack C-point 8.74 cm-1 6.77 cm-1 5.89 cm-1 4.01 cm-1

Base on these results, the effect of dominant surface on film are discussed. Dominant surface vs. Crack patterns: Corresponding to the XRD results, it is easy to find that the network-shape patterns occur in the films with (111)-crystalline surface being dominant on all of the three detected surfaces, i. e., growth side, nucleation side and fractured section. The circle-shape patterns occur in the films with (220)-crystalline surface being dominant in all of the three detected surface. The river-shape patterns occur in the films with (111)-crystalline surface being dominant on one of the sides at least. AFM results indicate that the circle-shape patterns are mainly due to intergranular fracture, and network-shape patterns are due to transgranular fracture. The river-shape patterns seem due to the combination of inter- and trans- granular fracture. Therefore, we can conclude that different fracture mechanism results in different crack patterns, and different dominant crystalline surface is ready to excite different fracture. Dominant surface vs. Fracture strength We assume that the diamond film body is made of the pack of (111)- or (110)- crystalline surface. In 3-point-bending test, the film body will sustain shear stress. Fracture will occur when the shear stress is over fracture strength. It can be seen that the shear stress, τ, is 1/2P (P: load) in both (111)-surface and (110)-surface. According to reference (52), the cleavage energy is related to surface energy. It is known that the surface energy increase according to the order as v(111)
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analysis of stress state in film body suggests that the (111)-surface is easier cleaved than (110)-surface under the same load. Crack propagation depends on the crystalline quality. It will be blocked by the crystal with little point default. By controlling the crystalline structure of the films, crack-free self-standing diamond films with 2mm in thickness and with 120mm in diameter are successfully grown.

Figure 13. Crack-free diamond wafer with 2mm of thickness and 120mm or 60mm of diameter.

4. Fabrication of Single Crystal CVD single crystal diamond was pursued almost since Spitsyn [36] and Matsumoto [37] succeeded in deposition of diamond under low pressure [38]. Flame CVD first showed the possibility of single crystal fabrication [54], and then microwave CVD fabricated the single crystal diamond [55]. It is recently that carat-grade single crystal diamond was fabricated by microwave CVD [56,57]. So far, however, no report on DC Arcjet is found. As mentioned in part 1, morphology instability easily occurs in DC Arcjet method. This may be one of the important reasons to keep this method from fabrication of single crystal diamond. However, carat-grade single crystal diamond fabrication indicated that the very high growth, ~100μm/hr, is necessary by Microwave method. DC Arcjet, in other hand, is well known for its very high growth rate. Therefore, DC Arcjet may be more feasible to fabricate single crystal if morphology instability can be controlled. From the analysis of D.S. Dandy [15], we can see that the morphology instability is due to the change of radical concentrate and the change of growth tip curvature of crystal grain. If these two changes can be eliminated, the morphology instability can be controlled. Therefore, a method, named as “stable-tip”, is proposed to try to fabricate single crystal diamond in DC Arcjet. This method is that the substrate holder is descent during growth so that the growth tip of fabricated crystal is fixed to a position referenced to the sample holder. By this method, we fabricated a 1×1×0.6mm3 single crystal diamond. The optical microscope result is shown in Fig. 14. To distinguish the substrate, synchrotron radiation topography was applied to detect the CVD fabricated diamond and nature diamond. Fig. 15 shows the synchrotron radiation topography result. It can be seen that there are very clear dislocation net in {111} surface in the substrate. However, the Laue spots becomes deformed, and there are very high dense dislocation in {111} surface of the CVD fabricated diamond.

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Figure 14. Single crystal diamond grown by “stable-tip” method. a. lateral view; b. top view.

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Figure 15. Synchrotron radiation topography results of nature diamond and grown diamond. a. nature diamond single crystal (substrate); b. enlarged (111) surface’s Laue spot in a.; c. CVD diamond single crystal; d. enlarged (111) surface’s Laue spot in c..

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Conclusion 30 kW DC Arcjet CVD system was used to fabricate self-standing diamond film. Layerstructured, i.e., nano-/micro-layer, films were fabricated by fluctuating the ratio of methane to hydrogen. The film structure transformation could be ascribed to the plenty of secondary nucleation. The film with this layered structure possessed higher fracture strength than that with microcrystalline mono-layer. Very high concentration of methane in hydrogen, from 10% to 25%, was surveyed. Morphology containing nice faceted micro-sized grains were obtained with CH4/H2 up to 17%, and high (111)-oriented films were deposited under the condition of CH4/H2=15% at the maximum growth rate about 50μm/h. The distribution and style of dominating crystalline surface could influence the strength of self-standing film. The Large area film with 120mm of diameter and 2mm of thickness was successfully deposited by controlling of dominating crystalline surface in the film.. “Stable-tip” method was used to overcome the morphology instability. Large single crystal was successfully fabricated by this method. The sychrotron radiation was one of effective testing technique to characteristic CVD single crystal diamond.

Acknowledgment This work was supported by NSFC (No.50472095), Beijing Novel Project (No. 2003A13) and Beijing NSF(No.2062015).

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In: Nanotechnology… Editors: C.J. Dixon and O.W. Curtines, pp. 459-477

ISBN: 978-1-60692-162-3 © 2010 Nova Science Publishers, Inc.

Chapter 14

NANOSHELL ARRAYS: FABRICATION AND ENHANCED PHOTOLUMINESCENCE Zhipeng Huang and Jing Zhu* Beijing National Center of Electron Microscopy, Tsinghua University, Beijing, 100084 (P. R. China) Laboratory of Advanced Materials, Department of Materials Science and Engineering, Tsinghua University, Beijing, 100084 (P. R. China)

Abstract Low emission efficiency of silicon (Si) based light emission devices (LED) still blocks application of Si-LED. Therefore, studies focusing on improving light emission of Si-LED still attract researches’ passion. Recently, we have developed a new method combining nanosphere lithography and pulsed laser deposition to fabricate Si-based arrays nanostructures, and have obtained remarkably enhanced photoluminescence (PL) from these structures. The Si based nanostructures are hemisphere shell arrays (HSSAs) or nanoflower arrays assembled by silicon-germanium (SiGe) alloy. These structures include non-closepacked and close-packed ones, single layer and multilayer ones, as well as arrays on different substrates. We investigated the photoluminescence of these arrays structures, and found that all these structures could enhance the photoluminescence intensities. Among them, the enhancement of light emission from SiGe double layer HSSAs (DL-HSSAs), which is as high as 700 folds, is the highest among those of all structures. Employing transmission electron microscopy (TEM), scanning electron microscopy (SEM), time-resolved PL, and electromagnetic simulation etc, we found the enhancement of light emission in Si based nanostructures originated mainly from the increase of extraction efficiency of photons from the nanostructures. The electromagnetic simulation of enhancement matched well the experiment data. We also found that these enhancements are related to degree of order of arrays. In highly order arrays, the enhancement is higher than that in other arrays.

*

E-mail address: [email protected]

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1. Introduction In past decades, scaling-down has been a common scheme to improve performance of integrated circuit. However, scaling-down scheme is facing varied problems, such as interconnects latency, bandwidth limitation and crosstalk etc [1,2,3,4], which become more prominent as the decrease of line-width in large-scale integrated circuit. These problems are intrinsic features of circuit employing electron as carrier of information. Optic or optoelectronic interconnection, in which photon is carrier of information, will help to overcome these obstacles, and therefore attract researchers’ attention [5,6,7]. Among candidates, silicon-based optic interconnect is noteworthy because of its low cost and compatibility with current complementary metal-oxide-semiconductor process [8,9,10]. Guide, modulation, and transmit device of photon are basic units for optoelectronic interconnect. Silicon-based optic guide and modulation devices have been successfully demonstrated [8]. However, silicon-based photon transmitter remains in development. The challenge is to increase especially low intrinsic emission efficiency [11] related to indirect band gap nature of silicon. Many efforts have been attempted to increase emitting efficiency of silicon. The approaches include the modification of electronic structure and optical structure. Fabricating nano-crystal silicon [12,13,14,15], doping silicon with reactive center [16,17,18], tuning band gap by formation of silicon germanium alloy [19,20], employing defects such as dislocation [21,22] or point defects [23,24], increasing local electron density [25,26], and so on, are common ways to modify electronic structure of silicon. The popular methods to modify optical structure are fabricating surface pattern like inversed pyramid arrays [27] or SOI photonic crystal structure [28,29,30]. Here we introduce a simple and low-cost method to fabricate silicon-based nanostructures, and show the enhancement of room temperature photoluminescence of these structures. We will firstly introduce the fabrication details, photoluminescence properties, morphology and microstructure characterization, and electromagnetic simulation concerning single- and double-layer Si/Ge hemisphere shell arrays (SL-/DL-HSSAs) on silicon substrate [31], and then will introduce the fabrication and photoluminescence of multilayer Si/Ge HSSAs (ML-HSSAs) on Si substrate as well as on SiO2 nanosphere arrays on Si substrate.

2. Single- and Double-Layer Hemisphere Shell Arrays 2.1. Fabrication Procedures A five-step procedure was employed to fabricate HSSAs, as shown in Figure 1 and introduced in details below: 1) Single layer polymer nanosphere arrays on silicon substrate were used as template in our experiment, as shown in Figure 1a. Arrays of polymer nanosphere, which was polystyrene (PS) nanosphere with nominal diameter of 400 nm in our experiment, was formed by selfassembly on Si substrate. Before self-assembly process, the substrate was washed in turn in boiling Piranha solution and RCA solution, and then rinsed with copious amount deionized water. These treatments of substrate produced a hydrophilic surface suitable for the self-

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assembly of PS nanosphere. Drop-coating was employed for the self-assembly of PS nanosphere. The substrate was horizontally placed, and a drop of PS nanosphere/water solution was applied onto the substrate. With the control of PS nanosphere concentration and dose of PS nanosphere solution, large-area single layer PS nanosphere arrays with highly ordering on Si substrate were obtained. However, double layer or multilayer arrays were also inevitable in this simple self-assembly method. The single layer, double layer or multilayer region of arrays showed different colors, and could be distinguished by naked eye. 2) The close-packed PS nanosphere arrays fabricated in step 1 were transferred into nonclose-packed ones via reactive-ion-etching (RIE), as shown in Figure 1b. For RIE, the etching gas was oxygen, the flow rate was 40 SCCM, the pressure was 5 Pa and the applied RF power was 30 W. The interstice between PS nanospheres can be adjusted by varying the etching duration of RIE. For the resulting introduced in this chapter, the duration of RIE was 105 seconds. 3) The non-closed-packed PS nanosphere arrays were used as template for the deposition of Si/Ge alloy, as shown in Figure 1c. The Si/Ge alloy was deposited via pulsed-laserdeposition (PLD) at room temperature. The temperature in PLD process was important, because temperature higher than glass transition temperature (Tg) of PS caused the melting of polymer and thus the failure of template. Tg of PS is about 110 oC [32]. In the deposition, a sintered Si1-xGex (x=50%) target was used, the pressure was lower than 5x10-5 Pa, the wave length of KrF pulsed laser beam was 248 nm, the deposition pulse rate was 2 Hz, and the power of laser pulse was 180mJ. The deposition was conducted at room temperature for 10~30 minutes. Si/Ge alloy deposited at room temperature, which would act as scaffold in the following process, was amorphous and had no contribution to the luminescence property. Therefore, the thickness of this scaffold should be as thin as possible. After this procedure, SiGe hemisphere shells were formed above the PS spheres and SiGe mesh was formed between interstices of non-close-packed PS nanosphere arrays. 4) The template, PS nanosphere arrays, was removed by heating the substrate at 700~900 o C for more than 30 minutes. During the elevation of temperature, PS spheres were removed and SiGe hemisphere shells moved down to contact with SiGe mesh, leading to the formation of continuous SiGe scaffold, as shown in Figure 1d.

Figure 1. Scheme for the fabrication procedures of Si/Ge HSSAs. Low-left part of d,e have been drawn transparent to show clearly the shell structure. (Reprinted with permission from Ref. [31]. Copyright AIP 2007).

5) The Si/Ge alloy was deposited at 700~900 oC and proper duration, which is 55 minutes for SL-/DL-HSSAs samples in this chapter, to obtain Si/Ge alloy with desired thickness.

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Except for the deposition temperature and duration, the other parameter of PLD was identical to what descript in step 3. After this step, a Si/Ge alloy film with hexagonal ordering protrude arrays was formed, as shown in Figure 1e.

2.2. Characterization of Morphology Surface morphologies of HSSAs were investigated by scanning electron microscope (SEM). Large-area SiGe HSSAs could be seen on the surface of silicon substrate. Figure 2a shows a low magnification bird-eye view SEM image of SL-HSSAs. The tilted angle of sample was about 15o. HSSAs were a replication of the top surface morphology of PS sphere arrays, and therefore exhibited highly hexagonal ordering. The number of layer in HSSAs structure was tuned by that of the PS sphere arrays. By employing single layer template, SLHSSAs could be obtained, and DL-HSSAs by double layers or multilayer template. The details of structures could be seen in bird-eye view and cross-sectional view SEM images of SL- and DL-HSSAs. In single layer structure, the patterned film showed a hexagonal ordering, and there was a hollow hemi-sphere in each protrudes. The center-to-center distance between shells was about 400 nm, matched well the average diameter of untreated PS sphere. For each hemi-sphere shell, the diameter was about 320 nm, and the thickness of the shell was about 80 nm. The diameter and thickness of each shell was determined respectively by the duration of RIE process and deposition time of SiGe alloy. Exposed to the Si/Ge flux during the deposition, the interstice between non-close-packed PS spheres also filled with SiGe film.

Figure 2. SEM images of HSSAs. a, A low magnification image of monolayer HSSAs; b, High magnification bird-eye view and c, cross-sectional view images of monolayer HSSAs. d, High magnification bird-eye view and e, cross-sectional view images of double layers HSSAs. (Reprinted with permission from Ref. [31]. Copyright AIP 2007).

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Thus, in single layer each hemi-sphere shell was connected firmly to its nearest neighbors. In double layers structure, each layer was also hexagonal arrays of hemi-sphere shell, and the stacking manner of the two layers was similar to (111) plane stacking arrangement in FCC structure. Different from single layer arrays, in the upper layer double layers structure the shells were isolated to each other, while they were interconnected by flat film of SiGe in single layer arrays. The other difference was that in lower layer of double layers structure, the shell was not a complete hemi-sphere, as shown in Figure 2d and Figure 2e. As shown in Figure 2a and Figure 2b, there were some little protrudes on the surface of hemisphere shell. For sample presented in Figure 2, the deposition duration at room temperature was 30 minutes. When the deposition duration of Si/Ge alloy at room temperature was decreased to 10 minutes, HSSAs with different morphologies were obtained, as presented in Figure 3. These SEM images exhibited distinguished elongated protrudes, which had diameters of ca 70 nm. These structures will be termed nanoflower arrays in this chapter. Figure 3 gives low and high magnification images for single layer and double layer nanoflower arrays, from which morphology characters of nanoflower arrays can be found. Similar to HSSAs, single layer and double layer nanoflower arrays could be obtained. White broken circles in Figure 3a sketched up the hexagonal ordering of single layer nanoflower arrays. As marked by white arrow in Figure 3c, nanoflower from bottom layer in double layers structure were shown.

Figure 3. SEM images for single layer (a,b) and double layer nanoflower arrays.

Except for the deposition duration at room temperature, the other experimental parameters of samples presented in Figure 3 were identical to those in Figure 2. Therefore, the difference on the morphology between two samples should be caused by the deposition duration at room temperature. The longer duration at room temperature, the thicker Si/Ge alloy film would be deposited. In previous experiment, it was found that annealing at high temperature (>700 oC) would roughen the surface morphology of Si/Ge alloy film. A thinner

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film would be more susceptible to high temperature treatment, and got rougher after high temperature. We believe a rougher surface will be favorable for the growth of nanoflower arrays.

2.3. Photoluminescence Properties The room temperature photoluminescence (PL) of HSSAs and nanoflower arrays were studied by a micro-raman spectroscopy (Renishaw RM2000). The wave length of exciting laser was 514.5 nm, the power of exciting laser was 25 mW, and the spot diameter of exciting laser was 4 μm. Raman signals were measured in a back-scattering geometry and no polarizer was used. A 50x objective lens was used. The optic microscope system of this instrument was helpful to observe surface morphology and seek single layer structure region, double layer structure region and flat film region, because different surface morphologies showed different colors in the optic microscope. Aided by the optical microscope, the PL properties of single layer, double layer structure, and flat film were studied. The results of PL measurement were presented in Figure 4. Two kinds of reference samples were measured. One was a flat Si/Ge alloy film deposited under the experiment parameters identical to those of HSSAs, and the other was a flat region on the sample of HSSAs. It was shown that two reference samples showed similar PL spectra at room temperature. HSSAs and nanoflower arrays structure show stronger PL intensities and different shapes of PL spectra, compared to the reference samples. The most intensive PL spectrum came from DL-HSSAs. The spectra shapes were related to layer number of HSSAs or nanoflower arrays structures. For single layer structures the PL spectra show the most intensive peaks between 600~700 nm, while the most intensive peaks of PL spectra for double layer structure lied between 800~900 nm. Beside the spectra shapes, the intensities of PL spectra were also layer-number dependent. The double layer structure always showed PL spectra more intensive than single layer structure.

Figure 4. Room temperature photoluminescence spectra for HSSAs, nanoflower arrays, and reference samples. (Reprinted with permission from Ref. [31]. Copyright AIP 2007).

If the PL intensities of the strongest peaks were compared, the PL intensity from DLHSSAs is more than 50 folds of that from reference samples. If the PL intensities at the wave

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length of 850 nm were compared, the PL intensity from DL-HSSAs is more than 700 folds of that from reference samples. In recent publications concerning Si LED, the samples were fabricated by expensive optical- or electron beam lithography, and SOI structure was employed to facilitate the emission. The published enhancement of enhancement of Si LED varied from 70 to 250 folds [28,29,30]. Here, enhancement more than 700 folds was easy to obtain via a convenient and low cost method. Furthermore, this enhancement was obtained from a Si-substrate structure.

2.4. Discussion The obvious relationship between shapes and the most intensive peaks of PL spectra and layer number of HSSAs and nanoflower arrays implies that the enhancements from HSSAs and nanoflower arrays arise from some kind of optical effects. We will discuss in details each parameter that will affect the PL behavior of HSSAs in this section, and show that the enhancement is indeed an optical effect but not a quantum effect. For PL spectra shown in Figure 4, no character of stimulated emission was found. Thus the PL results should be studied in scenario of spontaneous emission. The luminescence intensity of an object can be described by the flowing formula: 2

I (ω ) ∝ nΓ (ω )η (ω ) = n[= 2 < f | H | i > ρ (ω )]η (ω )

equation 1

where I(ω) is PL intensity, n is density of carriers in the structure, Γ is the spontaneous emission rate, η is extraction efficiency of structure, ω is frequency of photon, f and i denote respectively initial and final state of photon, H represents the Harmiltonian of transition interaction, ρ is density of states (DOS) of emitted photons. The item in square brackets describes spontaneous emission rate given by Fermi’s Golden Rule [33]. We will discuss the difference concerning each term equation 1 between different samples, and find the origin of the enhancement in intensities and difference in PL spectra shape. Photogenerated carriers were the major source of carriers available for radiative transition in our experiment. There were three factors that induced the difference of carrier density among samples, which were the possibility that an electron would be excited to excited state under the same illumination, the volume of materials within the illumined area of exciting laser, and the interface within this area. More materials within the illumined area, more photo-carriers generated. The interface of materials affected the diffusion of photo-carriers and therefore the density of photo-carriers. Because the HSSAs and the reference samples were fabricated with the identical parameters, the microstructure of these samples should be identical, and then the possibility that an electron will be excited to excited state under the same illumination should be identical. The microstructure of HSSAs and reference sample were characterized by Transmission Electron Microscope (TEM), and the results are presented in Figure 5. The high resolution TEM (HR-TEM) images showed that Si/Ge nanocrystals were embedded in amorphous Si/Ge matrix in two samples, and the microstructure of two samples was very similar. The similarity was confirmed by counting of particle size distribution from a series of HR-TEM images, which were shown in Figure 5c, and Figure 5d.

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For the same projected illumined area, the volume of Si/Ge in DL-HSSAs and SL-HSSAs was 1.85 and 1.17 times of that in flat film sample, respectively. While the enhancement of PL spectra from SL- and DL-HSSAs exceeded 8 times of reference samples. Therefore, the extra amount of Si/Ge within illumined area could not respond for the enhancement of PL spectra.

Figure 5. HR-TEM images of HSSAs (a) and reference sample (b), and the corresponding particle size distributions of HSSAs (c) and reference sample (d). (Reprinted with permission from Ref. [31]. Copyright AIP 2007).

Interfaces were common in reference samples, HSSAs, and nanoflower arrays. Interfaces blocked the diffusion of excited carriers, or confined the excited carriers, increased the carriers’ concentration, and therefore enhanced the PL intensity. For single layer structure and reference samples, the interface that could confine the carriers was interface between film and substrate, while another interface existed in double layers arrays structure, which was the interface between unit in the upper layer and the bottom layer. The contact between nanoshells in the upper layer and the bottom layer was not so firm, as shown in Figure 2d in which the nanoshells in the upper layer were easily detached from the bottom layer. The weak contact meant that the confinement effect of excited carriers was more pronounced in double layers arrays structure, therefore the PL intensities of double layers structures should stronger than that of single layer structure and reference samples. However, carriers’ confinement could not induce the difference in shapes of PL spectra, because it increased PL intensity at all photon frequencies, as predicted by equation 1 and shown in Figure 2 of Ref. 25. Besides the carriers’ concentration, there was another electron-related factor that impacted the PL intensity in equation 1, which was the transition ratio of carriers at excited state. As shown in Figure 5, the microstructures of HSSAs and reference sample were similar,

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and thus the transition ration of excited carriers in HSSAs and reference sample were similar. Therefore, the Hamiltonian H in equation 1 was similar for HSSAs and reference sample, and the transition ratio was not the origin of PL enhancement and PL shape difference. The overall spontaneous emission efficiency, which was the term in bracket in equation 1, could be deduced from lifetime measurement of emitted photons. The lifetime measurements of photons from HSSAs and reference sample were conducted on an Edinburgh instruments (FLS 920) spectrometer. In the measurement, the excitation light source was a Xe lamp, and the samples were placed in a 45o angle both with the incident and emergent beam. The decay curves of PL intensities at 780 nm were recorded, as shown in Figure 6. Lifetimes was obtained by fitting the PL decay curves using the expression I=I0exp(-t/τ). Lifetime for shell sample was 6.36 μs and lifetime for flat sample was 6.48 μs. There was no obvious difference between lifetime for SiGe nanocrystal in shell arrays and in flat film. The measured PL lifetime (τmeasure) was defined by the relative strength of radiative (τr) and nonradiative (τnr) transitions: 1/τmeasure=1/τr+1/τnr. τnr for SiGe nanocrystal in shell arrays and flat film should be similar, because the shell arrays and the flat film were deposited with the same PLD parameter, and thus had the same microstructure, which was confirmed by TEM characterization as shown in Fig. 4. Therefore, the radiative lifetime, τr, of Si1-xGex nanocrystal in shell arrays and in flat films had no large difference. This meant that the overall spontaneous emission efficiency of Si1-xGex in HSSAs and reference samples had no obvious difference, and could not induce the PL enhancement and PL shape difference.

Figure 6. Decay spectra of PL intensities for HSSAs and reference sample. (Reprinted with permission from Ref. [31]. Copyright AIP 2007).

The carriers concentration (n) and spontaneous emission efficiency ( = 2 < f | H | i > 2 ρ (ω ) ) could not induce the shape difference of PL spectra from HSSAs and reference samples. Therefore, the enhancement and shape difference of PL spectra should be ascribed to the extraction efficiency of structure (η), as shown in equation 1. It is well known that a rough surface facilitates the extraction of photon from the surface. However, the enhancement and shape difference of PL observed in our experiment should not be simply ascribed to the roughness of surface, because it was obvious that the nanoflower arrays was rougher than HSSAs, while the PL intensity of nanoflower arrays was weaker than that of HSSAs.

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The enhancement and shape difference should originate from the arrays structures. In a two dimensional pattern structure, when the wave vector satisfied certain Bragg diffraction conditions, the lateral propagating wave could couple with ex-plane mode and then extracted from the two dimensional structure. As discussed in Ref. [34], the Γ direction in photonic band diagram was particularly important because at this direction wave propagating in six equivalent Γ-X directions coupled each other and coupled to a mode normally diffracting from surface. The normally diffracting mode remarkably enhanced the extract of light from pattern structures. Ref. 28 also showed an extraordinary enhancement of PL intensity at Γ direction of photonic band diagram. Photonic band diagram of SL- and DL-HSSAs were calculated using MEEP, an implement of three dimensional (3D) Finite-Difference Time Domain (FDTD) method [35]. Due to the limit of MEEP, the basis vector of hexagonal array was re-defined in a Cartesian coordinate system. The re-defined basis vector in patterned surface was illustrated in Figure 7a, in which a denoted center-to-center distance between nanoshells. In the 3D-FDTD 3

computation, the volume of unit cell was (1× 3 × 10)a , where (1× 3)a

2

stood for

rectangular couple of re-defined basis vector lying on the surface of HSSAs structure, and 10a was the length of vector normal to surface of HSSAs structure, as the geometric models shown in Figure 7b,c. Periodical boundary condition was adopted for vectors parallel to the surface of HSSAs, and Perfectly Matched Layers (PML) were adopted for vectors normal to the surface of HSSAs. The geometry parameters of HSSAs were measured from SEM images. The center-to-center distance a was 400 nm, the outside diameter of nanoshell was 320 nm, and the thickness of nanoshell was 80 nm.

a

b c

cd

1

1

1. Top view

a

1. Top view

a1

3 a2

2. Front view

2

3

2. Front view

2

3d view

3. Side view

3d view

3. Side view

Figure 7. a, illustration of re-defined basic vector in lateral direction, where a is original basis vector, and a1, a2 are re-defined basis vectors. b, geometric model for band structure computation of SLHSSAs. c, geometric model for band structure computation of DL-HSSAs. b and c only draw the interface of different refractive index.

Photonic band diagram of SL- and DL-HSSAs were presented in Figure 8, accompanied with the corresponding PL spectra. Hatched regions were drawn to highlight the relation between PL peaks and Γ direction in photonic band diagram. For both single layer structure and double layers structure, the wave frequencies of PL peaks were close to frequencies at Γ direction in photonic band diagram. The computation implied that the positions of peaks in PL spectra were determined by the photonic band structure in arrays structures.

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Figure 8. a, photonic band diagram of SL-HSSAs and PL spectra of single layer arrays structure; b, photonic band diagram of DL-HSSAs and PL spectra of double layer arrays structure.

However, the origin of shapes of PL spectra, or the relative intensities between peaks in PL spectra, was still not clear. Further investigation of PL properties was conducted by simulating the light propagating in and extracting out the SL-HSSAs, DL-HSSAs and flat film, and obtaining enhancement factor of SL-HSSAs and DL-HSSAs. MEEP was employed to perform the simulation. The geometric model of DL-HSSAs for the simulation was shown in Figure 9. The geometric model of SL-HSSAs was similar, except that there was only one layer of nanoshell arrays. The geometry size of simulation was 3.2x3.2x2.4 μm3, and the grid size was 25 nm. The boundary condition was PML, and the thickness of PML layer was 0.4 μm. A Gaussian pulse was used as radiating source, and was placed at the center for the flat film and center of SL- and DL-HSSAs. The radiating source was lighted, and light began to propagate in the structures. The lights extracting from the structure were recorded at 5 observation planes, which was 5 planes surrounding the blue volume in Figure 9, and were summed as PL intensity.

Figure 9. Geometric model of DL-HSSAs used in the simulation of light propagation. (Reprinted with permission from Ref. [31]. Copyright AIP 2007).

The enhancement factor was defined as the ratio of the PL intensity from the HSSAs to that from flat film. The simulation results were presented in Figure 10. It was shown by simulation that both SL-HSSAs and DL-HSSAs enhanced the extraction of photons, and that the enhancement of extraction from DL-HSSAs was stronger than that from SL-HSSAs. The

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simulation predicted that the greatest enhancement of PL from DL-HSSAs occurred at wavelength of ca. 850 nm, which matched well the experiment. The simulation results qualitatively explained relative intensity of PL between SL-HSSAs and DL-HSSAs, and predicted the wavelength of strongest PL peaks from SL-HSSAs and DL-HSSAs. As shown by simulation, the patterned surface structure remarkably enhanced the light extracting, and therefore induced enhanced light emission.

Figure 10. The simulated enhancement factor of PL. (Reprinted with permission from Ref. [31]. Copyright AIP 2007).

Each term in equation 1 was discussed, and the origin of enhancement and shape difference of PL spectra was ascribed to the difference in the abilities of extracting light from difference structures. The reference samples, flat films, suffered the internal reflection which blocked the light extraction in many high refractive index LED device, and therefore had low extraction efficiency. After patterned into hemisphere shell or nanoflower arrays, some lateral propagating light waves coupled each other and coupled to light wave normally extracting from the surface. Thus, periodically patterned surface facilitated the light extraction. In nanoflower arrays, the surfaces were rough, the symmetry of arrays was disturbed by the elongated protrudes on the surfaces of repeating unit. Accordingly, the couples in laterally and normally propagating light wave were weaker than those in HSSAs. Therefore, the PL intensities of nanoflower arrays were weaker than those of HSSAs, while they have very similar shape of PL spectra. Double layers structures could extract more light, because more couple of light waves took place. In conclusion, light extraction ability was important for the PL properties. It was also emphasized by Yablonovitch et al. that “In most cases of room temperature LED’s, it appears that the light extraction effect, will be more important.” (than Purcell spontaneous emission rate enhancement ) [36].

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3. Controllable Multilayer Hemisphere Shell Arrays on Different Substrate 3.1. Background As introduced in section 2, HSSAs remarkably enhanced the PL intensities of Si/Ge alloy, and DL-HSSAs showed enhancement far beyond SL-HSSAs. However, it is lack of convenient method to control layer number in PS arrays. So to obtain DL-HSSAs, great efforts have to pay to prepare double layer PS arrays as template. Further more, the results introduced in section 2 imply that the enhancement of PL intensity is closely related to layer number in HSSAs, and that more remarkable enhancement can be obtained if further increase the layer number. In this section, we will introduce the preliminary results concerning the fabrication of layer-number-controllable multilayer hemisphere shell arrays (ML-HSSAs) and corresponding PL properties. Our attempts followed a layer-by-layer assembly scheme, which had been successfully demonstrated on incorporating defects into 3d colloid crystal [37,38,39]. The key point of layer-by-layer assembly is to control layer number of arrays in each round of assembly. Our effort is schematically presented in Figure 11. After self-assembly by a simple drop coating method, PS arrays might showed different layer numbers at different region, as shown in Figure 11a. Without etching by RIE process, the PS arrays were used as scaffold to deposit a thin layer of Si/Ge by PLD at room temperature. After this process, only the utmost top surface of PS arrays was covered by Si/Ge, as shown in as shown in Figure 11b. Thus, only one layer of Si/Ge alloy was obtained despite what is the layer number of PS arrays. In succession, the PS arrays were removed by heating at high temperature, 700 oC for example. After the removal of PS arrays, SL-HSSAs were obtained, as shown in Figure 11c. Finally, Si/Ge alloy was deposited at desired temperature for appropriate duration to demanded thickness, as shown in Figure 11d.

Figure 11. Experimental scheme for the control of layer number in an assembly round.

The sample after procedures listed in Figure 11 was used as substrate for the selfassembly of PS sphere, and then began to be initiated another round of fabrication. Because the layer number of HSSAs can be controlled in single round, the overall layer number of HSSAs is well controlled just by repeating the round for desired repeats.

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3.2. Geometry Characterization Surface morphology of SL-HSSAs fabricated by the new scheme was characterized by SEM, and the results were presented in Figure 12. SEM images were taken from the sample after the removal of PS by heating. Large area homogeneous ordered thin Si/Ge HSSAs were obtained, as shown in Figure 12a. Different from what was introduced in section 2, nanoshells fabricated by new scheme were close packed. Figure 12b shows cross-sectional view SEM image of SL-HSSAs, and Figure 12c shows a high resolution image corresponding to region enclosed by white rectangle in Figure 12b. It is shown by Figure 12c that the PS sphere had been totally removed, and the thickness of nanoshell was less than 10 nm.

Figure 12. SEM images of thin Si/Ge alloy HSSAs.

Figure 13. Cross-sectional view SEM image of FL-HSSAs (a,b,c) and the formation scheme of AB stacking (d,e,f) and AA stacking.

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Four-layers HSSAs (FL-HSSAs) had been fabricated by repeating 4 times procedures shown in Figure 11. Figure 13a-c presented cross-sectional SEM images of FL-HSSAs. As shown by SEM images, the layer number was homogeneous in the samples. There were two kinds of stack manner between nanoshell in adjacent layers, which were AB type as shown in Figure 13b and AA type as shown in Figure 13c. Because each layer of nanoshell was a replica of top surface of PS arrays, so the stacking manner between adjacent layer of nanoshells was determined by the relative position between the top layer PS spheres in PS arrays and the latest layer of nanoshell. If the top layer PS spheres were located at the interstice of nanoshells after self-assembly, as shown from Figure 13d to Figure 13e, the latest deposited nanoshells also located at the interstice of nanoshells in previous nanoshell arrays, as shown from Figure 13e to Figure 13e. Figure 13f to Figure 13g illustrated the process leading to AA type stacking.

3.3. Photoluminescence Properties The photoluminescence properties of FL-HSSAs were investigated by a micro Raman spectroscopy. The details about experiment setup and parameters were the same as what introduced in section 2.3. For the evaluation of enhancement properties, SL-HSSAs were fabricated with the same experimental parameter as each single round for FL-HSSAs. The results were presented in Figure 14. It was shown that the PL intensities of FL-HSSAs were about 7 times higher than those of SL-HSSAs.

Figure 14. Photoluminescence spectra of FL-HSSAs and SL-HSSAs on Si substrate (The sharp peaks marked by the circle are Raman scattering peaks).

The amount of Si/Ge alloy in FL-HSSAs within the projected illumined area during PL measurement was four times that of SL-HSSAs. Thus the actually enhancement arisen from structure difference was about 2 times. The enhancement in this approach was lower than what was introduced in section 2. There were two possible reasons. The structural difference may be one reason. As shown by SEM images, in ML-HSSAs, the hemisphere nanoshell is closed-packed to each others, while the hemisphere nanoshell is non-close-packed in samples introduced in section 2. The variation of refractive index was more pronounced in non-closepacked arrays, and had more powerful ability to tune the optical properties. This opinion was also implied by the difference between positions of strongest peaks from close-packed SL-

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HSSAs, as shown in Figure 14, and non-close-packed SL-HSSAs, as shown in Figure 4. The disorder in the stacking manner between adjacent layers in ML-HSSAs, as shown in Figure 13, can also contribute to the low enhancement ability.

3.4. ML-HSSAs on Different Substrates All samples introduced above were fabricated directly on Si substrate, which has a relative high refractive index. The higher refractive index of substrate makes smaller the total reflection angle of substrate, and makes lower the extraction efficiency of device. Thus, many published Si LED devices were fabricated on a low refractive index SiO2 film [28,29,30]. We also tried this scheme to further enhance the luminescence behavior of HSSAs. Multi-layers SiO2 nanosphere arrays were used as substrate for the fabrication of MLHSSAs. The SiO2 nanosphere arrays self-assembled to substrate often show (111) direction normal to the substrate surface. Because SiO2 nanosphere array has an optical band gap along (111) direction, the crystalline orientation of SiO2 nanosphere array relative to substrate helps to block the light propagating normally deep into the substrate and facilitate the extraction of light from the substrate [40].

Figure 15. Plan view (a), low magnification (b) and high magnification (c) cross-sectional view SEM image of FL-HSSAs.

The SiO2 nanospheres had a nominal diameter of 500nm, and were assembled to arrays by drop-coating technique. SEM images of FL-HSSAs on SiO2 nanosphere arrays were shown in Figure 15. Region 1 in Figure 15a showed top surface of FL-HSSAs, while region 2 showed SiO2 nanosphere arrays. Because diameter of SiO2 nanosphere was larger than that of PS sphere used as mask for the deposition of Si/Ge nanoshell, the ordering of PS sphere on SiO2 nanosphere arrays was a little bit bad, as shown in Figure 15a. From the cross-sectional view image, Figure 15b, different layers of structure was easy to be distinguished, which were 1 for Si substrate, 2 for SiO2 nanosphere arrays, and 3 for FL-HSSAs. SEM images showed

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that multi-rounds of fabrication for ML-HSSAs did not disturb the original ordering of SiO2 nanosphere arrays and the ordering of HSSAs fabricated by previous step. PL spectra of FL-HSSAs on Si substrate and SiO2 nanosphere arrays substrate were presented in Figure 16. The SiO2 nanosphere substrate increased 4 times the PL intensities. It should be noticed that the order of FL-HSSAs on SiO2 nanosphere arrays was not good and the diameter of SiO2 nanosphere had not been designed to open the band gap at the wave length coincided to the strongest peaks of FL-HSSAs. The order of FL-HSSAs can be improved by using PS nanosphere and SiO2 nanosphere with the same diameter. Combining the practicable design of photonic band gap, the enhancement effect can be further improved.

Figure 16. PL PL spectra of FL-HSSAs on Si substrate and SiO2 nanosphere arrays substrate.

4. Conclusion In this chapter, we introduce the fabrication and the PL properties of HSSAs. Remarkably enhancement, as high as 700 folds, can be obtained by DL-HSSAs. By the TEM, SEM, photon lifetime measure, and electromagnetic simulation, we ascribe the origin of PL enhancement to the improvement of extraction efficiency in HSSAs. Basing on this finding, we conducted preliminary research on the fabrication and PL properties of controllable MLHSSAs. The enhancement of PL from controllable ML-HSSAs proved the practicability of this idea. We also tried to incorporate controllable ML-HSSAs to SiO2 nanosphere arrays substrate, and obtained 4 times further enhancement of PL intensity. The results presented in this chapter shows that HSSAs is a promising route for Si based LED device.

Acknowledgement The authors are grateful to National 973 Project of China, Chinese National Nature Science Foundation and National Center for Nanoscience and Technology of China.

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[25] Choi, S. H.; Kim, J. N.; Kim, H. Y.; Hong, Y. K.; Koo, J. Y.; Seok, J. H.; Kim, J. Y. Appl. Phys. Lett. 2002, 80, 2520-2522. [26] Huang, Z. P.; Wu, Y.; Fang, H.; Deng, N.; Ren, T. L.; Zhu J. Nanotechnol. 2006, 17, 1476-1480. [27] Green, M. A.; Zhao, J.; Wang, A.; Reece, P. J.; Gal, M. Nature 2001, 412, 805-809. [28] Zelsmann, M.; Picard, E.; Charvolin, T.; Hadji, E.; Heitzmann, M.; Dal’zotto, B.; Nier, M. E. Appl. Phys. Lett. 2003, 83, 2542-2544. [29] Galli, M.; Politi, A.; Belotti, M.; Gerace, D.; Liscidini, M.; Patrini, M.; Andreani, L. C.; Miritello, M.; Irrera, A.; Priolo, F.; Chen, Y. Appl. Phys. Lett. 2006, 88, 251114. [30] Cluzel, B.; Pauc, N.; Calvo, V.; Charvolin, T.; Hadj, E. Appl. Phys. Lett. 2006, 88, 113120. [31] Huang, Z. P.; Zhu, J. Appl. Phys. Lett. 2007, 91, 013108. [32] Gates, B.; Park, S. H.; Xia, Y. N. Adv. Mater. 2000, 12, 653-656. [33] Scully, M.; Zubairy Quantum Optics. Cambridge University Press:Cambridge. 1997. [34] Masahiro, I.; AlongKarn, C.; Susumu, N.; Masamitsu, M. Phys. Rev. B 2002, 65, 195306. [35] Farjadpour, A.; Roundy, D.; Rodriguez, A.; Ibanescu, M.; Bermel, P.; Joannopoulos, J. D. ; Steven G. J.; Burr, G. Optics Lett. 2006, 31, 2972–2974. [36] Boroditsky, M.; Vrijen, R.; Krauss, T. F.; Coccioli, R.; Bhat, R.; Yablonovitch, E. J. Lightware Technol. 1999, 17, 2096-2112. [37] Vekris, E.; Kitaev, V.; Fremann, G. V.; Perovic, D. D.; Aitchison, J. S.; Geoffrey A. O. Adv. Mater. 2005, 17, 1269-1272. [38] Yan, Q. F.; Zhou, Z. C.; Zhao, X. S.; Chua, S. J. Adv. Mater. 2005, 17, 1917-1920. [39] Raul, P.; Agustin, M.; Manuel O.; Hernan, M. Adv. Mater. 2006, 18, 1183-1187. [40] Reynolds, A.; Lopez-Tejeira, F.; Cassagne, D.; Garcia-Vidal, F. J.; Jouanin, C.; Sanchez-Dehesa, J. Phys. Rev. B. 1999, 11422-11426.

In: Nanotechnology… Editors: C.J. Dixon and O.W. Curtines, pp. 479-508

ISBN: 978-1-60692-162-3 © 2010 Nova Science Publishers, Inc.

Chapter 15

A STRATEGY FOR THE INCORPORATION OF TRIVALENT LANTHANIDE IONS INTO ANATASE TIO2 NANOCRYSTALS Wenqin Luo,a, b Chengyu Fu,a, b Renfu Lia, b and Xueyuan Chen1,a, b a

Key Laboratory of Optoelectronic Materials Chemistry and Physics, Fujian Institute of Research on the Structure of Matter, Chinese Academy of Sciences, Fuzhou, Fujian 350002, P.R. China b State Key Laboratory of Structural Chemistry, Fuzhou, Fujian 350002, P.R. China

Abstract Trivalent lanthanide (Ln3+) ion-doped semiconductor nanocrystals have attracted extensive attention due to the ability to tailor their optical properties via size control and to achieve highly efficient luminescence through sensitization by the host. To date, finding a way to dope the “undopable” Ln3+ ions into semiconductor nanomaterials via chemical methods remains a challenge. In this chapter, recent progress in the doping of Ln3+ ions in TiO2 nanomaterials has been reviewed. A novel sol-gel-solvothermal method has been developed to effectively incorporate Ln3+ ions (Eu3+, Er3+, Nd3+ and Sm3+) into anatase TiO2 nanoparticles via the selfassembly and crystallization process of previous amorphous nanoparticles, in spite of a large mismatch in ionic radius and charge imbalance between Ln3+ and Ti4+. The crystallization process of Ln3+ doped TiO2 nanoparticles were systematically studied by means of thermogravimetric-differential thermal analyses (TG-DTA), powder X-ray diffraction (XRD), and transmission electron microscope (TEM). Photoluminescence (PL) spectra of Ln3+:TiO2 samples exhibit resolved and sharp emission and excitation lines from the intra f-f transitions of Ln3+ ions (Ln=Nd, Sm, Eu, Er), indicating regular crystalline surroundings of Ln3+ ions. Multiple sites of Eu3+, Sm3+ and Nd3+ ions in anatase TiO2 were detected by means of highresolution site-selective spectroscopy at 10K, whereas only single site emission of Er3+ in TiO2 were observed. Very intense near-infrared luminescence around 1.53 µm was also observed, which originated from the single lattice site of Er3+ ions incorporated in TiO2 nanocrystals. The luminescence dynamics and CF levels of Ln3+ at different sites have been analyzed. Highly efficient emissions of Nd3+ and Sm3+ sensitized by the TiO2 host were observed upon the excitation above the TiO2 band gap energy at room temperature (RT), 1

E-mail address: [email protected]. Author to whom correspondence should be addressed.

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1. Introduction Recently, Ln3+ ion-doped semiconductor nanocrystals have attracted extensive attention due to their potential applications in photoelectric devices, flat plane displays, solid state laser materials and bio-labels [1-6]. One of the advantages offered by these materials is the ability to tailor their optical properties via size control and to achieve highly efficient luminescence through the sensitization by the host. Moreover, they provide the possibility of excitation with electrical current. Titania is a well-known wide band-gap semiconductor (band-gap of 3.2 eV for anatase [7]) and a good candidate to be used as the host material of Ln3+ ions because of its good chemical, mechanical, optical and thermal properties. There are three crystalline forms of TiO2, i.e., anatase (tetragonal), rutile (tetragonal) and brookite (orthorhomibic). Anatase is a main phase of TiO2 when synthesized via soft chemistry from titanium tetravalent precursors. After high temperature calcination, anatase phase can be transformed to rutile, which is the most stable phase of titania. However, due to a large mismatch in ionic radius between Ln3+ and Ti4+ (0.086–0.103 nm vs. 0.061 nm, respectively, for coordination number VI [8]), charge imbalance and the existence of intrinsic self-purification processes at the nanoscale [9-10], it is difficult to incorporate Ln3+ ions into the TiO2 nanocrystal lattice through chemical methods. The means to effectively accommodate Ln3+ ions in TiO2 systems is a precondition to achieving high luminescence performance of this kind of material. In the past few years, many efforts have been made to synthesis Ln3+ ion-doped TiO2 nanomaterials in order to yield intense and host-sensitized Ln3+ emissions in various forms of the materials [11-22]. Although whether Ln3+ ions could be incorporated into the semiconductor lattice is controversial, some important results have been achieved. These results indicated that trivalent Ln3+ ions can be introduced into the TiO2 nanocrystals, and energy transfer (ET) from the host to Ln3+ ions can be achieved if the synthesis method is well designed. In this chapter, different chemical approaches to accommodate Ln3+ ions in anatase TiO2 host have been reviewed. Much attention is focused on a new approach recently developed by us to incorporate Ln3+ in TiO2 nanoparticles lattice. The crystallization behavior of Ln3+ doped TiO2 nanoparticles was systematically studied. Recent progress on the PL properties of Ln3+ ions (Ln=Eu, Sm, Nd and Er) in anatase TiO2 nanocrystals prepared by sol-gelsolvothermal method has been reviewed. A mechanism for the incorporation of Ln3+ ions into TiO2 lattice is presented.

2. Progress in the Doping of Lanthanides in TiO2 Nanomaterials 2.1. Amorphous-Crystalline Composites Amorphous phase, such as glass, is a relatively loose structure. The solid solubility of Ln ions in amorphous phase is much larger than that in the crystalline counterpart. Moreover, the lattice distortion caused by the replacement of Ti4+ ions with the larger Ln3+ ions and the charge imbalance can be easily compensated in amorphous phase. The 3+

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amorphous-crystalline composite material is a promising host for Ln3+ ions with efficient PL, in which the amorphous phase provides an ideal environment to accommodate Ln3+ ions and the crystallite with semiconducting nature acts as a sensitizer (energy antenna).

Figure 1. TEM and EDS of Eu3+-doped (8 mol%) cubic mesoporous titania thin films: a) view along the [100] zone axis; b) magnified view along the [100] zone axis, c) view down the [111] zone axis, d) EDS confirming the presence of europium in the mesopore walls. The measured europium content is 12±4 mol%; this is an average value taken over several large areas. (Reprinted with permission from [12], Frindell, K. L. et al., Angew. Chem. Int. Ed. 41, 959 (2002). ©2002, Wiley-VCH Verlag GmbH & Co. KGaA.)

A cubic mesostructured matrix of titania with three dimensional array of embedded anatase nanocrystals has been utilized to accommodate Ln3+ ions, and the sensitized Ln3+ emissions by TiO2 host were obtained [12, 23]. Figure 1 shows the TEM images of Eu3+ (8 mol%) doped cubic mesoporous titania thin films. The TEM images viewed along different directions ([100] and [111]) revealed that the films have a cubic arrangement of pores with a size of 7-8 nm. As shown in Figure 1d, the high content (12±4 mol%) of Eu3+ present in the walls of the sample is confirmed by energy dispersive X-ray spectroscopy (EDS) measurements. The excitation and emission spectra of Eu3+ (8 mol%) and Sm3+ (2 mol%) ions doped mesoporous titania films are shown in Figure 2a, b and c, d respectively [23]. Upon excitation above titania band gap, both samples exhibited intense red emissions due to

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intra-4f transitions of Eu3+ and Sm3+ ions. The obtained emission lines are unresolved as a result of the location of Ln3+ ions in the amorphous titania regions. The excitation spectra of both samples by monitoring the emissions at 614 nm for Eu3+ ions and 601 nm for Sm3+ ions exhibited a broad band at 330 nm, which was associated with the band gap of titania films, whereas excitation lines from intra-4f transitions of Eu3+ and Sm3+ can not be detected, indicating that the main contribution to the excitation is from the band gap of the titania.

Figure 2. Excitation and PL spectra of (a, b) 8 mol% europium and (c, d) 2 mol % samarium doped mesoporous titania films. (Redrawn after [23], Frindell, K. L. et al., J. Solid State Chem. 172, 81 (2003). © 2003, Elsevier).

Similar amorphous-crystalline structure was also adopted to accommodate Eu3+ ions in a monodisperse spherical mesoporous TiO2 particles prepared by a simple nonionic surfactantassisted soft-chemistry method [19]. Figure 3 shows the SEM and TEM images of Eu3+ doped TiO2 nanoparticles. The SEM image (Figure 3a) reveals that the particles are spherical and monodisperse with a mean size of 250 nm. The presence of Eu3+ (4.95 mol%) was confirmed by EDS in the inset of Figure 3a. The TEM image illustrated that the Eu3+ doped TiO2 nanoparticles possess mesoporous structure with the pore size of 7–10 nm. The high resolution TEM (HRTEM) image (Figure 3c) reveals that the pore wall of the mesoporous nanoparticles is semicrystallite and there are many anatase nanocrystallites embedded in the amorphous titania regions. The RT PL spectra of the Eu3+ doped nanoparticles under different heat treatments are compared in Figure 4. Under ultraviolet excitation at 360 nm, broad emission lines from the excited state of 5D0 to 7F1, 7F2 and 7F3 states can be observed. As shown in Figure 4, the sample calcined at 400 °C shows more effective luminescence compared to the as-prepared sample without further calcination. An important factor for the increase of sample’s luminescence intensity may be the formation of titania nanocrystallites in the pore walls under calcination which plays the role of sensitizer to transferring the absorbed excitation energy to Eu3+ ions. When the sample is further annealed at 500 °C, the PL intensity was found to decrease. This may be due to the fact that the amorphous titania were completely transformed into anatase and thus the previously well dispersed Eu3+ ions tended to aggregate and the PL of Eu3+ was quenched.

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Figure 3. (Color online) (a) SEM image and EDS spectra, (b) low magnification TEM, and (c) HRTEM images and corresponding selected area electron diffraction (SAED) pattern of monodisperse mesoporous Eu3+ doped TiO2 phosphor particles after 400 °C calcinations. (Reprinted with permission from [19], Yin, J. B. et al., Appl. Phys. Lett. 90, 113112 (2007). © 2007, American Institute of Physics).

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Figure 4. (Color online) RT PL spectra of the mesoporous Eu-doped TiO2 phosphor particles (a) before calcinations, (b) after 400 °C calcinations, and (c) after 500 °C calcinations. (The inset is the emission photograph of the corresponding samples.) (Reprinted with permission from [19], Yin, J. B. et al., Appl. Phys. Lett. 90, 113112 (2007). © 2007, American Institute of Physics).

2.2. Aggregates of Layered Nanostructure Layered oxide materials are repeating two dimensional units of host oxide layers between which a guest cation exists to hold the layers together. The unique structure of the layered oxides makes them an ideal host to accommodate the Ln3+ ions. In this system, Ln3+ ions are not really incorporated into the TiO2 lattice. As a consequence, no crystal-field (CF) transition of Ln3+ ions can be observed in PL spectra. Xin et al. reported an exfoliation-restacking route to accommodate Ln3+ ions in the lamellar aggregates of titania nanosheets [24]. The TEM image (Figure 5a) clearly shows the lamellar structure of these materials. The content of Eu3+ ions in the flocculated sample can be detected by EDS analysis (Figure 5b) and the concentration of the dopant can be roughly determined to be 10±1 mol%. The PL spectra of ex-Ti0.91O2/H, ex-Ti0.91O2/Tb and eTi0.91O2/Eu (“ex- ” means “exfoliated”) samples are shown in Figure 6. Under the excitation at 250 nm, ex-Ti0.91O2/Eu sample emits lights both from the Ti0.91O2 host and Eu3+ ions, whereas only Ti0.91O2 host emission is detected for ex-Ti0.91O2/Tb sample. Xin et al. ascribed the broad band peaking at 250 nm (Figure 6c) to the Ti0.91O2 host absorption. However, the band gap of Ti0.91O2 host was determined to be 380 nm by UV/VIS absorption spectra, which deviates significantly from 250 nm. Instead, the charge transfer between Eu3+ and O2- ions, i.e., an electron transferring from the O2- (2p6) orbital to the empty orbital of Eu3+ (4f6), usually locates at 200−290 nm [25]. Therefore it is likely that the observed excitation band between 220 and 380 nm results from the Eu3+−O2- charge transfer in ex-Ti0.91O2/Eu. Their conclusion on the ET from Ti0.91O2 nanosheet to Eu3+ ions on the basis of the observation of

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emissions of Eu3+ under the excitation at 250 nm seems questionable and deserves further evidence.

Figure 5. (a) TEM image and (b) EDS spectra of the composite ex-Ti0.91O2/Eu. (Reprinted with permission from [24], Xin, H et al., Appl. Phys. Lett. 85, 4187 (2004). © 2004, American Institute of Physics)

Layered titanate oxides intercalated with hydrated Eu3+ ions have been synthesized by electrostatic self-assembly deposition (ESD) and layer-by-layer assembly (LBL) method [2627]. Figure 7 shows the XRD pattern of an as-deposited Eu/TiO film by ESD method, where Eu3+ ions and water molecules are sandwiched between titanate nanosheets. The interlayer distance of TiO is 6.8 Å. The composition with content was determined to be Eu0.31Ti0.81O4·2.1H2O, indicating that the intercalated Eu3+ ions exist as aqua ions and are coordinated with 7−10 water molecules under ambient conditions. ET from TiO layer to intercalated Eu3+ ions has been achieved in this material. The interlayer water that contributed to ET plays a key role in the sensitized emission of Eu3+ ions. As shown in Figure 8, the excitation spectra by monitoring the 5D0→7F2 transition at 614 nm exhibit a broad band at

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250-350 nm, which was associated with the band gap of the titanate nanosheet. In addition, the peak at 395 nm due to the 7F0→5L6 intra-4f transition of Eu3+ ions is also observed. Upon excitation above TiO band gap at 300 nm, emission lines at 570, 593 and 614 nm assigned to the transitions from 5D0 to 7F0, 7F 1 and 7F 2 are detected, respectively. The intensities of excitation and emission peaks of the as-deposited Eu/TiO film are stronger than that of Eu/TiO film treated at 100 °C for 1 h due to the elimination of some of interlayer water by heat treatment. Emission intensity and layer distance are also found to decrease with the decrease in humidity for the same reason as illustrated in Figure 8.

Figure 6. PL excitation and emission spectra of (a) ex-Ti0.91O2/H, (b) ex-Ti0.91O2/Tb and (c) ex-Ti0.91O2/Eu. The excitation spectra were measured by monitoring the emissions at 395, 544 and 612 nm for ex-Ti0.91O2/H, ex-Ti0.91O2/Tb and ex-Ti0.91O2/Eu, respectively. The excitation wavelengths are also shown. (Reprinted with permission from [24], Xin, H et al., Appl. Phys. Lett. 85, 4187 (2004). © 2004, American Institute of Physics).

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Figure 7. XRD pattern and structure model of as-deposited titanate layered oxide intercalated with Eu3+ ions. (Reprinted with permission from [27], Ida, S. et al. J. Phys. Chem. B 110, 23881 (2006). © 2006, American Chemical Society).

Figure 8. RT excitation (λem=614 nm) and emission (λex=300 nm) spectra of Eu/TiO films of asdeposited, treated at 100 oC for 1 h and under 5% humidity for 2 days. (Reprinted with permission from [27], Ida, S. et al. J. Phys. Chem. B 110, 23881 (2006). © 2006, American Chemical Society).

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2.3. Sol-Gel Thin Films Sol-gel method is commonly used to prepare Ln3+ ions doped TiO2 thin films due to the possibility of architectural design of coating and multilayered thin films in several types of geometries [11, 14, 28-29]. The substrate of the thin films may significantly affect the crystallization and PL performance of the samples [11, 30]. Figure 9 shows the XRD patterns of TiO2 thin film deposited on glass and Si respectively. For thin films deposited on glass, a broad diffraction peak at around 26° is observed, indicative of amorphous nature for the sample. In contrast, the thin film deposited on Si substrates exhibits diffraction peaks of a

Figure 9. XRD of TiO2 thin film deposited on (a) glass and (b) Si before and after annealing at 600 °C. Films deposited on Corning glass substrates are amorphous, while those deposited on Si present a mixture of crystalline phases. After annealing, the rutile phase (see the peaks at * positions) is better developed in films deposited on Si. (Redrawn after [30], Palomino-Merino, R. Thin Solid Films 401, 118 (2001). © 2001, Elsevier).

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mixture of crystalline phases (anatase, rutile and TiO). With the increase of annealing temperature at 600 °C, no improvement was achieved for the crystallinity of amorphous thin film deposited on glass, while the crystallinity of the films deposited on Si was improved slightly. Upon the excitation by a 325 nm Ar laser beam, Eu3+ ions embedded in both samples exhibit broad emission lines from the 5D0→7FJ (J=0, 1, 2, 3, 4). Moreover, for the same nominal concentration of Eu3+ and similar film thickness, a higher PL intensity is observed for film deposited on Si than on glass substrate. A possible explanation of this effect was ascribed to the non-negligible contribution from the reflected light of the PL signal at the TiO2/Si interface, which was not present in the transparent TiO2/glass interface [11]. However the contribution of better crystallinity of the TiO2 on silicon to the improvement of PL intensity should also be taken into account. It was pointed out that the PL intensity of Ln3+ in thin films can be enhanced by codoping some foreign cations such as Ce3+ and Y3+ [28]，which act as dispersers of the dopants and thus reduce the possibility for the cross relaxation of close Ln3+ ions. Strong visible luminescence of Tb3+ ions due to intra-4f shell transitions were observed for Tb3+ doped titania films under the excitation at 340 nm line of Xe lamp [28-29]. The excitation spectrum presents a broad line ranging from 340–390 nm with the strongest peak at 370 nm, which is ascribed to the band gap absorption of TiO2 host, suggesting that the emissions of Tb3+ are mainly achieved via ET from TiO2 host to Tb3+ ions. The PL intensity of Tb3+ is significantly enhanced by co-doping Ce3+ ions (2.6 mol%) in the thin film. The Ce3+ ions bridged the ET from TiO2 host to Tb3+ ions, i.e., the energy absorbed by TiO2 host is transferred to Ce3+ ions firstly and then to Tb3+ ions. Furthermore, the co-doped Ce3+ ions can disperse the clustering of Tb3+ ions, prevent nonradiative de-excitation among Tb3+ ions and enhance the PL intensity of Tb3+. The co-doping of Y3+ ions was also found to enhance the PL intensity of Er3+ doped TiO2 thin film [29]. With the increase of the co-doped Y3+ ions (030 mol%), the PL intensity of Er3+ ions at around 1.54 μm was gradually enhanced, meanwhile the emission lines of Er3+ were broadened as a result of worse crystallinity of TiO2 host induced by the co-doping of Y3+ ions.

2.4. Nanocrystals and Nanotubes Besides the different ways to accommodate Ln3+ ions in TiO2 materials mentioned above, Ln ions were also doped into the nanocrystals [13, 18, 20-21, 31] and nanotubes [25, 32]. Stouwdam and coworkers reported the sensitized emissions in Ln3+ (Ln=Eu, Er, Nd, Yb) ions doped TiO2 nanoparticles with a size of 3-5 nm prepared by the thermal decomposition of [Ti(OiPr)3(dmae)] (ipr=isopropyl; dmae=dimethylaminoethoxide) in hot trioctylphosphine oxide (TOPO). The sensitized emissions of Sm3+ in TiO2 nanoparticle aggregates prepared by sol-gel method were also addressed [13]. As shown in Figure 10, the excitation spectra show a broad band centered at 358 nm, which is assigned to titania band gap absorption, thus verifying the ET from TiO2 to Sm3+ ions. The relatively weak band at 240 nm in the spectra can be attributed to the charge transfer transition from the O2- ligand to the Sm3+ ions. The band gap of TiO2 can be tailored by co-doping Bi3+ and Zr4+ ions. Upon excitation above the TiO2 band gap at 355 nm, slightly resolved emission lines from 4G5/2 state to 6H5/2,7/2,9/2 multiplets are detected, indicating that Sm3+ ions were probably located at a quite regular CF environment. 3+

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Recently, Wang et al. synthesized Eu3+ doped nanocrystalline titania microspheres by ultrasonic spray pyrolysis (USP) and solvent evaporation-induced self-assembly method [20]. As shown in Figure 11, the TiO2:Eu3+ (8 mol%) sample heated at 400 °C shows spherical morphology with diameters ranging from 0.1 to 1.4 μm. The spherical particles exhibit porous structures (Figure 11b) and anatase polycrystalline nature (Figure 11c). Emission bands centered at 590, 614, 651 and 698 nm corresponding to the characteristic 5D0→7FJ (J=1, 2, 3, 4) transitions of Eu3+ are observed under the excitation at 330 nm. These lines are inhomogeneously broadened suggesting that Eu3+ ions were located at much distorted CF environments. Due to the porous structure of the sample, high concentration of doped Eu3+ ions up to 16 mol% can be achieved. The excitation spectra by monitoring the 5D0→7F2 transition at 614 nm exhibit a broad band at around 330 nm which can be associated with the band gap absorption of TiO2 host, verifying effective ET from TiO2 host to Eu3+ ions.

Figure 10. (a) Excitation spectra (λem=610 nm) for (i) TiO2:Sm3+ (0.75 mol%), (ii) TiO2:Sm3+ (0.75 mol%), Zr3+ (0.5 mol%), and (iii) TiO2:Sm3+ (0.75 mol%), Bi3+ (0.5 mol%). (b) Emission spectrum of TiO2:Sm3+ (λex=355 nm). (Redrawn after [13], Hu, L. Y. J. Lumin. 127, 371 (2007). © 2007, Elsevier).

Intense PL of Eu3+ ions was obtained in TiO2 nanotubes prepared by hydrothermal method in NaOH aqueous solution with the precursor of Eu3+:TiO2 nanocrystals obtained by the sol-gel method [25, 33]. Figure 12a shows the TEM image of the obtained nanotubes. It can be seen that the outer diameters and lengths of the nanotubes are about 9 and 200 nm, respectively. By further annealing at 700 oC, as shown in Figure 12b, the nanotubes were transformed to nanorods. The excitation spectra of precursor TiO2/Eu nanocrystals, nanotubes and nanorods are compared in Figure 13. By monitoring the 5D0→7F2 emission of Eu3+ ions at 612 nm, the excitation lines at about 394, 465 and 535 nm corresponding to the typical Eu3+ ff transitions can be observed in all three samples. However, the shapes and positions of excitation lines were quite different from each other. In addition to the above Eu3+ characterized absorption lines, two intense broad excitation peaks with the maxima at about

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275 and 312 nm are observed in Eu3+:TiO2 nanotubes, which can be ascribed to the charge transfer between Eu3+ and O2- and the band gap absorption of titania host respectively, illustrating the existence of ET from the TiO2 nanotubes to Eu3+ ions. In this preparation method, the TiO2 precursors may have great impact on the PL properties of final nanotubes. The PL of titania nanotubes prepared via a two-step hydrothermal treatment [32] behaved very differently from that of the counterparts above-mentioned, in that the ET from TiO2 to Eu3+ ions was not observed and the PL lines were worse resolved.

Figure 11. Eu-doped nanocrystalline titania microspheres. a) SEM image. b) TEM image. c) SAED pattern recorded on the sphere shown in b). The sample shown in a), b), and c) contains 8 mol % Eu. d) PL spectra of the spheres doped with varying amounts of Eu. The excitation is at 330 nm. e) Corresponding PL excitation spectra for the 614 nm emission. The samples were thermally treated at 400 °C. (Reprinted with permission from [20], L. Li et al., Adv. Mater. 20, 903, (2008). ©2008, Wiley-VCH Verlag GmbH & Co. KGaA.)

Figure 12. TEM images of Eu3+:TiO2: (a) nanotubes, (b) nanorods (Redrawn after [25], Zeng, Q. G. et al., Scripta Mater. 57, 897 (2007). © 2007, Elsevier).

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Figure 13. Excitation spectra of the Eu3+ ions in TiO2 nanotubes, nanorods and nanocrystals by monitoring the 612 nm emission line. The inset is the ultraviolet–visible spectrum of the nanotubes. (Redrawn after [25], Zeng, Q. G. et al., Scripta Mater. 57, 897 (2007). © 2007, Elsevier).

3. Sol-Gel Solvothermal Approach for Incorporation of Lanthanides in TIO2 Nano-Lattices 3.1. Synthesis and Characterization Experiments of Ln3+:TiO2 Nanoparticles Recently, we have developed a facile sol-gel solvothermal approach to incorporate Ln3+ ions into anatase TiO2 nano-lattices [34-35]. In this method, we take advantages of the hydrolysis and condensation reactions of tetra(n-butyl)titanate to precipitate TiO2 nanoparticles from the solution in the absence of catalyzers such as acid or alkali. During the sol-gel process, Ln3+ ions can be co-precipitated with the TiO2 nanoparticles. In detail, 1 mL tetra(n-butyl)titanate was dissolved in 20 mL ethanol, to which was added the required amount of Ln(CH3COO)3·6(H2O) solution in 0.2 mL diluted de-ionized water and 20 mL ethanol. After stirring for about 30 min, white particles began to precipitate from the solution. Under continuous stirring for 3 h, the obtained cloudy solution was transferred into 20 mL Teflon-lined autoclaves and subjected to solventhermal treatment for 5 h at 120 °C. The final obtained precipitates were isolated by centrifugation, washed with ethanol several times, and dried at 40 °C. The as-prepared sample was further annealed in air for 2 h at 400, 500, 600, and 700 °C, respectively, with a heating rate of 5 °C/min to get the final products. TG-DTA experiments were conducted on a Netzsch STA449C thermal analysis system with pure air flow of 20 cm3/min. The sample was heated at a rate of 10°C/min. XRD patterns were collected using a PANalytical X’Pert PRO powder diffractometer with Cu Kα1 radiation (λ=0.15187 nm). The morphology of the sample was characterized by a JEOL-2010

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TEM equipped with the EDS detector. Samples for TEM studies were prepared by dispersing Ln3+:TiO2 powders onto a holey carbon support. Emission and excitation spectra and transient decays were recorded on an Edinburgh Instruments FLS920 spectrofluorimeter equipped with both continuous (450 W) and pulsed xenon lamps. For low temperature measurements, samples were mounted on a closed cycle cryostat (10-350 K, DE202, Advanced Research Systems). For site-selective spectroscopy, the excitation (or emission) monochromator’s slits were kept as small as possible to improve the instrumental resolution. The best wavelength resolution is 0.05 nm. The line intensities and positions of the measured spectra were calibrated according to the FLS920 correction curve and standard mercury lamp. Laser spectroscopic measurements were carried out upon excitation by a mode-locked ps Ti: sapphire laser (Spectra-Physics, Tsunami), which provides a pulse width of 1.5 ps, a repetition rate of 82 MHz and a tunable range from 700 to 1000 nm. The sample was mounted on an optical cryostat (Janis SHI-950, 4-300 K). The fluorescence was dispersed by a 1-meter monochromator (Jobin-Yvon 1000M) and detected with a cooled Hamamatsu R943-02 photomultiplier (PMT) or InGaAs detector. The signals were recorded by SpectrAcq2 data acquisition system with DM302 photon counting module (Jobin-Yvon).

3.2. Crystallization Behavior It has been revealed that the crystallization process has great effect on the incorporation of Ln3+ ions into TiO2 lattice sites [34]. So the systematic study of crystallization behavior of nanoparticles is of great importance to achieving high performance of nanophosphors. In the following, the Eu3+ doped TiO2 system is used as an example to study impact of dopant concentration on the crystallization and the thermodynamics of TiO2 nanoparticles.

3.2.1. Effect of Dopant Concentration on the Crystallization of Nanoparticles Figure 14 shows the TG-DTA curves of the as-prepared Eu3+:TiO2 (x=2, 4, 6 mol%) nanoparticles. The weight loss between 50–150 °C can be ascribed to the evaporation of absorbed water and ethanol in nanoparticles, and the weight loss between 150–400 °C accompanying with an exothermal peak at around 250 °C may be due to the decomposition of organic residues such as butyl alcohol byproduct. The exothermic peaks at approximately 400, 520, and 560 °C are related to the crystallization temperatures for samples with Eu3+ contents of 2, 4 and 6 mol% respectively, suggesting that the crystallization of TiO2 was hindered by the doping with Ln3+ ions. Figure 15 shows the XRD patterns of Eu3+:TiO2 annealed at 400 °C for 2 h with different dopant concentrations. As revealed in patterns, under annealing at 400 °C for 2 h, the samples with nominal Eu3+ concentration below 6 mol% exhibit pure anatase phase (JCPDS No. 711166, space group I41/amd), and the higher Eu3+ contents the worse crystallinity of the sample is observed. The sample of Eu3+:TiO2 (6 mol%) shows amorphous structure, which indicates a higher crystallization temperature required for its heat-treatment, in consistence with the results obtained by TG-DTA experiments.

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  3+

Eu 6 mol% Exo.

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DTA

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TG

Eu 4 mol%

80 100 3+

Eu 2 mol%

90 80

100

200

300

400

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o

Temperature ( C)

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Figure 14. TG-DTA curves of the as-prepared xEu3+:TiO2 (x=2, 4, 6 mol%) nanoparticles.

Table 1. Lattice parameters of Eu3+:TiO2 (x=0, 1, 2, 4 mol%) nanoparticles annealed at 400°C a (nm)

c (nm)

Volume (nm3)

Bulka TiO2:Eu3+ (0%) Eu3+:TiO2 (1%)

0.37845 0.37864 0.37893

0.95143 0.94186 0.94254

0.13627 0.13503 0.13534

Eu3+:TiO2 (2%) Eu3+:TiO2 (4%)

0.37909 0.38009

0.94527 0.95240

0.13584 0.13759

Samples

a

Calculated from JCPDS No. 71-1166. See also Ref. [36]

Because of the large ionic radius of Eu3+, the occupation of Eu3+ at Ti4+ site may lead to a slight lattice expansion and lattice strain. The crystal lattice parameters of Eu3+:TiO2 (x=0, 1, 2, 4 mol%) nanoparticles annealed at 400 °C are calculated from the (101) and (200) planes of XRD patterns, which are compared to the values of the bulk anatase in Table 1. Compared to the bulk anatase, an increase of a and a decrease of c are observed for the undoped anatase nanocrystals, where a and c denote the lattice parameters of the tetragonal unit cell. Similar variation was reported by Swamy et al., where a nonlinear size dependence of lattice parameters was observed for the neat anatase crystallites with a size less than ~15 nm [36]. For the samples doped with Eu3+ ions, the lattice parameters of a and c and the unit cell volume increase with the increase of Eu3+ content (Table 1). This variation can be explained

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by taking into account the fact that the higher Eu3+ contents the more Ti4+ ions could be replaced by Eu3+. The slight lattice expansion for samples with high Eu3+ content indicates that Eu3+ ions to some extent were incorporated into the TiO2 lattice.

3+

Intensity (a. u.)

Eu 6 mol%

3+

Eu 4 mol%

3+

Eu 2 mol%

3+

Eu 1 mol% 20

30

40

50

60

2θ (degree)

70

80

Figure 15. XRD patterns of xEu3+:TiO2 (x=1, 2, 4, 6 mol%) nanoparticles annealed at 400 °C for 2 h.

(224)

(116) (220) (215)

(204)

(105) (211)

o

700 C

(200)

(004)

Intensity (a. u.)

(101)

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o

600 C o

500 C o

400 C as-prepared

10

20

30

40 50 60 2θ (degree)

70

80

Figure 16. XRD patterns of Eu3+:TiO2 (2 mol%) nanoparticles annealed at different temperatures. (Reprinted with permission from [34], Luo, W. Q. et al., J. Phys. Chem. C 112 10370 (2008). © 2008, American Chemical Society).

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Figure 16 shows the XRD patterns of Eu3+:TiO2 (2 mol% ) nanoparticles annealed at different temperatures [34]. The as-prepared sample exhibits amorphous structure. Pure diffraction peaks of anatase TiO2 are obtained in the samples annealed at 400 to 700 °C. With the increase of annealing temperatures, the diffraction peaks become stronger and sharper, indicating an improvement of crystallinity of the samples. By means of Debye-Scherrer equation, the average sizes of samples annealed at 400, 500, 600, and 700 °C are estimated to be 12, 17, 19, and 27 nm, respectively.

Figure 17. TEM images of TiO2:Eu3+ (2 mol%) at different experimental stages: (a) samples obtained before solvothermal step; (b) samples obtained after solvothermal step (dried at 40 oC); (c) samples annealed at 400 oC for 2 h; (d) samples annealed at 600 oC for 2 h. The insets show the SAED patterns of corresponding samples.

To investigate the growth process of nanoparticles, samples taken from different stages of preparation process were subjected to the TEM characterization. Figure 17a shows TEM image of the precursor prior to hydrothermal treatment, the morphology of the obtained particles are irregular and the size is about 800 nm. The SAED pattern (inset of Figure 17a) shows amorphous nature of the precursor. Near monodisperse spherical amorphous TiO2 nanoparticles with the size of about 600–800 nm are obtained after solvothermal treatment at 120 °C for 5 h as shown in Figure 17b. After annealed at 400 °C for 2 h, clear polycrystalline rings indexed in anatase crystalline structure (inset of Figure 17c) is observed and the morphology and size of particles remain unchanged. As revealed by HRTEM observation, the

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spherical nanoparticles are the aggregates of 8–12 nm anatase TiO2 nanocrystals [34]. When the annealing temperature increases to 600 °C for 2 h, clearer polycrystalline rings are observed (inset of Figure 17d), suggesting a better crystallization of nanoparticles. The spherical morphology of nanocrystals remains unchanged and the size of nanoparticles slightly decreases to 500–650 nm as a result of sample densification under higher temperature heat treatment.

3.3. PL Properties of Ln3+:TiO2 Nanoparticles 3.3.1. Eu3+ Doped TiO2 Nanoparticles Recently, we have provided spectroscopic evidence of the multiple site structure of Eu3+ ions incorporated in the 8-12 nm TiO2 aggregates [34]. By means of site-selective spectroscopy at 10 K, as shown in Figure 18, three kinds of luminescence sites of Eu3+ are identified. One site (site I) exhibits broadened fluorescence lines with most intense emission at 613.3 nm similar to that of Eu3+ ion in glasslike phase, which is associated to the distorted lattice sites near the surface. The other two sites (Sites II and III) exhibit sharp emission and excitation lines with most intense emission lines at 616.7 and 618.1 nm respectively, which

Intensity (arb. units)

10K Emission

7

F0

7

F1

7

7

F2

F3

7

F4

(d)

(c) (b) ( X8)

(a) 600

650

Wavelength (nm)

700

Figure 18. (Color online) The 10 K emission spectra of Eu3+:TiO2 (2 mol%) annealed at 400 °C, with (a) λexc=343.0 nm, corresponding to the band-gap excitation; (b) λexc=464.6 nm for Site I; (c) λexc=470.7 nm for Site II; (d) λexc=472.1 nm for Site III. Color photographs of Eu3+:TiO2 and Eu3+:Y2O3 (2 mol%) are compared in (e) and (f), respectively. Both nanophosphors were excited with a xenon lamp at 465– 472 nm under the same experimental condition. To eliminate the influence of excitation light, a 495-nm long-pass glass filter was used when taking these photos. (Reprinted with permission from [34], Luo, W. Q. et al., J. Phys. Chem. C 112 10370 (2008). © 2008, American Chemical Society)

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T=10 K

τSite I =(0.37ms)

Intensity (arb. units)

1000

τSite II =(0.33ms) τSiteIII =(0.39ms)

100

10

Exp. Fit 0.0

0.5

1.0

1.5

2.0

2.5

3.0

time (ms) Figure 19. (Color online) The 10 K luminescence decays of 5D0 for Eu3+ at Sites I, II and III respectively. (Reprinted with permission from [34], Luo, W. Q. et al., J. Phys. Chem. C 112 10370 (2008). © 2008, American Chemical Society).

are ascribed to the lattice site with ordered crystalline environment. The value of full width at half maximum height (FWHM) of Sites II and III is much smaller than that of Site I, decreasing from ~9.0 nm (the 613-nm peak) to 0.58 nm (the 617-nm peak). As shown in Figure 18a, Eu3+ luminescence from sites II and III plus other minor sites can be observed when excited above TiO2 band-gap at 343 nm, indicating a weak ET from TiO2 host to Eu3+ ions. The weak ET can be understood because of the energy mismatch between TiO2 band gap and Eu3+ ion excited states. Since no energy level of Eu3+ is located in an energy range approximately from 28,500 to 31,000 cm-1 and thus this nonresonant host-to-Eu3+ ET can only be accomplished with the assistance of lattice phonons [37]. The PL intensity of Eu3+:TiO2 is comparable to that of Eu3+:Y2O3 (2 mol%) nanophosphors (prepared by sol-gel combustion method) under the xenon light excitation as shown in the inset of Figure 18. Note that the Eu3+:Y2O3 is a commercial red phosphor. The decay curves of 5D0 for Eu3+ at sites I, II and III are plotted in Figure 19. The three curves can be well fitted with single exponential and similar 5D0 lifetimes of 0.37, 0.33 and 0.39 ms for sites I, II and III are obtained, respectively. Eu3+ ion is a sensitive optical probe to detect local symmetry around Eu3+ ions. By carefully analyzing the high resolution emission and excitation spectra of Eu3+ ions at three sites, the local symmetries of three sites can be determined. The intensity ratio of 5D0→7F2 lines of Eu3+ with electric-dipole (ED) nature to that of 5D0→7F1 lines with magnetic-dipole (MD) nature can provide some structural information such as distortion of ligand environment and site symmetry of local environment around Eu3+ ions. As shown in Figure 18, the ED transitions of Eu3+ ions at three sites are all much stronger than the MD transitions, suggesting that Eu3+ ions occupy low-symmetry sites without an inversion center. Ti4+ ions sit at a D2d site in the anatase lattice. The substitution of Ti4+ with larger Eu3+ leads to a descent of the

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intrinsic D2d to a lower site symmetry (S4, C2v or D2), according to the branching rules of the 32 point groups [38]. Theoretically, if Eu3+ ions situate at D2d or S4 lattice sites, only two lines for J=0 to J=1 transition and three lines (S4) or two lines (D2d) for the J=0 to J=2 CF transition are theoretically allowed. However, three lines for 5D0→7F1 transition and four lines for the 5 D0→7F2 transition of Eu3+ ions at Site II can be clearly identified. Moreover, according to the ED selection rule, the 5D0→7F0 (0-0) transition is only allowed in the following 10 site symmetries, Cs, C1, Cn and Cnv (n=2, 3, 4, 6) [39-40]. The appearance of the 0-0 line suggests that Eu3+ ions at Site II may occupy a C2v symmetry. By contrast, for Eu3+ at Site III, the absence of 0-0 emission and three resolved lines from the emission of 5D0→7F1 plus three lines from the emission of 5D0→7F2 indicate a possible D2 site symmetry for Site III. Furthermore, according to the high resolution spectra, Eu3+ ions at Site II (or III) are very likely located at a distorted C2v (or D2) site, since one of the 5D0→7F2 lines was observed to be split into two neighboring peaks for Site II (or III). As for Eu3+ at site I, it resides in a disordered environment, therefore, should have the lowest site symmetry C1. Consistently, we observed all three lines for 5D0→7F1 of Eu3+ at site I.

Figure 20. Illustration of the site symmetry of (a) pure anatase nanocrystal, where Ti4+ occupies a D2d symmetry; (b) Eu3+ doped anatase nanocrystal, where Eu3+ occupies a C2v symmetry (Site II); (c) Eu3+ doped anatase nanocrystal, where Eu3+ occupies a D2 symmetry (Site III). The six nearest neighboring oxygens are labeled as O1 to O6. The top-view projection plane of these nearest neighboring atoms and the positions of symmetry operators (C2 and σv) are schematically plotted. Lattice expansion is vividly represented by slightly moving outward O1, O2, O3, O4, O5 and O6 in (b) and (c).

A possible model describing the lattice distortion from D2d site symmetry to C2v or D2 is depicted in Figure 20. Recently, a charge transfer vibronic exciton (CTVE) model was proposed to interpret the multi-site formation in Eu3+:BaFCl crystals [40]. Similarly, an

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oxygen vacancy or self-trapping CTVE may be formed to compensate for the charge imbalance when a trivalent Eu substitutes for the tetravalent Ti in the anatase, accompanied by a lattice relaxation. As shown in Figure 20b, the lattice expansion occurs evenly at various directions and an oxygen vacancy is created around Eu3+ ion at Site II, thus a C2v site symmetry is formed. For Site III, however, an oxygen vacancy is not physically formed; instead, the charge compensation may be collectively accomplished by CTVE. As shown in Figure 20c, two vertical symmetric planes (σv) at the original Ti4+ location were broken due to the uneven lattice expansion along specific directions when Ti4+ is replaced by Eu3+, thus a D2 site symmetry is formed. The thermal stability of Eu3+ in sites I, II and III was investigated by site-selective spectroscopy experiments at RT. Figure 21 shows the influence of annealing temperature on the emission intensities of 5D0→7F2 transition, which is highly sensitive to the local environment, of Eu3+ at three different sites. Multiple sites can be observable for the sample annealed at a temperature up to 700 °C. Whereas, the emission intensity of Eu3+ in sites I and III increased gradually with the annealing temperature increased up to 600 °C at the expense of the decrease of emission intensity of Eu3+ at site II, illustrating a site transition of site II to site I and III. When the annealing temperature increased above 700 °C, only emission lines from site I can be observed. The above facts indicate that Eu3+ ions at sites II and III in the samples annealed at high temperature are unstable, and may be expelled from the lattice to the grain or surface site.

Intensity (a. u.)

Site I Site II Site III

400

500

600

700 o

Annealing temperature( C) Figure 21. Influence of the annealing temperature on the emission intensities of 5D0→7F2 transitions of Eu3+ at different sites in Eu3+:TiO2 (2 mol%) nanocrystals.

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3.3.2. Sensitized Emissions of Sm3+, Nd3+ Doped TiO2 Nanoparticles Sensitization is an important process for the efficient PL of Ln3+ ions, which can overcome the inefficient direct excitation of parity forbidden intra-4f transition. Unlike Eu3+ ions doped samples, due to the match between the energies of the TiO2 band gap and Sm3+, Nd3+ excited states, the efficient sensitization emission of TiO2 host to Sm3+ or Nd3+ can be expected [41-42]. Figure 22 shows the RT excitation spectra of Sm3+ and Nd3+ doped TiO2 nanoparticles annealed at 500 °C for 2 h. Both samples exhibit broad bands at around 340 nm, which are associated with the band-gap of TiO2 nanoparticles, whereas only weak excitation lines originating from the f-f transitions of Sm3+ and Nd3+ ions are detected, suggesting more efficient PL of Sm3+ and Nd3+ ions by ET from TiO2 host to Ln3+ ions than direct excitation of Ln3+ ions. Upon the excitation above the TiO2 band gap at around 340 nm at RT, as shown in Figure 23a and c, intense sensitized emissions from Sm3+ and Nd3+ ions are observed. As shown in Figure 23b, upon excitation above TiO2 band gap at 10 K, sharp CF splitting emission lines related to the transition from 4G5/2 to 6H5/2, 6H7/2, 6H9/2 and 4H11/2 multiples of Sm3+ centered at 584.1, 612.8, 664.1 and 727.0 nm are obtained, indicating that Sm3+ ions were incorporated into a regular environment at TiO2 nanocrystal matrix. Due to Kramers degeneracy for f 5 configuration, 3, 4, 5 and 6 emission lines from the lowest CF level of 4 G5/2 to 6H5/2, 6H7/2, 6H9/2 and 6H11/2 are theoretically expected respectively for Sm3+ ions sit at a lattice site. As clearly identified in Figure 23b, more emission lines (i.e. at least 5, 6, 7 and 4 lines for the transitions from 4G5/2 to 6H5/2, 6H7/2, 6H9/2 and 6H11/2 respectively) than expected are detected, indicating that at least two different CF environments around Sm3+ ions exist in

400 (a)

RT

(a) TiO2:Nd (b) TiO2:Sm

Intensity (a. u.)

300 (b) 200

100

0 200

250

300

350 400 450 500 Wavelength (nm)

550

Figure 22. RT excitation spectra of (a) TiO2:Nd3+ (2 mol%) annealed at 500 oC for 2 h by monitoring 4 F3/2→4I11/2 transition at 1094 nm; (b) TiO2:Sm3+ (2 mol%) annealed at 500 oC for 2 h by monitoring 4 G5/2→6H7/2 transition at 612.8 nm. (Redrawn after [42], Luo, W. Q. et al., J. Phys. Chem. C 113 8772 (2009). © 2009, American Chemical Society).

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Intemsity (a.u.)

TiO2:Sm

6 6

4

G5/2→ H5/2

6

H11/2

(a)

RT

(b)

10K

500 550 3+ TiO2:Nd 4 4 F3/2→ I9/2

900

H7/2 6H 9/2

600

650

700

4

I11/2

4

I13/2

(c)

RT

(d)

10K

1000

1100 1200 1300 Wavelength (nm)

1400

Figure 23. Emission spectra of Sm3+ and Nd3+ ions doped TiO2 nanoparticles annealed at 500 °C for 2 h upon excitation above TiO2 band gap. (a) RT emission spectrum of TiO2:Sm3+ (2 mol%) under excitation at 343 nm; (b) 10 K emission spectrum of TiO2:Sm3+ (2 mol%) under excitation at 332 nm; (c) RT emission spectrum of TiO2:Nd3+ (2 mol%) under excitation at 338 nm; (d) 10 K emission spectrum of TiO2:Nd3+ (2 mol%) under excitation at 339 nm. (Redrawn after [42], Luo, W. Q. et al., J. Phys. Chem. C 113 8772 (2009). © 2009, American Chemical Society).

the samples. The multiple site emissions are also detected in Nd3+ doped TiO2 nanoparticles. As shown in Figure 23d, upon excitation above TiO2 band gap at 339 nm at 10K, at least 9, 7 and 4 emission lines centered at 915, 1094 and 1385 nm for the transitions from the lowest CF level of 4F3/2 to 4I9/2, 4I11/2 and 4I13/2 can be clearly identified respectively, more than that theoretically expected due to Kramers degeneracy for f3 configuration (i.e., 5, 6 and 7 emissions for 4F3/2 to 4I9/2, 4I11/2 and 4I13/2 respectively).

3.3.3. Er3+ Doped TiO2 Nanoparticles By modifying the synthesis condition, single site emissions of Er3+ in TiO2 with a nominal dopant concentration of 0.75 mol% were obtained [35]. Figure 24 shows the 10 K excitation spectrum of Er3+ doped TiO2 by monitoring the near-infrared (NIR) 4I13/2→4I15/2 emission at 1532.6 nm. Abundant sharp excitation lines centered at 380.6, 407.6, 489.4, 523.4, 550.5, and 654.0 nm can be assigned to the direct excitation from ground state of 4I15/2 to the upper excited states of 4G11/2, 2H9/2, 4F7/2, 2H11/2, 4S3/2 and 4F9/2, respectively. Fine CF splitting of the excited states of Er3+ can be easily identified, indicating that Er3+ ions are incorporated into a regular TiO2 nanocrystal lattice. It should be noted that a broad excitation band centered at 358 nm can be observed, which is associated with the band gap of anatase TiO2 nanocrystals, indicating the ET from TiO2 host to Er3+ ions. According to the Kramers

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degeneracy for the f 11 configuration, two excitation lines from the lowest CF level of 4I15/2 to 4 S3/2 are theoretically expected for Er3+ ions sit at a lattice site at low temperature. As clearly seen in the inset of Figure 24, there are only two lines assigned to the excited state of 4S3/2 (with an energy gap of 18 cm-1) and no trace of CF splitting due to another site can be observed. The CF splitting of other excited states in Figure 24 is also in good agreement with theoretical analysis, indicating that the doped Er3+ ions were very likely located at the same site in TiO2 nanocrystals. 2

T=10K

H11/2← I15/2

600

4

4

S3/2← I15/2

F ←I

7/2

2

556

560

4

F9/2← I15/2

4

4

300

552

S3/2← I15/2

4

← I15/2

0

548

4

F

15/2 4

3/2,5/2 4

200

544 4

H9/2← I15/2 4 G11/2← I15/2

Intensity (a.u.)

4

4

4

400

400 500 Wavelength (nm)

600

700

Figure 24. Excitation spectrum of Er3+:TiO2 nanocrystals at 10 K, and the inset enlarges the excitation lines for the transition of 4I15/2→4S3/2. (Reprinted with permission from [35], Fu, C. Y. et al. Opt. Lett. 33, 953 (2008). © 2008, Optical Society of America).

As shown in Figure 25a and b, similar NIR emission lines are obtained upon excitation either to the 2H11/2 state of Er3+ or above the TiO2 band gap, suggesting a homogeneous CF environment for the doped Er3+ ions. When excited above the band-gap energy, as shown in Figure 25c, a broad band (peaking at ~550 nm) attributed to defects is observed. These defects may be related to oxygen vacancies originating from the substituting Er3+ for Ti4+ [15]. In addition, sharp emission lines from 4S3/2→4I15/2 are observed being superimposed on the broad band. The eight CF levels of 4I15/2 are experimentally determined to be located at 0, 15, 95, 166, 210, 378, 454 and 504 cm-1 according to the 4S3/2→ 4I15/2 emission at 10 K. Interestingly, self absorption lines correspond to the hypersensitive transition of 4I15/2→2H11/2 are observed due to the large rank-2 reduced matrix elements (RMEs) of the unit tensor of the transition which could result in strong absorptions. Similar self absorption lines were also observed in Er3+ doped Gd2O3 nanocrystals [43]. Figure 25d shows the upconversion luminescence of Er3+ under a 976 nm laser excitation at RT. An intense green emission due to the transition from 2H11/2 and its thermally coupled 4S3/2 states to 4I15/2 is observed. As illustrated in the inset of Figure 25, the

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fluorescence decay curve of 4I13/2 is slightly deviated from a single exponential under the excitation of 523.4 nm, which may be caused by a nonradiative ET process from Er3+ to the neighboring defects that have close energy levels to the 4I13/2 state. Assuming the electric dipoledipole interaction between donor and acceptor, the decay curve can be well fitted by the Inokuti-Hirayama model [44]:

[

I (t ) = I 0 exp − t /τ 0 − C(t /τ 0 )

1/ 2

], where I is the time

dependent PL intensity and I0 is the initial intensity, C is a freely varied parameter, t is the time and τ0 is the intrinsic luminescence lifetime. The intrinsic fluorescence lifetime of 4I13/2 is fitted to be 1.56 ms at 10 K.

300

t=1.56m s t

4

200

4

100

I 13/2 → I 15/2

10

100

0

(a) (b)

0

1500

Intensity (a.u.)

1000

T=10K

300

1520

0 120 80

1560 4

T=10K

200 100

1540

2

2

3

4

t (m s)

1580

5

6

1600

S 3/2 → 4 I 15/2

H 11/2 ← 4 I 15/2

(c)

400

450

500

550

600

650

700

RT 2

H 11/2 , 4 S 3/2 → 4 I 15/2

40 0

1

4

(d)

400

450

500

550

F 9/2 → 4 I 15/2

600

650

700

W avelen g th (n m ) Figure 25. (Color online) NIR (a) (b), visible (c), and upconverted (d) luminescence spectra of Er3+:TiO2 nanocrystals, where (a)-(c) were measured under the 523.4, 358 and 358 nm excitation at 10 K and (d) was measured under the 976 nm laser excitation at RT, respectively. The inset shows the luminescence decay curve of (a), with the experimental (dotted) and fitted (solid) results. (Reprinted with permission from [35], Fu, C. Y. et al. Opt. Lett. 33, 953 (2008). © 2008, Optical Society of America).

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3.4. Growth Mechanism of Ln3+ Doped Nanoparticles For material synthesis, the study of growth mechanism of the materials is of great importance for the material design and performance. It has been demonstrated that, in the solgel solvothermal method, Ln3+ ions can be incorporated into the anatase TiO2 lattice and intense emission lines of Ln3+ ions can be obtained by either directly excited of Ln3+ ions or energy transfer from host matrix. However, Ln3+ ions in TiO2 lattice are metastable and will be expelled to the surface of TiO2 nanocrystals under the high temperature annealing (above 700 oC). Moreover, the solvothermal temperature was found to play a vital role in the PL performance of the samples. For example, the emission intensity from Eu3+ at TiO2 lattice gradually weakened with the increase of solvothermal temperature above 160 oC and eventually vanished at 170 oC, under which condition, TiO2 nanoparticles has been crystallized to anatase phase. No improvements in PL performance were achieved even after further annealing the nanocrystals at high temperature (above 400 oC), indicating that the crystallization process of amorphous TiO2 nano-aggregates plays a important role in the successful incorporation Ln3+ ions into TiO2 lattice. Based on the experiment results of TGDTA, TEM, XRD and PL, we have proposed a mechanism to illustrate the growth process of Eu3+ doped TiO2 nanoparticles, as shown in Figure 26 [34]. First, the hydrolysis of tetra(nbutyl)titanate followed by the condensation of the resulting titanium hydroxide species affords the -Ti-O-Ti- and -Ti-O-Eu- clusters which undergo the further condensation to form amorphous TiO2 nanoparticles with their surfaces covered with hydroxyl and Eu3+ ions which hindered the crystallization of the nanoparticles under post heat treatment. After solvothermal treatment, the amorphous TiO2 nanoparticles are coalesced to form amorphous submicron spherical aggregates (Figure 17b). By further heat-treatment (above 400 °C), two or more neighboring TiO2 nanoparticles are merged into one single anatase nanocrystals, and those Eu3+ ions originally embedded in the interfaces between nanoparticles may enter the Ti lattice

Figure 26. A schematic illustration showing the growth mechanism of Eu3+ doped TiO2 nanocrystals. (Reprinted with permission from [34], Luo, W. Q. et al., J. Phys. Chem. C 112 10370 (2008). © 2008, American Chemical Society).

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to lower the interfacial energy of nanoparticles during the crystallization process. However, compared to the surface site, the compatibility of Eu3+ ions at lattice sites II and III is relatively poor. If provided enough thermal energy, Eu3+ ions at the lattice site are prone to be expelled from the TiO2 matrix during the crystal growth.

4. Conclusion Ln3+ ion-doped anatase TiO2 nanoparticles are promising materials for phosphor devices. Various approaches to accommodate Ln3+ ions into TiO2 host materials have been reviewed. ET from TiO2 host to Ln3+ ions was observed in different TiO2 nanomaterials, illustrating that TiO2 is a good host and sensitizer for Ln3+ ions. However, in most cases, only unresolved PL lines were obtained due to the poor incorporation of Ln3+ into TiO2 lattice. A facile sol-gel solvothermal approach was proposed for the synthesis of Ln3+ doped TiO2 nanocrystals. Through the sol-gel process and self assembly of the nanoparticles, the Ln3+ ions can be doped into the nanoparticles. With the crystallization of the nanoparticles, the Ln3+ ions can be metastably incorporated into TiO2 lattices. The strategy in this method provides a possible way to design the syntheses of other semiconductor nanocrystals with Ln3+ ions incorporated into the lattice, which is usually difficult to do in other methods. The lattice incorporation of Ln3+ ions in TiO2 nanocrystals and efficient host-to-Ln ET obtained in these systems is of great significance for further material applications in the fields of optics and electronics. Moreover, the nearly monodisperse spherical morphology is ideal for luminescence devices because high packing densities and scattering of light can be obtained. Although some advances have been made in Ln3+ doped TiO2 nanomaterials, to gain deep insight into the understanding of chemistry and physics of Ln3+ ions in TiO2 nanocrystals, there is still much work to be done in further research in the future, such as the band gap engineering of TiO2 via size control, to achieve more efficient host-to-Ln PL, and systematic CF analysis of local environment and distortion around lanthanides.

Acknowledgements This work is supported by the One Hundred Talents Program of the Chinese Academy of Sciences (CAS), Knowledge Innovation Program of CAS for Key Topics (No. KJCX2-YW358) and Young Scientists, Instrument Developing Project of CAS (No. YZ200712), the NSFC (Nos. 10504032, 10774143 and 10804106), the 973 program (No. 2007CB936703), the National High-Tech R&D Program of China (863 Program, No. 2009AA03Z430), Fujian Provincial Science Fund for Distinguished Young Scholars (No. 2009J06030), and the Key Project of Science and Technology Foundation of Fujian Province (No. 2007I0024).

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[30] Palomino-Merino, R.; Conde-Gallardo, A.; Garcia-Rocha, M.; Hernandez-Calderon, I.; Castano, V.; Rodriguez, R. Thin Solid Films 2001, 401, 118-123. [31] Falcomer, D.; Daldosso, M.; Cannas, C.; Musinu, A.; Lasio, B.; Enzo, S.; Speghini, A.; Bettinelli, M. J Solid State Chem 2006, 179, 2452-2457. [32] Chi, B.; Victorio, E. S.; Jin, T. Nanotechnology 2006, 17, 2234-2241. [33] Zeng, Q. G.; Ding, Z. J.; Zhang, Z. M. J. Lumin. 2006, 118, 301-307. [34] Luo, W. Q.; Li, R. F.; Liu, G. K.; Antonio, M. R.; Chen, X. Y. J. Phys. Chem. C 2008, 112, 10370-11377. [35] Fu, C. Y.; Liao, J. S.; Luo, W. Q.; Li, R. F.; Chen, X. Y. Opt. Lett. 2008, 33, 953-955. [36] Swamy, V.; Menzies, D.; Muddle, B. C.; Kuznetsov, A.; Dubrovinsky, L. S.; Dai, Q.; Dmitriev, V. Appl. Phys. Lett. 2006, 88, 243103. [37] Carnall, W. T.; Goodman, G. L.; Rajnak, K.; Rana, R. S. J. Chem. Phys. 1989, 90, 34433457. [38] Butler, P. H. Point Group Symmetry Application: Method and Tables; Plenum: New York, 1981. [39] Chen, X. Y.; Liu, G. K. J. Solid State Chem. 2005, 178, 419-428. [40] Chen, X. Y.; Zhao, W.; Cook, R. E.; Liu, G. K. Phys. Rev. B 2004, 70, 205122. [41] Luo, W. Q.; Li, R. F.; Liu, Y. S.; Chen, X. Y. J. Nanosci. Nanotechno. 2009, in press. [42] Luo, W. Q.; Li, R. F.; Chen, X. Y. J. Phys. Chem. C 2009, 113, 8772-8777. [43] Chen, X. Y.; Ma, E.; Liu, G. K. J. Phys. Chem. C 2007, 111, 10404-10411. [44] Inokuti, M.; Hirayama, F. J. Chem. Phys. 1965, 43, 1978-1989.

In: Nanotechnology… Editors: C.J. Dixon and O.W. Curtines, pp. 509-524

ISBN: 978-1-60692-162-3 © 2010 Nova Science Publishers, Inc.

Chapter 16

NANOCRYSTALLITE SUPERHARD TITANIUM NITRIDE FILM IN MULTI-ARC ION PLATING Xiang Yu1, Chengbiao Wang1, Meng Hua2, Yang Liu3 and Shengli Ma4 1

School of Engineering and Technology, China University of Geosciences, Beijing100083, China 2 MEEM Department, City University of Hong Kong, 83 Tat Chee Ave., Kowloon Tong, Kowloon, Hong Kong, China 3 Beijing Powertech Co. Ltd., Beijing 100072, China 4 State-Key Laboratory for MBM, Xian’s Jiaotong University, Xian 710049, China

Abstract Titanium nitride (TiN) films synthesized by multi-arc ion plating (AIP) normally have a columnar microstructure, and are likely to induce surface defects due to the formation of macroparticles and neutral particles in the vicinity of cathode arc sources. Hence, the achievable microhardness of the normal AIP TiN films only ranges between 20~30 GPa. A systematic study for fabricating an adherent nano-superhard titanium nitride (TiN) film on M2 high speed steel substrate by a vacuum cathode multi-arc ion-plating (AIP) system was initiated. To understand the relationship of the film processing-structure-property, their microhardness, film-to-substrate adhesion, frictional property, and microstructure of the film were investigated using Vickers hardometer, scratch tester, ball-on-disc tester, X-ray diffractometer, and transmission electron microscope. Results show that: (i) the achievable film microhardness ranges between 35 GPa and 45 GPa; (ii) the critical load (Lc) of the superhard TiN film is at 64 N approximately; (iii) the friction coefficient, under a high-load and a high rotating-speed, of the film is ranging from 0.5 to 0.8; and (iv) the nm scale mean main grain-sizes of the film are approximately 12.7 nm for TiN111, 19.7 nm for TiN200 and 9.6 nm for TiN220. The maximum achievable microhardness 45 GPa is more than twice of the 22 GPa for standard TiN film. Such hardness enhancement is anticipated as mainly due to: (a) the formation of nanoscaled crystalline grains; (b) the preferential orientation and growth of

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Keywords: Superhard TiN film; mechanical property; crystalline size; preferential orientation

1. Introduction TiN film has been extensively studied and successfully used for decades due to its favorable properties of high melting point, high hardness and high thermal conductivity [1]. Its application ranges from hard and protective coating on mechanical tools to decorative coating, and diffusion barriers in microelectronic components [2,3]. Currently, the rapid development of engineering and high technologies brings strict demand on the quality of TiN film, and therefore the performances of the film still need to be improved further. Multi-arc ion plating (AIP) with features of high ionization rate, ion kinetic energy and deposition rate has been rapidly developed in the last two decades. During depositing AIP TiN film, macroparticles and neutral particles are usually accompanied with the ions in the vicinity of cathode arc sources. This subsequently roughens relatively the TiN film structure with the presence of columnar grains and other defaults. Hence, this type of films has low erosion and corrosion resistance, and poor brightness. Their microhardness is normally in the order of 20~30 GPa [4,5]. Conventionally, the approach to improve the hardness, wear and corrosion resistance performances is to increase the thickness of the film. However, this may not be usually effective. Improvement of the film intrinsic and/or extrinsic properties like grain size, morphology, texture, etc. can also be the alternative approach to enhance these performances [6]. Resent interest in nanotechnology stimulates numerous studies in films, specifically aiming at understanding the atomic level of their formation and growth mechanisms. The growing demand for superior protective coatings to operate in severe conditions results in the novel development of hard and superhard coatings like multilayered coatings, and superlattices and/or heterostructures, nanostructured composite coatings [7,8]. This study was initiated to synthesize an adherent superhard TiN film on the substrate of M2 high-speed steel using multi-arc ion-plating coater. It also investigated the relationship among the deposition parameters, the mechanical properties and the nanostructure so as to understand the forming mechanism of the film. Recent research in preparation of industrial materials has led to coatings with hardness comparable to diamond and cubic boron nitride. This has subsequently aroused a great deal of interest in synthesizing nanostructured superhard TiN films with Vickers hardness ≥40 GPa. The nanostructures of these superhard materials are identified to be either ternary or quaternary. However, techniques to synthesize these films are generally rather complex because the films usually: (i) constitute of alternating layers of nm-scale transitional metal nitrides [9] and (ii) are embedding the transitional metal nitride nanocrystals within the amorphous phase of a covalent nitride (e.g. nc-TiN/a-Si3N4, nc-(Ti1-xAlx)N/a-Si3N4, nc-TiN /TiB2, nc-TiN/BN, etc.) with the element-doped nanocomposites [10]. As binary TiN films become outmoded and replaced by superhard films, it is difficult to achieve microhardness as beyond 40 GPa due to inadequate control of the film processing parameters. Our recent deposition method has allowed superhard nanocrystallite TiN films to be accomplished with

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microhardness values consistently up to 45 GPa, with none of the complexity (i) and (ii) as mentioned above.

Figure 1. Schematic cross-sectional view of SP0809AS coater.

2. Experimental Details 2.1. Coating Equipment and Sample Preparation Table 1 tabulates the operating conditions of the SP0809AS type vacuum coater (Figure1) used in fabricating the films. The surface of steel substrate was firstly polished to an average surface finish of Ra=0.08 µm, and was thoroughly cleaned in an acetone ultrasonic bath for 15 min and blow-dried with nitrogen. The so-prepared steel coupons were hung vertically on the steel racks of a rotational substrate holder in the centre area of the vacuum coater. Eight Ti rod targets were fixed around the vacuum chamber. Ar gas was injected to each of the two vertically collocated arc evaporation sources from the 1 mm diameter holes uniformly arranged on a double-layer steel bar, and nitrogen gas was injected towards the substrate from the two ion sources for improving ionization efficiency. The flow of Ar and N2 gases was controlled by two mass flowmeters. The bias voltage and the bias current to the substrate were controlled by a unipolar pulse DC bias power supply. The duty ratio of the pulsed bias voltage in Ar ion cleaning was 70% and in deposition was 50 %. The temperature of vacuum chamber before deposition was raised up gradually to 300 oC by two 4 KW group heaters, followed by further cleaning the coupon surfaces for 15 min by bombarding Ar ion at a bias voltage of –1200 V in a 4 Pa argon atmosphere. By maintaining the depositing pressure at 2×10-1 Pa in the argon atmosphere, Ti ion etching was performed for 5 min under the – 1000 V bias voltage, which was followed by 10 mins pre-deposition of a 0.2 µm Ti interlayer

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using the eight Ti rod targets with 80 A arc current and under –300 V. The deposition of TiN layer was initiated for accomplishing a film thickness of approximately 3 µm by introducing nitrogen gas with the target voltage stabilized at 20 V. A PLC (Programmable Logical Controller) system was used to control automatically the process parameters in each coating. TiN layers were deposited with individual bias voltages varying from –100 V to -600 V in step of –100 V under the 2×10-1 Pa depositing pressure. Table 1. Setting of operational conditions of deposition equipment Deposition method

Vacuum cathode multi-arc ion-plating

Arc target Substrate Working gases Substrate bias power Base pressure Arc current /Target voltage Heating temperature Ar ion bombardment Ti ion etching Depositing Ti interlayer pressure TiN layer

Ti×8, purity of 99.99% Cr12Mo4V high-speed steel, 20×20×5 mm3 Ar + N2, purity of 99.999% 0–1500V (unipolar DC pulse), 5–95% (duty ratio), 0–20 A (bias current) 1.0×10-3 Pa 80 A/20 V 300 oC Bias voltage -1200 V and duty ratio 70%, 4 Pa, 15 min Bias voltage -1000 V, duty ratio 50%, 5 min, 0.2 Pa 0.2 Pa (Ar 80 Sccm) 0.2-0.6 Pa (Ar/N2, Ar 80 Sccm, N2 0-165 Sccm)

2.2. Methods of Surface Analysis A MH-6 microhardness tester with a Vickers indenter under 20 g loading was used to measure the hardness of the films. A MS-S3000 scratch tester with initial load 3 N, loading rate 100 N/min and transverse speed 4 mm/min was used to investigate the adhesion between the steel coupons and the TiN films in air atmosphere. Optical microscope with a video camera was used to observe the wear traces in scratch-testing. A DD-92 ball-on-disc tester was used to investigate the frictional property of the films. Toughness was evaluated by ballpressing test under a vertical force of 100 kg with a Φ12 SiC ball pressing and dwelling onto a TiN filmed disc coupon for 5 mins. A JEM3010 high-resolution transmission electron microscope (HRTEM) and a X’pert Pro X-ray diffractometer (XRD) were used to investigate the crystallographic structures of the TiN film. A CuKα radiation operating at 40 kV and 40 mA was used to record the scanning results in the 2θ range from 20 o to 75o (by XRD detection) with a step size of 0.02o and a dwell time of 4 s. All results were taken with a sampling depth of 200 nm and the sampling area of about 10×10 mm2. Subsequently, the level of preferred orientation (POD) of the TiN films was determined by the following formulae [11].

[POD]

0o k

⎛ I k0 =⎜ s ⎜ Ik ⎝

o

⎞ ⎟ ⎟ ⎠

1 ⎛⎜ I k0 ∑k n ⎜ I s ⎝ k n

o

⎞ ⎟ ⎟ ⎠

(1)

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Where: Ik0 is the relative intensity measured at an orientation angle of α=0o; Iks is the relative intensity in the same plane of the standard PDF (which is the abbreviation of Joint Committee for Powder Diffraction Standards-International Centre for Diffraction Data (JCPDS-ICDD)) card; and n is the total number of reflection peaks obtainable from the film.

Figure 2. Bias voltage vs. microhardness of TiN films.

3. Results and Discussions 3.1. Mechanical and Tribological Properties 3.1.1. Microhardness To measure accurately the Vickers microhardness of TiN film thickness in micron-scale, the mathematical mean of five measurements, under the same measuring condition, at four different points was taken as the experimental value. Figure 2 shows the relationship of the bias voltages and the micorhardness of the TiN films. The micrographs in Figure 2 illustrate: (1) the specimen of the intended superhard TiN film of 45 GPa and (2) the pressing-ball used to perform the indentation. The achievable microhardness values, indented under 20 g loading, of the TiN film coupons deposited with a bias voltage in the range from -100 to -600 V were between 35 GPa and 45 GPa. The largest limit of 45 GPa was above twofold of the films deposited by the arc-ion-plated TiN films [10]. The morphology of the film surface was clearly seen on the indentation micrograph (1). The micrograph (2) for the ball-pressing test

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suggested that the toughness of the film was high and beneficial in accomplishing good adhesion. The high ion current used in the deposition of the TiN films resulted in high substrate bias current. Typical substrate bias current was in range of 6 A ~14 A, which was almost fourfold that in the magnetron sputtering [11]. The deposition pressure in scale of 10-1 Pa was low, and would provide purer space and longer mean free path for the film-to form ions than that in several-Pascal scale as in the ordinary arc ion-plating process. This activated high-energy ion bombardment in scale of several-hundred Volts for forming fine and dense structure to enhance the microhardness of the films. Biasing the substrate increases the density of excited radicals and brings along high-energy suitably for ion bombardment onto the substrate. This also strengthens the microhardness of the films. However, the magnification of residual stress in the TiN films [12] when the bias voltage is too large, typically beyond -350 V, may degrade the film mechanical properties. Hence, the achievement of superhard TiN films requires the adequate optimization of the process parameters like the substrate bias voltage and the gas pressure. Following analyses were performed on a 45 GPa TiN film deposited with a -200 V substrate bias voltage and a 0.2 Pa gas pressure.

3.1.2. Film-to-Substrate Adhesion The surface of the specifically deposited superhard TiN film was scratch-traced across by a diamond stylus under the condition of gradually increasing the normal loading. During the loading process, the noise signals and images of the friction force produced in the scratch test were monitored and plotted (Figure 3). Subsequently, the critical load (Lc) was determined by studying the variation of the curves for the friction forces and noise signals along with the images of typical morphologies, and the scratch history was calculated from the images of typical scratch traces (Figure3). The flat trend with very mild variation of the scratch trace in the image 1 of Figure 3 resulted in the relatively low friction force and friction noise. A few

Figure 3. Curves of changes of friction forces and acoustic signals during scratching test and typical morphology images of superhard TiN film.

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wear cracks in those visible scratch traces were observed on the morphology of the image 2 (Figure 3) when the loading was increased. These cracks were not large enough to propagate through the thickness of the TiN film layer and they did not break the film. Under such circumstance, the friction force and noise were gradually increased. Increasing further the loading would lead to the commencement of forming chip in the film, as seen on the image 3 in Figure 3. When the applied load was approaching the critical breaking value Lc＝ 64 N of the film, some visible traces of peeling were detected on the film surface and sudden raise in the friction force and noise was observed. When the applied load was increased beyond this critical value, enlargement of cracks and chips along the scratch trace (see Image 4 in Figure 3) took place. Such scratching phenomena led the curves of friction forces and noise to fluctuate sharply. The film was peeled off at the end stage when a large load of 100 N was applied.

3.1.3. Tribological Property In order to investigate the frictional property of the superhard TiN film working under high load and high-speed conditions, ball-on-disc tests were performed loaded under 10g and 20g, respectively, were performed. Such loading conditions were generally considered as relatively high. The Al2O3 ceramic ball used in the tests was φ8 cm and the rotating test-speed was 2000 revolutions/min. The two curves of friction coefficients of the superhard TiN film under the two individual loads were shown in Figure 4. The image inserted in the figure illustrates the morphology of the disc wear track under the corresponding loading conditions.

Figure 4. Change curves of friction coefficients of superhard TiN film under different loads.

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The frictional coefficients of superhard TiN film were found to be between 0.5 and 0.8. The high coefficients of friction would be due to the formation of macroparticles during AIP deposition. The average friction coefficient under the 10 g loading was 0.56 and that under the 20 g was 0.67. This confirms that larger load is susceptible to frictional deterioration. The inserted image shows sight of some small cracks in the wear areas. Such cracks may be brought along by the drop-out and repetitively frictional rubbing of the ball against those tiny asperity peaks on the film surface. Tests under the high-speed and high-load conditions indicated that the superhard TiN film: (i) did not fail in the one-hour testing; and (ii) did give a normal range of friction coefficients. Such noteworthy results imply that using the suprhard TiN film can improve the wear-resistant performance of the steel substrate.

3.2. Microstructure The microstructure analyses of the superhard TiN film for comparing with those of standard TiN film in PDF card 6-642 were performed by XRD technique. The XRD spectrum (Figure 5) of the superhard TiN film displayed a strong preferred crystalline orientation in the (111) orientation plane and weak ones in the orientation planes of (200) and (220). The relative intensity I/Io values of the standard TiN film in PDF card 6-642 showed: (I) the maximum characteristic peak of 100 for the (200) plane at 2θ ≈ 42.4o; and (ii) the decreasing of individual intensities for the other two characteristic peaks from 75 for the (111) plane to 55 for the (220) plane. The XRD analyses showed that the relative intensity of the maximum characteristic peak for the superhard TiN film was 100 for the (111) plane at 2θ ≈ 36.5o and its minimum characteristic peak was 2.5 for the (200) plane. Obviously, the comparison

Figure 5. XRD spectrum of superhard TiN film.

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illustrates that for the superhard TiN films deposited using AIP system: (I) has given a preferred orientation of (111) plane; (ii) has moved its maximum characteristic peak to 2θ ≈ 73.7o in the (111) plane instead of at 2θ ≈ 42.4o in the (200) plane for the standard TiN film; and (iii) has resulted in more striking intensity difference between the maximum and minimum characteristic peaks. As the structure of TiN material belongs to a NaCl face-centred cubic type in which Ti atoms constitute mainly to the skeleton of the crystal lattice with N atoms filling into its interspaces, the TiN films hence share the densest structure at (111) plane rather than in the planes of (200) and (220) [13].The formation of the preferred orientation and preferred growth at (111) plane for the superhard TiN film may thus be due to concentrating the crystalline grains at the plane (111) that subsequently constituting orderly and uniform arrangement. However, their arrangement at the planes of (200) and (220) is less uniform. Such orderly and preferred orientation greatly contributes to the super-hardness. When the crystal sizes of the deposited TiN films are larger than 10-4 cm, the depths of their diffraction peaks become insensitive to the change of crystal sizes. But, the broadened diffraction peaks of the films are sensitive with the decrescent crystal sizes when their sizes are below 10-5 cm. Consequently, the crystal sizes can be calculated using Scherrer formula [15]:

Dhkl = Kλ

B1 / 2

cosθ

(2)

Where: Dhkl is the crystal size along a direction normal to the plane; K is the Scherrer constant; λ= 0.154nm and is the X-ray wavelength of CuKα radiation; B1/2 is the full-width at half maximum of a Bragg peak; and θ is Bragg angle. Calculation gave the values of the mean nm-scale grain-sizes of the film as 12.7 nm for TiN111, 19.7 nm for TiN200, and 9.6 nm for TiN220. All these values are smaller than that of the other AIP TiN films as reported in [16]. The layer number of TiN grain lapped normal to plane (111) was determined by N111= D111 /d111 = 12.7/ 0.244 =52 (layers). The calculation gave 52 layers of crystallographic plane normal to plane (111) to constitute a TiN111 grain for the deposited TiN film. Their mutual interaction subsequently re-enforced the strength of the so deposited film.

Figure 6. Superhard TiN film’s (a) conventional TEM micrograph and (b) HRTEM micrograph.

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For illuminating further the film microstructure, TEM and HRTEM analyses (Figure 6) on the nanocrystalline superhard TiN film were conducted. The conventional TEM micrograph (Figure 6(a)) suggested that the nanometer scale crystal grains were finely and uniformly distributed. The HRTEM micrograph (Figure 6(b)) revealed clearly the distribution of the preferred orientation in most area on the plane (111) of the nanocrystalline. The above TEM analysis and the XRD analysis mutually and supportively confirm the existence of a microstructure characteristic of the preferred crystalline orientation in nanometer scale on the plane (111) of the superhard TiN film. Such nonopolycrystalline structure and preferred orientation at the adjacency of crystal growth create barriers to restrain the dislocation sliding and plastic deformation of crystal grains. As a result, it induces the superhardness in the film.

3.3. Hardness Enhancement Mechanism All the mean micorhardness values, as obtained by four data points method, for the coated TiN samples were in the range of 34 ∼ 45 GPa. Since some of these hardness values are above 40 GPa that is almost twice the reported value for tranitional arc-ion-plated TiN films, they are thus considered as super-hard coatings. As a result, it also confirms that the AIP techniques we used are able to produce super-hard coating under the properly controlled conditions. High ion current and ion energy values were generated in the filming process. The bombardment of these high-energy ions under several-hundred volts facilitates the deeper penetration and better fusion so that fine and dense structure can be formed. Our testing results also showed that the film microhardness varied notably within certain ranges of the depositing pressure and bias voltage. Such finding suggests that the depositing pressure and the bias voltage are the two major parameters influencing the formation of the film microstructure that, in turn, determines the microhardness of the film. Table 2. Comparison of XRD data of nano-superhard film with standard TiN film k (hkl)

6-642

Superhard TiN film

I/Io

Ιk

θ k(°)

d(nm)

D111(nm)

N(layer)

111

75

100

18.25

0.244

12.7

52

200

100

2.5

21.20

0.212

19.7

93

9.6 220 55 3.5 30.76 0.1496 Keynotes: I is for intensity, θ is the angle and subscript k stands for superhard film.

64

Table 2 compares the XRD data of the superhard TiN film with the standard TiN film in PDF card 6-642. Generally, the magnitude of the diffraction peaks is insensitive to the changes in crystal size when the crystal size is beyond 10-4 cm. But the broadening of the diffraction peaks is sensitive to the sizes of crystallites when their sizes are below 10-5 cm. Calculation based on the measured B1/2 gave the mean grain-sizes of the film as 12.7 nm for TiN111, 19.7 nm for TiN200, and 9.6 nm for TiN220. However, the peak of XRD spectrum for the (200) plane is much broader than that for its (111) plane as shown in Figure 7 for the XRD spectra obtained from a nanocrystalline superhard TiN film. This seems to contradict with the experimental result that has shown the grain-size for (200) plane to be larger than the grain-size for (111) plane. The peak broadening of a spectrum may have been resulted from

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the effect of either one or the combination of the three factors: (i) instrumental broadening, (ii) smaller-grain and (iii) microstrain. Identification of the respective contribution of these three factors to the broadening generally requires a series of carefully planned and tedious experiments. Analysis of the results (Figures 7 and 8) indicates that: (a) the peak intensity of TiN200 and TiN222 (Figure 7) is much weaker than TiN111 – reflecting a larger attribution from the grain size to the peak broadening; (b) the influence of Kα2 on Kα is negligible at low 2θ and becomes obvious at 2θ > 45o – implying the instrumental broadening is more significant for the (200) plane spectrum rather than for the (111) plane spectrum (Figure 7); (c) there is not any obvious plane distortion and bending respectively for the (111) and (220) planes in the HR-TEM micrograph (Figure 8(a)) – implying the negligible attribution from microstrain broadening. The above analyses therefore suggest that the peak broadening is mostly attributed to the smaller-grain size. They also affirm that the calculation by Eq.(2) really gives a reliable guide to the trend of the grain-sizes. The layer number N lapped normal to a crystalline plane can also be calculated theoretically. The calculated N lapped normal to the (111) plane of the TiN grain by N111=D111/d111=12.7/0.244 =52 (layer) implies the structure of a TiN111 grain in a crystallographic plane normal to the (111) plane to be constituted by 52 layers. Such structure gives the highest density for the atomic arrangement and induces hindering to the dislocation within the TiN nano-crystallines so as to enhance the film macroscopic mechanical behaviors.

Figure 7. XRD spectra of nanocrystalline superhard TiN film.

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The XRD spectra (Figure 7), HR-TEM micrograph (Figure 8(a)) and electron diffraction pattern (Figure 8(b)) confirmed the three main crystallographic planes (111), (200) and (220) of the TiN films. Our study (Table 2) showed that: the obtained value of the relative intensity I/Io of the maximum characteristic peak for a standard TiN film was 100 for the (200) plane at 2θ ≈ 42.4o, 75 for (111) plane, and 55 for (220) plane; whereas the obtained intensity value Ik for a superhard TiN film was 100 for the (111) plane at 2θ ≈ 36.5o, and the minimum characteristic peak of 2.5 for the (200) plane. Hence, it is derived that the superhard TiN films deposited by the ion-plating system have a preferred orientation in (111) plane. Moreover, the maximum characteristic peak is shifted to the plane (111) at 2θ ≈ 36.5o for the superhard TiN film from the (200) plane at 2θ ≈ 42.4o for the standard TiN film in PDF card 6-642. Comparing the relative intensity I/Io values in the column 6-642 for the standard TiN film and in the column Ik for the superhard TiN film (Table 2), it shows a striking difference of the intensity between their maximum and minimum characteristic peaks. Typically, their respective planes have mutually swapped over. The preferential crystalline orientation of the nano-crystalline TiN film is strong on the plane (111) and weak on the planes (200) and (220). With the aim of alleviating the effect of steel substrate on the diffraction peaks, STD diffraction at αo＝0.6º, 3º and 5º were respectively conducted and their XRD spectra were correspondingly labeled as (2), (3) and (4) in Figure 7. The location of lower azimuth angle and the shallower penetration depth in Figure 7 suggests that the elimination of the substrate effect was advantageous for the analyses. It can be observed from Figure 7 that: (i) there is an obvious peak for the substrate in the case (1); (ii) a reduction in substrate peak intensity, in cases (2) to (4), can be clearly seen; (iii) the lowest X-ray depth observed in the case (4), which suggests its substrate effect is minimum. The difference in intensities for crystallographic planes (200), (220) and (311) of the superhard TiN film as seen from the STD diffractions confirms that preferred growth has taken place at the close-packed plane (111) for the superhard TiN film. The TiN material belongs to NaCl type face-centred cubic structure that has a skeleton of Ti atoms as the crystal lattice and N atoms in its interstices. The densest film structure is in the (111) plane rather in the planes (200) and (220). The preferred orientation and preferred growth being in the close-packed plane (111) of the superhard TiN film may mainly be due to the concentration of crystalline grains at the (111) plane and its associated uniform arrangement. Such a high degree of order plane subsequently provides the densest and hardest structure when compared with the planes (200) and (220). Figure 8 shows the HRTEM micrograph and the electron diffraction pattern of a selected area of the multi-grains on a deposited TiN film. It illustrates the existence of preferred crystalline orientation on plane (111) and the occurrence of local orientation at the grain interface on (220) (Figure 8a). The electron diffraction pattern of the correspondingly selected area (Figure 8b) confirms the existence of an obvious nano-polycrystalline character for the deposited TiN film. As distinguished from the ring-shaped diffraction rings obtained from amorphous grain planes, Figure 8b shows a distinct arc-shape cluster of diffraction rings for the crystalline planes (111), (200) and (220) with the brightest ring being for the (111) plane.

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( c)

1μm

Figure 8. (a) HR-TEM micrograph, (b) selected area electron diffraction pattern and (c) cross-sectional micrograph of the TiN film.

It is thus reasonable to derive that these planes, particularly the close-packed (111) plane, are those for the TiN nanocrystalline grains to grow. Figure.8c shows a cross-sectional micrograph of a TiN film. It can be seen that the film has a dense and fine microstructure, and there is no sight of a columnar character. The HRTEM micrograph shows a fine, uniform and nanometer-scale crystalline size for the superhard TiN film, which is consistently in agreement with the traces of XRD analysis in Figure 7. During the formation of the film, the high flux and energy of the energetic ion bombardment enhances the mobility and dispersive capacity of the adsorptive particles and consequently facilitates the arrangement of the orderly atoms. Calculation by Eq.(1) with the relevant measurements gave the degree of preferred orientation (POD) on the plane (111) of the deposited superhard TiN film as 3.59. This calculated POD value is much higher than that for the reported sputtered TiN film [17] and serves to verify the methods used for the accomplishment of a preferred orientation of the TiN film.

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Figure 9. Correlation among film hardness, PODs and gas pressures & bias voltages.

To reveal the natural characteristics of the nanocrystalline TiN film, a correlation of the experimental hardness and structure to their process parameters was performed. Experimentally, it was found that by varying the process parameters of substrate bias voltage and gas pressure a superhard TiN film was produced which exhibited a microstructure with (i) a preferred orientation and preferred growth in the close-packed plane (111) (that can be expressed in terms of POD) and (ii) gave mean grain-sizes on a nm scale. A typical correlation among the TiN film microhardness, the PODs of TiN111, the gas pressure as well as the bias voltages is shown in Figure 9, on which an indent image illustrates the geometry and dimensions of an indented superhard TiN film having hardness of 45 GPa. The depth of indentation was made specifically in the range of 0.140 ∼ 0.34 μm so as to minimize the influence of substrate on the measured film hardness. The deviation of the microhardness measurements so obtained was estimated and found in an order of 10 %. The result in Figure 9 shows that: (i) the superhard TiN films have the structure character of POD beyond 3.1; and (ii) suitable selection and simultaneous adjustment of gas pressure (in the range of 0.10 Pa ~ 0.30 Pa) as well as pulsed bias voltages (in the range of -100 ~ -250 V) allowed the formation of superhard TiN films with the required structure. The AIP deposition method is likely to activate high ionization in the vacuum chamber. Hence, biasing the substrate leads to the generation of a glow discharge surrounding the ion sources and the substrate holder. The effect is to enhance further the ionization and reactivity of nitrogen gas. The majority of the nitrogen under such circumstance reacts with titanium atoms in substrate surface, and with

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discharging plasma and target surface. Although the application of biasing increases the density of excited radicals on the substrate and suitably activates high-energy ion bombardment to the substrate, it may however increase the residual stress in the TiN films and result in the loss of the required special structural properties when the bias voltage is above -350 V. Consequently, it degrades the mechanical properties of the film. In our experiment, the AIP deposition was set to a steadily low gas pressure of 0.1 Pa. It thus gave a relatively pure deposition environment and it also increased the mean free paths of the favorably excited particles. Optimization of the process parameters like the substrate bias voltage and the gas pressure thus facilitates the formation of superhard TiN films. The results of this study therefore indicate that the PODs of TiN111 are a good indication for judging whether the microhardness of the TiN films will increase or decrease. The phenomenon of hardness enhancement of the nanocrystalline TiN film is closely related with the ion flux and energy associated with the bombardment by the energetic particles in the film growing process. Estimation by Veprek et al [18] indicated that such a type of induced film superhardness would degrade the TiN film intrinsic hardness to, or below 22 GPa when its residual stress was reduced after the deposition period by an annealing treatment. Data in our experiments suggested that there was no visible decrease in the hardness values of the TiN films after one-month deposit under room temperature. Our experiments showed that a 45 GPa TiN film reduced its hardness to 36 GPa (i.e. 21 percent decrease) when it was annealed in a furnace at 600 oC for 0.5 h. It was only 21 percent decrease and the hardness value after annealing was still much higher than the intrinsic hardness of a standard TiN counterpart. The measured value of the residual stress of the film before and after annealing was 2.43 GPa and 1.54 GPa, respectively. Such a reduction suggests that annealing may result in grain growth and an associated decrease in hardness of the film. Consequently, it leads us to conclude that the residual stress induced in the film by the energetic ion bombardment can also make a partial contribution to the film hardness enhancement.

4. Conclusion An adherent nanocrystalline superhard TiN film was deposited on the substrate of M2 high-speed steel using a multi-arc ion plating system that was operated under (i) a low depositing pressure of 0.2 Pa and (ii) the energetic bombardment of the high ion flux and ion energy. The nanocrystalline superhard TiN film so deposited possesses a microstructure characteristic of nanometer scale preferred crystalline orientation on the plane (111), which is subsequently inducing superhardness of the film, as confirmed by the analyses of both XRD and TEM. Our studies have illustrated that it is possible to produce superhard naonosrystallite TiN films with microhardness of 45 GPa. The studies have also indicated that it is possible to correlate the preferred orientation on the close-packed plane (111) and the growth of TiN layers with the film hardness. Furthermore, the results of TEM studies on superhard TiN film are consistent with that of XRD analysis –This supports that our approach in using AIP system for producing superhard TiN film is feasible and that the relevant mechanisms found are valid. We believe that such findings contribute significantly to the scientific understanding in the field of thin film hardness enhancement.

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Acknowledgments The support from (i) Supported by Program for New Century Excellent Talents in University (NCET), (ii) A Foundation for the Author of National Excellent Doctoral Dissertation of PR China, (iii) Strategic Research Grant of City University of Hong Kong (CityU: 7002235), (iv) Major Project of International Scientific Cooperation Plan of China (2006DFB51260) is greatly acknowledged and (v) Open Research Foundation of National Scientific Drilling Laboratory of China University of Geosciences (NLSD200705).

References [1] G.S. Kim, S.Y. Lee, J.H. Hahn, B.Y. Lee, J.G. Han, J.H. Lee and S.Y. Lee, Surf. Coat. Technol. 2003, 171, 83-90. [2] W.M. Posadowski, Thin Solid Films 2001, 392 (2), 201-207. [3] U. Krause, M. List and H. Fuchs, Thin Solid Films 2001, 392, 196-200. [4] X.Yu, C. B.Wang, Y. Liu and D. Y. Yu, Acta. Metall. Sin. 2006, 42 (6), 662-666. [5] T. Hanabusa, K. Kusaka, T. Matsue, M. Nishida, O. Sakata and T. Sato, Vacuum 2004, 74, 571-575. [6] T. Nishikiori, T. Nohira and Yasuhiko Ito, Thin Solid Films 2002, 408, 148-154. [7] P. F. Mcmillan, Nature 2004, 430, 738. [8] S. Q. Hao, B. Delley, S. Veprek and C.Stampfl, Phys. Rev. Lett. 2006, 97, 086102. [9] K. Reuter and M. Scheffler, Phys. Rev. B 2002, 65, 035406. [10] S. Ma, J. Prochazkab, P. Karvankova, Q. Ma, X. Niu, X. Wang, D. Ma and K. Xu, Surf. Coat. Technol. 2005, 194, 143-148. [11] M. D. Huang, G. Q. Lin, Y. H. Zhao, C. Sun, L. S. Wen and C. Dong, Surf. Coat. Technol. 2003, 176, 109-114. [12] Novikov, R. Riedel, R. Solozhenko and Y. Zhao, Nat. Mater. 2004, 3, 576. [13] X.Yu, C. B.Wang, M. Hua, P. Tam, Y. Liu and D. Y. Yu, Surf. Rev. Lett. 2007, 14(4), 789-793. . [14] Ming-Hua Shiao, Sui-An Kao and Fuh-Sheng Shieu, Thin Solid Films 2000, 375, 163169. [15] T. Matsue, T. Hanabusa and Y. Ikeuchi, Vacuum 2004, 74, 647-653. [16] X.Yu, M.Hua, C. B. Wang, Z. Q. Fu and Y. Liu, Appl. Surf. Sci. 2007, 253 (7), 37053711. [17] X.Yu, C. B.Wang, M. Hua, Y. Liu and D. Y. Yu, Nanotechnology 2007, 18: 355710. [18] S. Veprek, G. J. Maritza, V. Heijman, P. Karvankova and J. Prochazka, Thin Solid Films 2005, 476, 1-29

In: Nanotechnology… Editors: C.J. Dixon and O.W. Curtines, pp. 525-557

ISBN: 978-1-60692-162-3 © 2010 Nova Science Publishers, Inc.

Chapter 17

EMBEDDED OPTICAL-ELECTRICAL NANOMATERIALES FABRICATED BY ION IMPLANTATION X.T. Zu and X. Xiang Department of Applied Physics, University of Electronic Science and Technology of China, Chengdu, 610054, People’s Republic of China

S. Zhu and L.M. Wang Department of Nuclear Engineering and Radiological Sciences, University of Michigan, Ann Arbor, MI 48109-2104, USA

Abstract Ion implantation provides a versatile and powerful technique for synthesizing nanometer-scale clusters embedded in the near-surface region of a variety of host materials. The embedded nanoparticles have attracted considerable attention because of their unique optical-electrical properties that are different from those of the bulk matrix. Metallic nanoparticles embedded in insulators have pronounced optical effects, including surface plasma resonance (SPR) absorption, and strong third-order nonlinear optical (NLO) susceptibility. The former suggests applications as optical filters, including eye-glass coatings. The latter has potential application in all-optical-memory or switching devices. Oxide nanoparticles have good photoluminescence. They have promising application in light-emitting devices. Magnetic metallic nanoparticles often show a ferromagnetic behavior with a larger coercivity than that of the corresponding bulk materials, which may provide potential application of the nanocomposite as magneto-optical materials for a high density magnetic data storage device. In this data review, nanoparticles embedded in insulators, e.g., Al2O3, MgO, YSZ and TiO2 single crystals, were fabricated by ion implantation and subsequent thermal annealing, including metallic Ni, Zn and their oxides, and intermetallic nanoparticles. Optical, magnetic and mircostructural properties of nanoparticles have been studied. The metallic nanoparticles have surface plasmon resonance absorption, and oxide nanoparticles show good photoluminenscence. The magnetic nanoparticles, e.g., metallic Ni and intermetallic CoxNi1-x nanoparticles, show strong ferromagnetism behaviors. The ion fluence can affect the concentrations and the intensities of the surface plasmon absorption of metallic nanoparticles. Ion flux is another important parameter to fabricate nanoparticles. An example of effects of ion flux on the nanoparticles has been presented in this data review. The relationship between

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X.T. Zu, X. Xiang, S. Zhu et al. annealing temperature and optical, magnetic and microstructural properties of nanoparticles has also been systematically studied.

I．Introduction Nanostructured materials are promising to play a dominant role in future technology as they possess different, and often unique, properties relative to their macroscopic counterparts. There has been a great deal of recent interest in the incorporation of nanoparticles into dielectric and semiconductor materials to form nanocomposites. Metallic nanoparticles embedded in insulators have been extensively studied because of pronounced optical effects, including surface plasmon resonance (SPR) absorption and strong third-order nonlinear optical (NLO) susceptibility [1]. These composites have drawn much attention due to applicability for all-optical-memory or switching devices and single electron transistors [2], etc. Magnetic metallic nanoparticles often show a ferromagnetic behavior with a larger coercivity than that of the corresponding bulk materials. Ferromagnetic nanoparticles have potential application in magnetic storage devices [3, 4]. Oxide nanoparticles have good photoluminescence. They have promising application in the light-emitting devices. Various synthesis methods have been attempted to synthesize these nano-phases, such as, surface sputtering [5], pulsed laser deposition [6, 7], sol–gel [8] and ion implantation [9–12]. Among these techniques, ion implantation is one of the most reliable and effective methods for the synthesis of nano-scaled particles. The implanted ions frequently precipitate as nanoparticles in controlled concentrations in the near surface regions of the host materials. Almost any element in the periodic table can be implanted into virtually any selected host material. The average precipitate size can be controlled by varying the implantation and annealing conditions (such as dose, dose rate, energy, temperature and annealing time) at pre-calculated depths of the host matrices [1]. Oxide crystals are often stable substrates used in a large range of technological applications. These high stabilities make them suitable candidates to allow the controlled formation of colloidal dispersions of metallic precipitates using ion implantation. In general, ion implantation techniques used to form nanoclusters may be categorized as follows [13]: (1) room temperature implantation, followed by high temperature annealing; (2) room temperature implantation at dosage above the threshold dose for spontaneous nanocrystals formation; (3) ion implantation at elevated temperatures. In this data review, metallic nanoparticles were prepared by means of the second; oxide nanoparticles were prepared by the second method of ion implantation and thermal annealing; the intermetallic nanoparticles participated spontaneously by sequential implantation of two ions at a high dose. This data review will be concerned primarily with the optical, magnetic and microstructural properties of metallic, oxide and intermetallic nanoparticles prepared in single crystals by room temperature ion implantation combined with subsequent thermal annealing. The substrates, metallic ion sources and implantation parameters used in this data review are listed in Table 1. X-ray diffraction (XRD), X-ray photoelectron spectroscopy (XPS), and transmission electron microscopy (TEM) analyses are used to characterize microstructural properties of nanoparticles. Optical absorption and room temperature photoluminescence (PL) measurements are used to obtain the optical properties of nanoparticles. The magnetic

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properties of the samples were characterized by a MPMS superconducting quantum interference device (SQUID) magnetometer. Table 1. Substrate and ion implantation parameters Material α-Al2O3 (0001) TiO2 (001) MgO (100) YSZ (001)

Ion Ion energy Ion fluence (cm-2) species (keV) Ni 64 1×1017 Zn 48 0.1, 0.5, 1, 5×1017 Ni 64 1×1017 Ni 64 1×1017 Ni 64 1×1017 Co+Ni

90

1×1017

Flux (μA/cm2) or dose rate Reference (ions/cm2 s) 2 5, 10 μA/cm [14, 15] 5 μA/cm2 [16, 17] 5 μA/cm2 [18] 5 μA/cm2 [19,20] 5 μA/cm2 [21] 2.93×1013 ions/cm2s for Co, [22] 2.63×1013 ions/cm2s for Ni

Abbreviations in This Data Review BF EDS EELS FC HAADF HREM MEVVA MPMS NLO PL SAED SPR SQUID STEM TEM UV-VIS XPS XRD YSZ ZFC

bright-field energy dispersive spectroscopy energy electron-loss spectroscopy field-cooling high-angle annular dark-field high-resolution electron microscopy metal vapor vacuum arc magnetic property measurement system nonlinear optical photoluminescence selected area electron diffraction surface plasma resonance superconducting quantum interference device scanning transmission electron microscopy transmission electron microscopy ultraviolet-visable X-ray photoelectron spectroscopy X-ray diffraction yttria-stabilized zirconia (with 9.5 mol. % Y2O3 in this datareview) Zero-field-cooling

II．Conclusion 2.1. Ni and NiO Nanoparticles Embedded in Single Crystals Ion implantations were conducted at room temperature using a metal vapor vacuum arc (MEVVA) implanter. The samples were tilted off-axis by around 7 degrees to avoid channeling implantation. After Ni ion implantation, the Al2O2 and YSZ samples were

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annealed in oxidization in order to study the formation of oxide nanoparticles, and the MgO sample was annealed in reducing atmosphere to study growing up of metallic nanoparticles.

2.1.1. Chemical Charge States of As-Implanted and Annealed Samples XPS measurements were performed to characterize chemical charge states of Ni element in as-implanted and annealed crystals. The C1S peak at 285.0 eV is used to calibrate the spectra. The samples had been etched 2 nm with Ar+ ion before measurements in order to remove surface contamination. 2.1.1.1. XPS Results of Al2O3 Figure 1 shows the XPS spectra of Ni2p3/2 energy level of the as-implanted and annealed Al2O3 crystals at etching depth of 2 nm, respectively [14]. The as-implanted spectra can be resolved into two Gaussian components. The peak at 852.3 eV (Figure 1a) is attributed to metallic Ni0, and the peak at 854.4 eV (Figure 1b) is due to Ni2+ (NiO). It is clear that the Ni element is prominent in charge state of metallic Ni0 in as-implanted crystals and Ni2+ in annealed crystals at 900 oC in ambient atmosphere. After etching 12 nm, there is no obvious change for the as-implanted and annealed samples.

a

864

b

856

848

864

856

848

Binding Energy/eV Figure 1. XPS spectra of Ni2p3/2 core level of Ni+-implanted Al2O3 at an ion flux of 5 μA/cm2 before (a) and after annealing at 900o for 1h in ambient atmosphere.

The XPS results above are obtained from the Ni-ion-implanted samples at ion flux of 5μA/cm2. The XPS result from Ni-ion-implanted Al2O3 at ion flux of 10μA/cm2 is shown in Figure 2. The peak at a binding energy of 856.6 eV can be attributed to Ni2+ in NiAl2O4 and the one at 862.8 eV is assigned to the well known shake-up satellite peak of Ni2+. The two peak positions and their energy difference 6.2 eV are just consistent with the previous study [23]. Thus, the XPS result indicates the formation of NiAl2O4 with the spinel structure when ion implantation conducted at a flux of 10μA/cm2.

Embedded Optical-electrical Nanomateriales Fabricated by Ion Implantation

873

864

855

529

846

Binding Energy/eV Figure 2. XPS spectra of Ni2p3/2 energy level of Ni+-implanted Al2O3 crystals at a flux of 10μA/cm2.

2 nm

a

12 nm

864

2 nm

b

12 nm

855

846 864 Binding Energy/eV

855

846

Figure 3. XPS spectra of Ni2p3/2 energy level of Ni ion-implanted (a) and annealed (b) YSZ crystals at 900oC in ambient atmosphere.

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2.1.1.2 XPS Results of YSZ Figure 3 shows the XPS spectra of Ni2p3/2 energy level of the as-implanted and annealed YSZ crystals at different etching depths, respectively [21]. The spectra can be resolved into Gaussian components. In the as-implanted spectrum at etching depth of 2 nm, the peak at 852.5 eV is attributed to metallic Ni0, the peak at 855.6 eV may be due to Ni2+ (NiO) or Ni3+ (Ni2O3), and the peak at 860.6 eV could be assigned to the shake-up satellite peak of Ni0, Ni2+ or Ni3+. But after etching 12 nm, the charge state of Ni is only Ni0. These results show that the implanted Ni in the surface is easy to be oxidized. This is not same as the Ni ion implanted αAl2O3 single crystals, whose Ni on the surface are mainly in charge state of metallic Ni0 [14]. In the annealed spectrum at etching depth of 2 nm, the peak at 855.5 eV shows Ni is only in charge state of Ni2+ (NiO) or Ni3+ (Ni2O3). The implanted Ni ion had entirely been oxidized in the surface of YSZ matrix. However, the metallic Ni0 (the Gaussian peak at 853.0 eV) appeared after etching 12 nm. This result shows that there is still some metallic Ni0 retained in the YSZ matrix even after annealing at 900 oC in ambient atmosphere. For the as-implanted and annealed TiO2 crystals, the XPS spectra of Ni are similar to those of YSZ. 2.1.1.3. XPS Results of MgO Figure 4 shows the XPS spectra of Ni2p3/2 energy level of the Ni ion implanted MgO crystals after annealing at 700 and 900 ºC in Ar+4% H2 atmosphere, respectively [20]. There is no obvious difference between the Ni2p3/2 energy level spectra at 700 and 900 ºC, i.e., the charge state of Ni did not change with the increasing annealing temperature. The peak at 852.9 eV is attributed to metallic Ni0, and the peak at 859.1 eV could be assigned to the shake-up satellite peak of Ni0. These results show that the implanted Ni is only in the charge state of Ni0 and the charge state of Ni was still metallic Ni0 after annealing at 900 ºC in Ar+4% H2 atmosphere.

a

864

b

855

846 864

855

846

Binding energy/eV Figure 4. XPS spectra of 1×1017 cm-2 Ni ion implanted MgO samples after annealing at 700 oC (a) and 900 oC (b) in Ar+4% H2 atmosphere, respectively.

531

5µA/cm2 Ni <111>

Al 2O3 <006>

unidentified

900oC

NiO <222>

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Intensity/ a.u.

800 oC

700 oC

600 oC

as-implanted

as-received

30

40

50

2θ/deg Figure 5. XRD traces (θ-2θ) of Al2O3 single crystals implanted at 5μA/cm2 and after annealing at different temperatures.

2.1.2. XRD Spectra of As-Implanted and Annealed Al2O3 Samples X-ray diffraction measurement was used to clarify the formation of metallic Ni, NiO and NiAl2O4 at ion flux of 5 and 10μA/cm2, respectively [15]. XRD traces (θ-2θ) of some samples were collected with a Cu Kα line of 1.54056 Å. Figure 5 shows XRD traces (θ-2θ) of asreceived and as-implanted crystals at a flux of 5μA/cm2 and annealed crystals at temperatures from 600 to 900 oC. In all the spectra, two diffraction peaks can be observed at ~40.6˚ and ~41.7˚, which were assigned to unidentified and (006) planes of the as-received Al2O3 crystals, respectively. For the as-implanted sample, a broad diffraction peak of metallic Ni appeared at ~44.3˚ indicating the formation of metallic Ni nanoparticles during ion implantation. After annealing at 600 oC in air, the XRD spectra show coexisted diffraction peaks of both Ni and NiO nanoparticles. As the annealing temperature increased, NiO nanoparticles grew up at the expense of Ni nanoparticles. When the annealing temperature reached 900 oC, most of the Ni nanoparticles were oxidized into NiO nanoparticles. The trace of the Ni nanoparticles was hardly detectable based on the XRD spectrum. This is consistent with the XPS result above.

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20

10µA/cm2

30

40

NiAl2 O3 <440>

Al 2O3 <006>

NiAl 2O3 <111>

Intensity/a.u.

532

50

60

70

2θ/deg Figure 6. XRD traces (θ-2θ) of Al2O3 single crystals implanted at 10μA/cm2 and after annealing at different temperatures.

In order to evaluated the mean grain size of Ni and NiO nanoparticles, the Scherrer formula [24] was used

D=

0.9λ B cos(θ B ) ,

where λ, θB, and B are the X-ray diffraction wavelength (1.54056 Å), Bragg diffraction angle and the full width at half maximum (FWHM) of diffraction peaks, respectively. The calculated grain sizes of Ni and NiO nanoparticles are listed in Table 2 for as-implanted and annealed samples. For the as-implanted sample, Ni nanoparticles have mean dimension of 4.9 nm, which will be proved by the following TEM measurement. The annealing effect on the growth of Ni nanoparticles is not obvious before the annealing temperature up to 800 oC, which is similar to Ni nanoparticles in the silica glass [9, 25]. The growth of NiO nanoparticles mainly occurred at annealing temperature above 800 oC. These results just explained the shift towards longer wavelength of UV absorption band, i.e., the optical band gap of NiO nanoparticles shifted towards lower energy due to quantum confinement effect as the grain size increased. Table 2. Average dimensions of Ni and NiO nanoparticles calculated from the XRD spectra after Ni ion implantation with an ion flux of 5μA/cm2. samples as-implanted annealed at 600°C annealed at 700°C annealed at 800°C annealed at 900°C

Ni particle size (nm) 4.9 5.1 5.5 9.2 —

NiO particle size (nm) — 10.8 11.6 19.9 21.5

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Figure 6 shows XRD traces (θ-2θ) of as-implanted crystal at an ion flux of 10μA/cm2. After ion implantation, two new diffraction peaks appeared at ~20.58˚ and ~64.56˚, respectively. According to the JCPDS card (No. 78-1601), these two diffraction peaks may be corresponding to (111) and (440) planes (2θ = 19.08˚ and 65.545˚) of NiAl2O4. The peak shifts with respect to the powder diffraction data may suggest an existing stress due to the lattice distortion after the Ni ion implantation. This XRD result proved the formation of NiAl2O4 spinel structure.

2.1.3. TEM Results of As-Implanted and Annealed Samples Nanoparticles of Ni precipitated spontaneously during ion implantation. In TEM measurements, a HAADF STEM (high-angle annular dark-field scanning transmission electron microscopy) technique was used besides a conventional bright-field imaging technique. For HAADF STEM imaging, intensity in the image is approaching a Z2 dependence on atomic number Z [26]. As an example, for the Ni ion implanted Al2O3, the local area where Ni element distributed will show brighter contrast because Ni has a larger Z than both Al and O. This suggests a HAADF STEM image provides chemical information on element distribution by its contrast (so called Z- contrast image). 2.1.3.1. Ni and NiO Nanoparticles in Al2O3 Figure 7 shows a bright-field and a HAADF STEM cross-sectional image indicating size and distribution of Ni nanoparticles embedded in Al2O3 formed by ion implantation at an ion flux of 5μA/cm2. In this study, Ni has a much higher Z than both Al and O; therefore the nanoparticles of Ni show bright contrast. As is shown in Figure 2, nearly spherical embedded nanoparticles are distributed from the surface to 30 nm below the surface, consistent with calculation results by SRIM 2000 code [27]. The size of nanoparticles ranges from 1 to 5 nm in diameter, which is consistent with the XRD result. Figure 8 is a HREM image showing the crystalline structure of nanoparticles of Ni in the surface of Al2O3 matrix. The image shows clearly the Ni-ion implanted area is amorphized entirely [14]. surface

a

surface

b

Ni Ni

Al 2O3

15 nm

Al2 O3

15 nm

Figure 7. A bright-field (a) and a HAADF STEM (b) cross-sectional image in the near surface of asimplanted Al2O3 matrix.

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Ni

Ni

Al2O3

22nm nm

Al2O3

Figure 8. A HREM image showing structure of Ni nanoparticles in the near surface of Al2O3 matrix. NiO

a

b

nanovoids

Al2O3 matrix

NiO nanoparticle

25 nm

5 nm Al2O3 matrix

Figure 9. A cross-sectional bright-field and a high-resolution TEM image from a sample annealed at 900 oC after Ni ion implantation with an ion flux of 5μA/cm2.

Figure 9 shows a cross-sectional bright-field (BF) and a high-resolution electron microscopy (HREM) image from a sample annealed at 900 oC after Ni ion implantation with an ion flux of 5μA/cm2. In the BF imaging (Figure 9a), the NiO nanoparticles grew to 6-20 nm in diameter with irregular shape, which is consistent with the XRD result. In addition, the nanoparticles migrated towards the surface of the crystal after annealing. A high density of voids formed below the nanoparticles during the thermal annealing process. The amorphous area of Al2O3 matrix was partially recrystallized. The HREM image (Figure 9b) demonstrates the single crystalline nature of the NiO nanoparticle [15]. 2.1.3.2. Ni and NiO Nanoparticles in YSZ TEM measurement showed that no obvious nanoparticles precipitated after Ni ion implantation. And YSZ matrix did not amorphize after ion implantation. Although XPS measurements have detected the metallic Ni, TEM measurement did not observe nanoparticles precipitated in as-implanted YSZ matrix. This may be due to the very small crystals or none, and even show amorphous conditions. Figure 10 shows a cross-sectional

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bright-field transmission electron microscopy image of the annealed sample at 900 oC. After thermal annealing at 900 oC for 0.5 h, nanoparticles with size ranging from 4 to 12 nm can be observed. The large particles distributed in the surface of the YSZ crystal. Figure 11 is a highresolution electron microscopy (HREM) image clearly showing the crystalline structure of nanoparticles in the surface of YSZ matrix. The nanoparticles are nearly in shape of sphere. In order to clarify which oxide the nanoparticles are, the selected area electron diffraction (SAD) pattern had been obtained in Ni-implanted region in YSZ sample after annealing at 900 oC (shown in Figure 12). After indexing the SAD pattern (the circled spots in Figure 12) the nanoparticles were identified to be NiO [21].

surface

NiO

25 nm Figure 10. A cross-sectional bright-field transmission electron microscopy imaging of the annealed sample at 900 oC.

NiO

YSZ

3 nm Figure 11. A HREM image showing the crystalline structure of nanoparticles in the surface of annealed YSZ matrix.

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200YSZ 022YSZ 111NiO

202NiO

Figure 12. A SAD pattern of the annealed YSZ matrix showing the formation of NiO nanoparticles.

2.1.3.3. Ni Nanoparticles in MgO

Figure 13. A cross-sectional HAADF STEM images indicating the nanoparticles of Ni in the surface region of the MgO single crystal after Ni ion implantation (a) and followed by annealing at 900oC for 0.5 h under Ar+4%H2 atmosphere.

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Figure 14. Composite selected area electron diffraction (SAED) patterns indicating orientation relationship between Ni nano-particles and the MgO matrix (a); A high resolution TEM (HREM) micrograph showing the crystalline characteristics of Ni nano-particles in the annealed MgO single crystal (b).

Figure 13 shows HAADF images in MgO matrix after Ni ion implantation (a) and subsequent annealing at 900 oC under Ar + 4% H2 atmosphere (b), respectively. In asimplanted MgO, the formation of Ni nano-particles is evident and the particle ranges 3–5 nm in size. The grown and coalescence of these nanoparticles occurred, after thermal annealing and the particles size increased to 8–10 nm. Moreover, larger rods of Ni with 20 nm in length precipitated in the surface of MgO matrix. Besides, the nanoparticles spread into the deeper regions of MgO matrix after thermal treatment. The Ni nanoparticles have specific orientation relationship with the MgO matrix in both of as-implanted sample and thermal annealed sample after ion implantation. As indicated by the composite selected area diffraction (SAD) pattern in Figure 14 (a), the orientation relationship between the Ni particles and MgO matrix was determined as: <001>Ni║<001>MgO and {010}Ni║{010}MgO. Extra satellite spots also appear around the strong spots from the original MgO single crystal in the composite SAD patterns. These satellite spots result from reflections attributable to double diffraction since Ni particles were embedded in the MgO single crystal with the same orientation relationship but different lattice spacing. The double-diffraction reflections are directly responsible for the Moiré fringes in the high resolution TEM images (Figure 14 (b)). As shown in Figure 14 (b), Morié fringes appeared in the place where embedded Ni particles were overlapped with the MgO matrix, due to lattice spacing mismatch between the two phases. The spacing of Morié fringes, dm, can be calculated by the following equation:

dm =

d2 d 1− 2 d1

,

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where d1 and d2 are lattice spacings of two overlapping crystals. The lattice spacing for Ni(002) is dNi,002 = 0.176 nm (lattice parameter aNi = 0.352 nm), and for MgO(002) dMgO,002 = 0.211 nm (lattice parameter aMgO = 0.422 nm). Thus according to the equation above, dm is 1.06 nm, which is consistent with the measured value in the image [19]. 2.1.3.4. Ni Nanoparticles in TiO2 Ni Nanoparticles precipitated spontaneously during ion implantation from the surface of TiO2 matrix to a depth of 50 nm, as revealed by the cross-sectional HAADF Z-contrast image shown in Figure 15(a). In this HAADF STEM image, Ni has a higher Z than both Ti and O, so the nanoparticles of Ni show a brighter contrast. The dimensions of nanoparticles ranged from 3 nm to 10 nm. Some elongated precipitates up to 20 nm in length were observed in the near surface of TiO2. These elongated Ni precipitates are observed with the round shape in a plan-view TEM image. The implantation region of TiO2 matrix has been damaged and amorphized to 60 nm in the depth below the surface. Figure 15(b) is a HREM image of Ni nanoparticles embedded in the TiO2 matrix, revealing a well-developed crystalline structure of Ni nanoparticles after ion implantation without a thermal treatment [18].

Figure 15. A cross-sectional HAADF STEM image of Ni nanoparticles in the near surface of a TiO2 single crystal (a); and a HRTEM micrograph showing a crystalline Ni nanoparticle in the amorphous TiO2 matrix (b).

2.1.4. Optical Absorption of As-Implanted and Annealed Samples The optical absorption spectra were measured by a SHIMADZU UV-2550 spectrophotometer at room temperature, with a deuterium lamp for UV and a tungsten halogen lamp for the visible region. The wavelength used in the experiment ranged from 200 to 1000 nm. 2.1.4.1. Optical Absorption of Ni-Implanted and Annealed Al2O3 The optical absorption spectra of as-received, as-implanted, and annealed Al2O3 single crystals with an ion flux of 5μA/cm2 are shown in Figure 16. The spectra have vertically been offset to avoid overlapping except that of the as-received crystal. The absorption spectrum of

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the pure crystal was a smooth line in the visible waveband due to a wide band gap ~9 eV. A broad absorption band peaked at 400 nm in the as-implanted crystal [14]. According to the XPS, XRD and TEM results above, the absorption band can be ascribed to the surface plasma resonance absorption of metallic Ni nanoparticles. As the annealing temperature increased, the absorption band shifted towards the longer wavelength. After annealing at 800 oC, this absorption band was absent. At the same time, the crystals turned colorless. However, a new absorption shoulder in the UV region began to appear after annealing at 600 oC. As the annealing temperature reached 800 oC, the UV absorption shoulder peak evolved into an absorption band and its peak position shifted to the longer wavelength of 306 nm (4.05 eV). There is no detectable change for the peak position of the absorption band after annealing at higher temperatures. This UV absorption band was related to the formation of NiO, since NiO is an insulator with a band gap of ~4eV (310 nm) [9]. These results have been proved by the XPS and XRD results. And TEM showed the microstructure of NiO and the recrystalization of Al2O3 matrix [15]. 0.6 5μA/cm2

as-implanted

0.4

400 oC

Absorbance/a.u.

500 oC 600 oC 700 oC 0.2

800 oC 900 oC 1000 oC as-received

0 250

400

600

800

Wavelength/nm Figure 16. Optical absorption spectra of Al2O3 single crystals after Ni ion implantation at 5μA/cm2 and annealing at different temperatures.

2.1.4.2 Optical Absorption of Ni-Implanted and Annealed YSZ The optical absorption spectra of Ni-implanted and annealed YSZ single crystals are partly shown in Figure 17. The spectra have vertically been offset to avoid overlapping except

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that of the annealed crystal at 900 ºC. In the spectrum of as-implanted crystal, there is a very broad and weak absorption band ranging from 400 to 700 nm. The absorption curves of annealed crystals are similar to that of as-implanted crystal up to 200 oC. The absorption intensity begins to decrease after annealing at 200 oC, and the absorption band disappears after 250 oC. There is no change for the absorption spectra at annealing temperature of 300~900 oC, similar to that of as-received YSZ [21]. 1.2 as-implanted 200 oC 250 oC 300 oC 900 oC

absorbance/a.u.

1.0 0.8 0.6 0.4 0.2 0

200

400

600

800

wavelength/nm Figure 17. Optical absorption spectra of Ni ion-implanted and annealed YSZ crystals at different annealing temperatures.

YSZ is known to have a band gap of 5.6 eV (222 nm). So its absorption spectrum is nearly a line in the visible region. According to the TEM results above, the broad and weak absorption band ranging from 400 to 700 nm is not due to the SPR absorption of metallic Ni nanoparticles. Although XPS measurements have detected the metallic Ni, TEM measurement did not observe nanoparticles precipitated in as-implanted YSZ matrix. This may be due to the very small crystals or none, and even show amorphous conditions. So the metallic Ni in the as-implanted crystals does not show metallic behavior. In addition, the broad absorption band disappeared just after annealing at 250 oC. At this temperature the metallic Ni clusters did not begin to grow. So the broad absorption band can be ascribed to the point defects and their clusters induced by ion implantation. Just as the Xe ion-implanted YSZ crystals [28], this absorption band may be associated with the combination of electrons trapped at oxygen vacancies and oxygen ions with trapped holes. Metallic Ni clusters grow with the increasing annealing temperature. At the same time, metallic Ni clusters are partly oxidized into NiO nanoparticles. However, the optical absorption spectra did not detect the absorption of Ni or NiO nanoparticles. A possible reason is that the Ni and NiO coexist in the YSZ single crystals and each concentration is low.

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2.1.4.3. Optical Absorption of Ni-Implanted and Annealed MgO The optical absorption spectra of as-implanted and annealed MgO single crystals are shown in Figure 18. The spectra have vertically been offset to avoid overlapping except that of the annealed crystal at 900 ºC. In the spectrum of as-implanted crystal, there are two weak absorption bands centered at ~360 nm and ~575 nm, respectively. They have been identified with the F2 centers and the V-type centers in magnesium sublattice (magnesium vacancies), respectively, which is consistent with the Ag+ and Ni+ ion implanted MgO single crystals [29,30]. Upon heat treatment above 700 °C these two absorption bands are completely annihilated by the recombination of point defects. During the annihilation of centers, a new absorption band at ~430 nm formed gradually. Form the XPS and TEM results, it can be concluded that the broad absorption band is related to the metallic Ni nanoparticles, i.e., the surface plasmon resonance (SPR) absorption. After annealing above 700 °C, the absorption maximum shifts to a longer wavelength with the increasing annealing temperature, which is related to the growth of Ni nanoparitcles [20]. 1.5

Absorbance/a.u.

1.0

as-implanted 200 °C 0.5

400 °C 600 °C 700 °C 800 °C 900 °C

0

200

400

600

800

1000

Wavelength/nm Figure 18. Optical absorption spectra of MgO single crystals after Ni ion implantation and annealing at different temperatures.

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2.1.5. Magnetic Properties of Ni Nanoparticles The magnetic properties of the samples were characterized by a magnetic property measurement system (MPMS) superconducting quantum interference device (SQUID) magnetometer at 10 and 300 K. The zero-field-cooling (ZFC) and field-cooling (FC) curves for detecting superparamagnetism of the nanoparticles were measured in an applied magnetic field of H = 100 Oe. The ZFC curve was achieved by cooling the sample initially in a zero field to 10 K, and magnetization was recorded in an applied magnetic field where H = 100 Oe as the temperature increased. The FC magnetization was measured by gradually cooling the sample from 300 K to 20 K, and the magnetization was recorded in the presence of a 100 Oe field. 2.1.5.1. Magnetic Ni Nanoparticles in MgO Figure 19 shows a magnetization plot as a function of magnetic field at 10 K in the Niimplanted MgO sample. The applied magnetization field H is parallel to (100) plane of the MgO single crystal, i.e. (100) of Ni nano-particles (Figure 14), owing to their orientation alignment. As shown in the hysteresis loop measured at 10 K, the coercivity, Hc, was about 195 Oe, which is larger than ~150 Oe of randomly oriented Ni particles in Al2O3 [31], because crystallographically oriented particles have stronger tendency to retain its magnetic moments than that of randomly oriented particles under a reversing magnetic field [31]. The magnetization curves at 10 K were completely saturated as the applied field increased to H = 5000 Oe (not shown in the figure). No coercive force Hc was observed in the sample at 300 K [19].

Figure 19. A magnetic hystersis loop of Ni nanoparticles in the MgO single crystal at 10 K after ion implantation. The coercivity Hc was about 195 Oe at this temperature.

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Figure 20. ZFC and FC magnetizations as a function of temperature for Ni nanoparticles in asimplanted and after subsequent annealed MgO samples. Curves were taken in the ZFC and FC processes at H=100 Oe.

The zero-field-cooled (ZFC) and field-cooled (FC) magnetization as a function of temperature are shown in Figure 20 for both as-implanted and the annealed samples. As shown in Figure 20, the FC magnetization increases monotonically with the decrease of temperature. However, the magnetization increases at first, then decreases with an increasing of temperature in ZFC curve. The temperature, at which the maximum in ZFC magnetization occurs, is characterized as the blocking temperature (TB). The blocking temperature TB of Ni nanoparticles was determined to be ~35 K in the as-implanted sample and above 300 K in the annealed MgO sample, respectively. These Ni nanoparticles exhibit superparamagnetic properties above the blocking temperature. This behavior is consistent with the result from the as-implanted sample that shows almost immeasurable coercivity and remanence at 300 K. However, superparamagnetism of the Ni nanoparticles in the annealed MgO sample persists to above 300 K, and these Ni nanoparticles remain ferromagnetic at room temperature [19]. TB has a relation with particle size, based on the equation: T B = Keff V/25kB, where Keff the effective anisotropy related to the magnetocrystalline anisotropy and to the shape anisotropy, V is the volume of a particle, kB is Boltzmann constant [32]. If we use the magnetocrystalline anisotropy K1 value (-8×105 erg/cm3) in [32] to replace Keff, the Ni particles size that contributed to the measured magnetic property can be roughly calculated to be ~3.3 nm in as-implanted sample and 6.8 nm in thermal annealed sample. This is close to

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the TEM observations above. It should be noted that the calculation neglects the shape anisotropy caused by large rod-shape Ni precipitates in the annealed sample, which may cause magnetic anisotropy and increase average particles size contributing to the magnetic property. 2.1.5.2. Magnetic Ni Nanoparticles in TiO2 Figure 21 shows the ZFC and FC magnetization as a function of temperature of Ni nanoparticles in TiO2 matrix. Figure 21 clearly shows the non-zero difference between the FC and ZFC data, indicating the hysteresis while eliminating any para- and diamagnetic contributions. In these ZFC/FC curves, the blocking temperature, TB, is ~85 K. Above the blocking temperature, the magnetization is unstable and the sample loses all its hysteric responses. The two insets in Figure 21 showing magnetization plots as a function of magnetic field (M-H) at 10 K and 300 K, respectively. The applied magnetization field H is parallel to (100) of the TiO2 single crystal. The magnetic hysteresis loop at 10 K after correcting paramagnetic contribution from TiO2 matrix shows a ferromagnetic behavior with coercivity, Hc, equal to ~ 210 Oe [18]. This value is larger than that of Ni nanoparticles in MgO (195 Oe) above. In addition, a coercivity, Hc, equal to ~ 270 Oe was obtained at 10 K in Ni-implanted YSZ crystal when the applied magnetization field H is parallel to (001) of the YSZ single crystal [21]. The variation of coercivity strongly depends on the grain size. Since nanoparticles less than a critical size are in single domain states, their coercivities are higher than that of a common bulk sample [33]. This large coercivity can be explained by the nanosize effect and the single domain structure (the critical size of the single domain for spherical Ni particles was ~ 42 nm) [34, 35]. The magnetization curves at 10 K was completely saturated at applied fields H = 4000 Oe. The remnant magnetization of the saturation magnetization (Mr/Ms) is about 40% at this temperature. The coercive force and remnant magnetization of Ni nanoparticles were not observed as the temperature increased to 300 K, confirming the superparamagnetic behavior of the nanoparticles above the blocking temperature. Since some nano-particles are not truly spherical, the shape anisotropy has contribution to the superparamagnetism. Langevin function can be used to calculate the true magnetic moment of each particle for superparamagnetic particles as [36]:

M ( H / T ) = M 0 [coth( M 0 mH / k BT ) − (k B T / M 0 mH )] where, M0 is the saturation magnetization (emu/g), m is the mass of individual particle (gram), and kB is Boltzmann constant. The saturation magnetization in our sample was about 25 emu/g from the hysteresis loop at 300 K. Figure 22 shows magnetization vs. applied magnetic field of Ni nanoparticles at 300 K (solid circles) and the best fit for the Langevin function (solid line). From the data fitting, we calculate the average grain size of 9.4 nm for the Ni nanoparticles contributing to magnetic moment, which is in the range of particle size measured by TEM observation. The mean-magnetic moment per particle of the sample was calculated to be 11064μB [18].

Embedded Optical-electrical Nanomateriales Fabricated by Ion Implantation

M (×10-4 emu)

1

FC 4.0

0.5 0 T=300K

-0.5 -1 -3 -2 -1

ZFC M (×10 -4 emu)

Magnetization (×10-5, emu)

6.0

545

2.0

0

1

2

3

H (kOe) 1 0 T= 10 K

1

-3 -2 -1 0 1 2 3

0

H (kOe)

0

50

100

150

200

250

300

Temperature (K)

Magnetization (emu/g)

Figure 21. Temperature dependence of the ZFC and FC magnetization curves for the Ni implanted TiO2 sample with magnetic hysteresis loops at 10 K (bottom inset) and 300 K (top inset).

20 10 0 -10 -20 -30 -20

T=300K

-10 0 10 Applied Magnetic field (kOe)

20

Figure 22. Measured (solid circle) and the Langevin function fitted (solid line) magnetization (M) vs. magnetic filed (H) at the room temperature.

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2.2. Zn and ZnO Nanoparticles Embedded in Al2O3 Ion implantations were conducted at room temperature using a metal vapor vacuum arc (MEVVA) implanter. The samples were tilted off-axis by around 7 degrees to avoid channeling implantation. After Zn ion implantation, the Al2O2 samples were annealed in oxygen atmosphere in order to fabricate zinc oxide nanoparticles. ZnO is well-known as a versatile wide band gap (~3.3 eV) semiconducting material with a large exciton binding energy of 60 meV, which allows excitonic recombination and optically pumped laser oscillations even at the room temperature. Following the demonstration of blue-green light emitting diodes (LEDs) and lasers using II–VI compounds, ZnO has been intensively studied for optoelectric applications in both the visible and ultraviolet (UV) regions. Optical absorption and TEM measurements are same as those of the Ni-ion-implanted samples above.

2.2.1. Optical Absorption of Zn Nanoparticles Fabricated with Different Fluences The optical absorption spectra of as-implanted α-Al2O3 single crystals at fluences of 1×1016, 5×1016, 1×1017, and 5×1017 cm-2 are shown in Figure 23. The spectra have been offset to avoid overlapping except that of the as-implanted spectrum at fluence of 1×1016 cm-2. After Zn+ ion implantation at a fluence of 1×1016 cm-2, a weak absorption peak appeared at ~260 nm. This absorption peak becomes clear gradually with the increasing fluences. At the same time, the peak wavelength shift to longer wavelength linearly (dash line in Figure 23). The absorption peak shifts to ~285 nm when the ion fluence up to 5×1017 cm-2. The strong and broad peak at 260-285 nm is due to surface plasma resonance absorption of metallic Zn nanoparticles [16]. Actually, the similar absorption peaks have been observed in SiO2 glass and MgO single crystal matrices (listed in Table 3). 0.8 5×1017 1×1017 5×1016 1×1016

Absorbance/a.u.

0.6

0.4

0.2

0

200

400

600

800

Wavelength/nm Figure 23. Optical absorption spectra of Zn+-implanted Al2O3 single crystals at different fluences.

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As a whole, the SPR peaks in α-Al2O3 and MgO lie at longer wavelength than that in SiO2 matrix. This is because that α-Al2O3 and MgO have larger refractive indexes (1.76 for αAl2O3 and 1.74 for MgO) than SiO2 (1.52). According to the Maxwell-Garnett (MG) [41] and Mie [42] theory, the SPR peak would shift towards low energy with increasing refractive index of surrounding medium. The size of nanoparticles is another factor to influence the SPR peak wavelength, which will shift to a longer wavelength with the increasing crystalline size. The SPR absorption peaks in Table 3 are observed in the as-implanted matrices except that in MgO which was observed after annealing at 1150 K, which shift to a longer wavelength due to the growth of nanoparticles. Table 3. SPR peaks of metallic Zn nanoparticles in several insulator matrices matrix

SiO2

∗

MgO

α-Al2O3

Zn+ ion/cm-2

SPR peak

reference

1×1017

4.8 eV/259 nm

[10]

1×1017

4.86 eV/255 nm

[37]

1×10

17

5.3 eV/234 nm

[38]

3×10

17

4.86 eV/255 nm

[39]

1×1017

4.2 eV/295 nm

[40]

1×1017

4.56 eV/272 nm

17

4.35 eV/285 nm

5×10

[17]

2.2.2. Optical Absorption of Zn-Ion-Implanted and Annealed Al2O3 The optical absorption spectra of the as-received, as-implanted with a fluence of 1×1017 cm-2 and annealed α-Al2O3 single crystals are shown in Figure 24. The spectra have been shifted vertically to avoid overlapping except that of the as-received sample. The SPR absorption peak slightly shifted towards the longer wavelength due to the growth of nanoparticles with the increasing annealing temperature. The spectrum of annealed sample at 500 °C is a transition one because it includes both the absorption peak of Zn nanoparticles and the absorption edge of ZnO nanoparticles. After annealing at 600 °C for 1 h, the SPR absorption peak disappears and a clear absorption peak appears at ~360 nm, which is consistent with the exciton absorption of ZnO [10, 17, 40]. The intensity of this exciton absorption peak decreases with the further increasing annealing temperature. Apparently, this intensity decrease does not indicate a decrease of Zn content in the sample. It may be due to the decreased concentration of ZnO nanoparticles because of the formation of ZnAl2O4 spinel during thermal annealing in O2 atmosphere [43]. The microscopic morphology of Zn and ZnO nanoparticles has been characterized by TEM imaging with the main results shown below.

∗

Observed after annealing at 1150 K.

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Absorbance/a.u.

0.8

0.6

0.4 600 oC 700 oC

0.2

800 oC 900 oC

0

400 600 Wavelength/nm

200

800

Figure 24. Optical absorption spectra of Zn+-implanted and annealed Al2O3 single crystals at different temperature. surface

amorphous

Crystalline matrix

5 nm

20 nm

Figure 25. A cross-sectional bright-field TEM image (a) and a high resolution TEM (HREM) image (b) of Zn nanoparticles embedded in α-Al2O3 formed by Zn+ ion implantation at a dose of 1×1017 cm-2. The polycrystalline ring in the bottom inset in (a) showing the random orientation of the Zn nanopartilces precipitated after the ion implantation.

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2.2.3. TEM Results of Zn and ZnO Nanoparticles Figure 25 shows a cross-sectional bright-field TEM image (a) and a high resolution TEM (HREM) image (b) of Zn nanoparticles embedded in α-Al2O3 formed by Zn+ ion implantation at a dose of 1×1017 cm-2. Nearly spherical embedded Zn nanoparticles of 3-10 nm in diameter are observed and the ion-implanted area is amorphized. The insert one in Figure 25a is a selected area electron diffraction (SAD) pattern. The polycrystalline ring can be observed in the SAD pattern showing the random orientation of the Zn nanopartilces precipitated after the ion implantation [16]. surface

ZnO

voids recrystallized Al2O3 100 101

20 nm

crystalline matrix

b a

Figure 26. A cross-sectional bright-field TEM image (a) and selected area electron diffraction (SAED) pattern (b) of ZnO nanoparticles embedded in α-Al2O3 formed after annealing at 600 °C.

Figure 26 contains a cross-sectional bright-field TEM image (a) and selected area electron diffraction (SAED) pattern (b) of ZnO nanoparticles embedded in α-Al2O3 formed after annealing at 600 °C. The ZnO can be confirmed by the SAED pattern. Figure 27 is a high resolution TEM image of ZnO nanoparticles, showing the nanoparticles of 10-12 nm in dimensions. The morié fringes indicate the precipitation of ZnO nanoparticles after annealing. It is clear that the ZnO nanoparticles formed close to the surface of the α-Al2O3 single crystal, with a depth shallower than the projectile range of implanted Zn atoms. It indicates that the Zn atoms migrated towards the surface of the crystal during the annealing in oxygen atmosphere. At the same time, some high density of large voids (labeled in Figure 26a) is observed in the near-surface region due to the migration and precipitation of irradiation induced vacancies. The recrystallized Al2O3 grains have different orientation relationship with the original matrix, which can be observed in Figure 28. However, the recrystallization was not complete after annealing at 600 °C for 1 h. The front surface area (labeled in Figure 28) remained amorpous. Apparently, longer annealing time or higher annealing temperatures are needed to complete the recrystallization [17].

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5 nm

Figure 27. A high resolution TEM image of ZnO nanoparticles, showing the nanoparticles of 10-12 nm in dimensions.

amorphous Al2O3

recrystallized Al2O3

2 nm

Al2O3 matrix

Figure 28. A high resolution TEM image showing the recrystallized Al2O3 with different orientation relation with the matrix.

2.2.4. PL of ZnO Nanoparticles Figure 29 shows the photoluminescence (PL) spectra of the as-implanted crystal and the annealed crystal at 600 °C using a He–Cd laser excitation at 325 nm line at room temperature. There is a very weak PL band peaked at ~470 nm in the as-implanted crystal. It may be due to the photoluminescence combination of both F (PL at 3.0 eV) and F2 (PL at 2.4 eV) centers coexisted in Al2O3 induced by ion implantation [44]. PL spectrum of the annealed crystal shows two PL peaks, one at 370 nm, and the other at 500 nm, which have also been observed

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in Zn+ ion implanted SiO2 and CaF2 [37, 39]. The UV emission is the characteristic PL peak ascribed to ZnO free-exciton recombination at room temperature, which confirms the formation of ZnO nanoparticles after thermal annealing. The green emission at ~500 nm may originate from the deep levels [45, 46]. The deep-level emission at around 2.5 eV is associated with either surface state emission or excess Zn interstitials (or oxygen vacancies). Up to now, both UV and green PL peaks of embedded ZnO nanoparticles have been observed in SiO2, CaF2 and Al2O3 matrices by ion implantation and thermal annealing (listed in Table 4). 400

PL intensity/a.u.

300

b

200 100

a 0 350 400 450 500 550 600 650 700 750 800

Wavelength/nm Figure 29. Photoluminescence spectra of the as-implanted (a) and the annealed (b) crystal at 600 °C at room temperature.

Table 4. PL of ZnO nanoparticles embedded in several insulator matrices SiO2 CaF2 Al2O3

700 oC 1h 700 oC 2h 400 oC 45min 500 oC 45min 700 oC 45min 600 oC 1h

375 377 384 372 379 370

500 500 500 500

1:2 1:7 2:1 3:1

[37] [39] [47] [17]

2.3. Intermetallic CoxNi1-x Nanoparticles Embedded in YSZ Sequential ion implantation in dielectric matrix determines three different cluster morphologies: separated systems, alloy clusters and core–shell clusters. Sequential ion implantations with different elements were performed mostly on silica substrates, but also on quartz and sapphire [48]. In this data review, the intermetallic CoxNi1-x nanoparticles embedded in YSZ were synthesized by sequential implantation of 90 keV Co and Ni ions at room temperature. TEM and magnetic measurements were utilized to analyses the magnetic CoxNi1-x nanoparticles [22].

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2.3.1. TEM Images of CoxNi1-x Nanoparticles A bright field cross-sectional TEM image (Figure 30a) shows that the nanoparticles precipitated spontaneously during ion implantation in the near surface of the YSZ matrix. The typical nanoparticle size ranged from 3 to 10 nm. In the high concentration region of implanted Co and Ni ions (~25–40 nm below the surface, consistent with the results from SRIM 2000 code [27]), some large precipitates up to 25 nm in length formed parallel to the surface of YSZ. High-resolution TEM images clearly revealed that the elongated particles consisted of individual crystals joined across twin boundaries. The (111) twin planes are shown in Figure 30b. The selected area diffraction pattern inserted in Figure 30a exhibits obvious extra diffraction spots from the precipitates, which are circled in the SAD pattern. The precipitates are cubic with a>0.35 nm. The phases consistent with this lattice parameter include cubic Co, Ni, solid solution CoxNi1-x, as well as intermetallic phases, CoNi and Co3Ni7 [49].

Figure 30. Bright-field cross-sectional TEM (a) and high resolution TEM (b) micrographs showing nanoparticles of CoxNi1-x in YSZ single crystal; SAD pattern (inset in (a)) indicating the orientation of the nanoparticles to the matrix.

Energy filtered elemental mapping images of Co, Ni, and O using the L2,3 or K edges in a corresponding energy electron-loss spectroscopy (EELS) spectrum, are shown in Figure 31. The bright contrast of the nanoparticles in the Co and Ni maps indicates that the nanoparticles contained both Co and Ni. These nanoparticles do not contain oxygen as indicated by the dark contrast in nanoparticles in the O map image. Energy dispersive spectroscopy (EDS) analysis showed the composition ratio of Co over Ni ranges from 0.8 to 1. Therefore, the nanoparticles are CoxNi1-x solid solution. The orientation relationship between aligned nanoparticles and the YSZ matrix are (200)YSZ║(200) CoxNi1-x, and [011]YSZ║[011] CoxNi1-x. Based on calculations using the SRIM code, the maximum level of damage reached 300 dpa in the YSZ matrix; however, the YSZ matrix still retained its crystallinity.

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Figure 31. (a) Bright-field image and elemental mapping images of (b) Co, (c) Ni, and (d) O, respectively, indicating the compositional distribution of the nanoparticles.

2.3.2. Magnetic Properties of CoxNi1-x Nanoparticles Magnetization plots as a function of magnetic field at both 10 and 300 K are shown in Figure 32. The magnetization field H applied is parallel to (001) of the YSZ single crystal, i.e. (001) of CoxNi1-x, due to its orientation alignment. The coercivity, Hc, at 300 K was measured to be ~100 Oe. The coercivity increased to 260 Oe as the temperature decreased to 10 K.

Figure 32. Magnetization vs field applied at temperatures of 10 and 300 K. The field applied is parallel to the (001) plane of the YSZ matrix.

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These values are consistent with previously reported particle-size dependent coercive forces in CoNi alloys [50, 51]. The magnetization curves at 10 and at 300 K were completely saturated at applied fields of H = 800 and 5000 Oe, respectively. The saturation magnetization, Ms, increased 15% as the temperature decreased from 300 to 10 K. The remanent magnetization Mr /Ms was 0.41 at 300 K and its value increased to 0.77 at 10 K. The variation of coercivity strongly depends on the grain size. Since nanoparticles less than a critical size are in single domain states, their coercivities are higher than that of a common bulk sample [33]. Previous investigations reported that the critical size with the largest coercivity lies in the 30–40 nm range for Co50Ni50 [50]. The coercivity of our samples can be attributed to this nanostructural effect. The mismatch between the matrix and the implanted CoxNi1-x particles probably induces stress, which also affects the coercivity [52]. Nanoparticles are expected to illustrate superparamagnetic properties.

Figure 33. Temperature dependence of the magnetization. The curves show that in the ZFC and FC processes at H = 100 Oe.

The blocking temperature can be readily characterized by ZFC and FC magnetization. Figure 33 shows the temperature dependent magnetization curves under ZFC and FC processes at H=100 Oe. No blocking temperature within 300 K was observed. Since TEM analysis has confirmed that the nanoparticles are not oxidized, the results from the ZFC and FC curves suggest that the sample might not be blocked at low temperature. This is reasonable if particle-size and shape effects are considered. As observed in the TEM micrograph (Figure 30), the particle size is widely distributed from 3 to 10 nm, and some longer particles up to 25 nm have formed from a sequence of twin planes roughly perpendicular to the length of the particle. They might not be in single domain states. The blocking temperatures therefore are different and may overlap. The critical temperature, i.e., the blocking temperature, TB, is also roughly given as the equation: TB = Keff V/25kB [32]. Based on this equation, TB = ~227 K, if we use the magnetocrystalline constant K1 value (1.5×105 J/m3) from Ref. [53] for Keff and a maximum diameter of 10 nm for the nanoparticles. The result of this calculation differs from the experimental observation because K1 only represents the magnetocrystalline anisotropy contribution. Since some nanoparticles are not truly spherical, the contribution by shape

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anisotropy to Keff should also be considered. On the other hand, the elongated particles formed by repeating twins may result in interparticle dipolar interactions. Thus, supermagnetism of the CoxNi1-x nanoparticles may persist to above 300 K, and the nanoparticles will remain ferromagnetic at room temperature.

References [1] Meldrum, A.; Haglund Jr., R. F.; Boatner, L. A.; White C. W. Adv. Mater. 2001, vol 13, 1431-1444. [2] Chakraborty, P. J. Mater. Sci. 1998, vol 33, 2235-2249. [3] Beecroft, L. L.; Ober,C. K. Chem. Mater. 1997, vol 9, 1302-1317. [4] Meldrum, A.; Boatner, L. A.; White, C. W. Nucl. Instrum. Meth. B, 2001, vol 178, 7-16. [5] Yang, C. J.; Kim, K. S.; Wu, J. M. J. Appl. Phys. 2001, vol 90, 5741-5746. [6] Gonzalo, J.; Perea, A.; Babonneau, D.; Afonso, C. N.; Beer, N.; Barnes, J. P.; PetfordLong A. K.; Hole, D. E.; Townsend, P. D. Phy. Rev. B 2005, vol 71, 125420. [7] Dempsey, N. M.; Ranno, L.; Givord, D.; Gonzalo, J.; Serna, R.; Fei, G. T.; PetfordLong, A. K.; Doole, R. C.; Hole, D. E. J. Appl. Phys. 2001, vol 90, 6268-6274. [8] Fonseca, F. C.; Goya, G. F.; Jardim, R. F.; Carreno, N. L. V.; Longo, E.; Leite, E. R.; Muccillo R. Appl. Phys. A: Mater. 2003, vol 76, 621-623. [9] Amekura, H.; Umeda, N.; Takeda, Y.; Lu J.; Kishimoto, N. Appl. Phys. Lett. 2004, vol 85, 1015-1017. [10] Amekura, H.; Umeda, N.; Sakuma, Y.; Kishimoto, N.; Buchal, Ch. Appl. Phys. Lett. 2005, vol 87, 013109. [11] Vallet, C. E.; White, C. W.; Withrow, S.P.; Budai, J. D.; Boatner, L. A.; Sorge, K. D.; Thompson, J. R.K.; Beaty, S.; Meldrum, A. J. Appl. Phys. 2002, vol 92, 6200-6204. [12] Sorge, K. D.; Thompson, J. R.; Schulthess, T. C.; Modine, F. A.; Haynes, T. E.; Honda, S.; Meldrum, A.; Budai, J. D.; White, C. W.; Boatner, L. A. IEEE Trans. Mag. 2001, vol 37, 2197-2199. [13] Ila, D.; Williams, E. K.; Zimm erman, R. L.; Poker D. B.; Hensley, D. K. Nucl. Instrum. Meth. B, 2000, vol 166-167, 845-850. [14] Xiang, X.; Zu, X.T.; Zhu, S.; Wang, L. M. Appl. Phys. Lett. 2004, vol 84, 52-54. [15] Xiang, X.; Zu, X. T.; Bao, J. W.; Zhu S.; Wang, L. M. J. Appl. Phys. 2005, vol 98, 073524. [16] Xiang, X.; Zu, X. T.; Zhu S.; Zhang, C. F.; Wang, L. M. Nucl. Instrum. Meth. B, 2006, vol 250, 192–195. [17] Xiang, X.; Zu, X. T.; Zhu S.; Wei, Q. M.; Zhang, C. F.; Sun K.; Wang, L. M. Nanotech, 2006, vol 17, 2636-2640. [18] Zhu, S.; Wang, L. M.; Zu X. T.; Xiang, X. Appl. Phys. Lett. 2006, vol 88, 043107. [19] Zhu, S.; Xiang, X.; Z u, X.T.; Wang, L.M.; Nucl. Instrum. Meth. B, 2006, vol 242, 114– 117. [20] Xiang, X.; Zu, X. T.; Zhu S.; Zhang, C. F.; Wang, L. M. Nucl. Instrum. Meth. B, 2006, vol 250, 229–232. [21] Xiang, X.; Zu, X. T.; Zhu S.; Wang, L. M. Physica B, 2005, vol 368, 88-93. [22] Zhu, S.; Sun, K.; Zhang, Q. Y.; Zu, X. T.; Wang, L. M.; Ewing, R. C. J. Appl. Phys. 2003, vol 94, 5648-5651.

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[49] Villars P.; Calvert, L. D. Peason’s Handbook of Crystallographic Data for Intermetallic Phases, ASM, Materials Park, OH 1991, Vol. 2. [50] Toneguzzo, P.; Acher, O.; Viau, G.; Pierrard, A.; Fievet, F.; Rosenman, I. IEEE Trans. Magn. 1999, vol 35, 3469-3471. [51] de Julián C.; Sangregorio, C.; Mattei, G.; Battaglin, G.; Cattaruzza, E.; Gonella, F.; Lo Russo S.; D'Orazio, F.; Lucari, F.; De, G.; Gatteschi, D.; Mazzoldi, P. J. Magn. Magn. Mater. 2001, vol 226–230, 1912-1914. [52] Morales, M. P.; Munoz-Aguado, M. J.; Garcia-Palacios, J. L.; Lazaro, F. J.; Serna, C. J. J. Magn. Magn. Mater. 1998, vol 183, 232-240. [53] Wijn H. P. J. Magnetic Properties of Metals, Springer: Berlin, 1991, pp 22.

In: Nanotechnology... ISBN 978-1-60692-162-3 c 2010 Nova Science Publishers, Inc. Editors: C.J. Dixon and O.W. Curtines, pp. 559-602

Chapter 18

S TRUCTURAL, DYNAMICAL AND O PTICAL P ROPERTIES OF S ELF - ASSEMBLED P ORPHYRINS AT THE M ESOSCOPIC S CALE Valentina Villari 1,∗, Norberto Micali1 and Luigi Monsu´ Scolaro2 1 CNR-Istituto per i Processi Chimico-Fisici, S.ta Sperone C.da Papardo, I-98158, Messina, Italy 2 Dipartimento di Chimica Inorganica, Chimica Analitica e Chimica Fisica, Universit´a di Messina Salita Sperone 31, I-98166 Vill. S. Agata, Messina, Italy

Abstract Organized self-assembly of molecules, driven by noncovalent intermolecular interactions, is the most versatile tool for accessing new materials with desired optical and electronic properties. Porphyrins are particularly attractive species to incorporate into supramolecular assemblies because their rich photochemistry may impart functionality, provide insight into the mechanisms of biological processes such as photosynthesis, serve as probes into the features of self-assembled structures and as models for molecular organization and energy/electron transfer processes. The close molecular packing in a self-assembled porphyrin aggregate leads to different electronic coupling and delocalization of the excitation energy, which can be exploited for applications in non-linear optical devices, photoelectric cells, recording devices. The possibility to control and tune either shape and size of the porphyrin clusters opens the way for their use as potential nanodevices. This review aims to collect some recent developments in the field of porphyrin self-assembly and to frame all the reported topics into the current theories.

1.

Introduction

Porphyrins constitute a wide class of natural and synthetic molecules whose photophysical properties can be modulated by changing the peripheral substituent groups and/or by ∗

E-mail address: [email protected]; [email protected]. (Corresponding author)

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inserting metal ions in the central core of the macrocycle.[1, 2] In nature, for instance, porphyrins are present under different forms and are responsible of many biochemical processes in animal and vegetable kingdom: specific examples are furnished by hemoglobin or by chlorophyll, as well as by many other biological molecules whose building blocks are porphyrin residuals, like cytochromes and hemocyanines. Along with natural porphyrins many synthetic porphyrins are exploited in biological and medical research for the study of nucleic acid conformation, of intercalation phenomena[3, 4, 5, 1, 6] and as drugs in anticancer photodynamic therapy.[7, 8] Moreover, the stability of these systems makes them especially interesting for photoionization processes, energy/electron transfer, photocatalysis and nonlinear optical properties.[9, 10, 11, 12] Numerous studies, for instance, have been carried out to investigate the DNA structure, by exploiting the ability of some porphyrin derivatives of intercalating into DNA of appropriate composition, while other porphyrins, depending on the nature of peripheral substituents or inserted metals, are limited to external, groove binding.[3, 4, 5, 1] The interaction of porphyrin-based moieties with different chemical species makes them useful for molecular recognition processes in the sensor field.[13, 14, 15] One important, promising, and newly developed practical application is the determination of the absolute configuration of various chiral compounds, even of biological importance like amino acids, and of the enantiomeric excess.[16, 17] In addition, the chemo-responsive behaviour of metalloporphyrins provides a way of reporting the presence of odors by changes in color; two-dimensional display of metalloporphyrins, as an example, was employed as sensor for the visual identification of a wide range of olfactants and solvent vapors.[18, 19] One of the most challenging aspect for specific application in materials science, condensed matter science, engineering, farmaceutics (drug delivery) as well as in nano-science and nanotechnology, consists in using porphyrins and their derivatives as building blocks for designing and accessing to supramolecular systems, by means of controlled self-assembly. From a general point of view, self-assembly is the autonomous organization of components into patterns or structures; it involves components from the microscopic to the macroscopic scale and many different kinds of non-covalent interactions, like van der Waals, electrostatic, hydrophobic interactions, hydrogen and coordination bonds. Thanks to the wide range of interactions when using components larger than molecules, often it is possible to adjust interactions themselves over wide ranges of strength and selectivity. Nonmolecular systems are, thus, in many aspects more versatile in their design than molecular systems. It is often easier to build-up nonmolecular components than it is to synthesize molecules, and easier to observe the process and products of self-assembly using conventional experimental techniques (like, for instance, scattering). Depending on their geometrical disposition in the assembly, porphyrins can form differerent kinds of aggregates called H- (face-to-face interactions) and J-type (side-by-side interactions). J aggregates have attracted a great deal of interest for their nonlinear optical properties originating from the close molecular packing: the stacking interactions together with electrostatic and hydrogen bonding interactions, in fact, leads to electronic coupling and delocalization of the excitation energy. Among different kinds of porphyrins the water soluble moieties are very interesting because their self-aggregation can be conveniently controlled by screening the charge repulsion through the ionic strength and pH and by varying concentration.[20, 21] More precisely,

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acting on the intermolecular interaction potential leads to structures which are different not only at a mesoscopic scale but also locally. Furthermore, porphyrin self-assembly on the DNA surface[22, 23, 24, 25, 26] or engineered viral particles[9] has been reported as a convenient method to control the size and extent of supramolecular porphyrin assemblies. Resonant energy transfer to the interacting porphyrins in these self-assembled structures has been also observed, occurrence which is of crucial importance in the design of new efficient artificial antenna systems. The local structural organization of self-assembled porphyrin aggregates can be conveniently studied by exploiting the features of the supramolecular chirality, which can arise from intrinsically chiral assemblies[27, 28, 29, 30, 31, 32] or from aggregation onto chiral templates.[34, 35, 36, 37] Selection of the chirality of a supramolecular structure, in the absence of any templating agents, was carried out by means of macroscopic chiral fields (i.e. vortex motion) during the aggregation process.[32, 33] Such an occurrence suggested to speculate about the role played by the sign of vorticity in relation to the origin of biological chirality. Also intriguing is the possibility to induce chirality on porphyrin aggregates which self-replicate in solution, that is retain the memory of their imprinted chirality.[37] These systems are useful in understanding the transfer of information in biologically relevant aggregation processes. The huge literature reporting the mesoscopic self-assembly of porphrins suggests that the shape and size of the mesoscopic (and also local) structure, as well as the aggregation kinetics and physico-chemical properties, of porphyrin aggregates in solution are not an intrinsic property of porphyrin itself. Rather they depend on the proper choice of the thermodynamic parameters of the solution and templating agent, which can be easily tuned and controlled. For these reasons porphyrin assemblies are extremely good candidates for applications in nanodevices, photoelectric cells, recording devices and as model for light harvesting in antenna systems. In this review article some recent developments in the field of porphyrin self-assembly is presented within the following topics: • theoretical and experimental aspects related to the characterization of structural and dynamical properties of the aggregating species; • kinetic mechanisms of the aggregation process and the consequent structure of the final aggregate; • dependence of the mesoscopic structure and geometry of the aggregates on the thermodynamic parameters of the solution (like porphyrin concentration, ionic strength and pH); • analogy between nonlinear optical properties of porphyrin aggregates and those of metal composites: delocalization of the excitons in submicrometric zones (”hot zones”) generating Raman and Rayleigh scattering enhancement; • supramolecular chirality induction by a templating agent or by an external field and dependence of the symmetry factor on aggregate’s size. All these subjects will be discussed and supported by numerous examples and experimental techniques (Static and Dynamic Light Scattering, Raman Scattering, Uv-Vis, Circular

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Differential Extinction and Circular Intensity Differential Scattering). Attention will be devoted to the experimental aspects concerning the mesoscopic systems like for instance the influence of the scattering on the absorption and circular dichroism measurements.

2.

Exciton Delocalization

In a porphyrin assembly the strong interactions give rise to the coupling between excited states (excitons) of the contituent molecules. The optical properties of these structures are described by the Frenkel exciton model, according to which the extent of the exciton wave functions is determined by the competition between intermolecular transfer interactions and (static) disorder. According to a purely exciton model, in fact, the energy and the broadening of the band for an aggregate of N two-level molecules can be obtained by the excitation Hamiltonian: H=

N X

(Emon + Em)b+ m bm +

m=1

N −1 X

+ J(b+ m bm+1 + bm+1 bm )

(1)

m=1

Here, b+ m and bm are the Pauli creation and annihilation operators for de-excitation and excitation on the site m, J is the nearest-neighbour excitation transfer interaction, Emon = hcνmon is the molecular (monomer) two level excitation energy and Em represents the energy offsets introduced by the static disorder. For perfectly ordered aggregates (distribution width of Em being zero), the exciton wave functions are delocalized along all the length of the aggregate. The one-exciton states, |ki (the ground state being |0i), which determine the absorption spectrum, and their k-th exciton energy are given by: |ki =

N X

sin



m=1



πkm b+ |0i N +1 m

νk = νmon + 2Jcos



πk N+1



(2) (3)

If molecules are arranged in a J-structure the exciton coupling constant J is negative, which leads to red-shifts in the spectrum relative to the monomer. By contrast, J is positive for H-type aggregates, with resultant blue-shifts in the spectrum. In disordered aggregates the exciton states are mixed due to the broken translational simmetry, and localize on a part of the aggregate; these states, which can be visible as separate lines when the disorder is small, merge into one broader line when the disorder increases. The width (half-width-at-half-maximum) of this absorption line, Γ, is directly related to the finite number of molecules over which delocalization occurs:[39, 38] Ndel =

2.1.

q

3π 2 |J| /Γ − 1

(4)

Resonance Light Scattering Effects

Exciton coupling and electronic communication among the molecules in an aggregate (especially in a J-type aggregate) can cause extraordinary optical effects: one example is resonance light scattering (RLS), i.e., an increase in the intensity of scattered light at the

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563

Figure 1. Experimental extinction (plot a) and scattering (plot b) spectra of an aggregated H2 T P P S4 solution ([H2 T P P S4 ]=3 µM , pH=1).

wavelength where the aggregate has an electronic absorption transition. For aggregates of sufficient size, the enhanced scattering overwhelms the absorption and a peak in the scattering spectrum appears. Besides in solutions of porphyrin aggregates[40], the phenomenon is observed also in a variety of small chromophores in solution[41, 42], chromophore-protein complexes[43, 44], chromophore-nucleic acid complexes[45, 46] and chlorophyll a aggregates[47]. Figure 1 displays, as examples, the absorption and scattering spectra of an aggregated porphyrin solution (H2 T P P S4 of figure 2). The isolated porphyrin has a Soret band at the frequency νmon , corresponding to 1/νmon = λmon = 434nm, whose width is Γmon = 875cm−1, and the J-aggregate a peak at λJ = 490nm (J = −1275cm−1 ), with Γ ≈ 100cm−1. According to the exciton theory the resonant band width leads to a delocalization over few tens of porphyrins. Even though the exciton theory allows for a calculation of the delocalization length from the spectral width of the absorption band of the aggregate, for large aggregates the scattering component affects the absorption spectra significantly. As a consequence the evaluation of the band width can be misleading. The distinction between scattering and absorption can be conveniently done by measuring independently the extinction and RLS spectra on the same sample and by performing a linear regression analysis of the data in the region close to the resonance.[48, 49, 50] By assuming that only a single excited state of the monomer contributes to the polarizability around the lower energy absorption frequency, α(ν), the quantum mechanical expressions

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Valentina Villari, Norberto Micali and Luigi Mons´u Scolaro

Figure 2. Scheme of 5,10,15,20-tetrakis(4-sulfonatophenyl)porphyrin ( H2 T P P S4 ) structure. for the absorption and scattering cross-section are: Cabs = 2πνIm(α) h i Csca = 8/3π 3ν 4 (Re(α))2 + (Im(α))2

(5)

Re(α) and Im(α) being the real and imaginary parts of the polarizability for a system where the dipole moment vector has one component[48, 51, 52]. According to this theory, the scattering cross-section Csca depends on the number N of interacting chromophores and the calculation for a J-aggregate can be performed according to the point dipole approximation as described in the literature.[51] Figures 3a) and b) describe the calculated dependence of the absorption and scattering cross-section, respectively, on the number of interacting chromophores. This approach was particularly useful for obtaining the size and the structure of a J-aggregate inside the water pool of a microemulsion on increasing the size of the inner core[48], as shown in the next subsection.

2.2.

Geometrical Arrangement of Porphyrins in a J-aggregate

Porphyrin molecules are optically isotropic, but when they aggregate in a Jarrangement, the large interaction between porphyrins (coherence length of the exciton coupling) makes the free displacements of electrons under incident radiation asymmetric, giving rise to optical anisotropy[21, 48, 51] and depolarization ratio ρdep = IV H /IV V 6= 0 (IV V and IV H being the polarized and depolarized scattered intensity, respectively). Unlike the scattering cross-section, the depolarization ratio is not dependent on the size of the aggregate, but it is related only to the principal values of the polarizability tensor at the resonance wavelength[51, 52]: 

ρdep =

h



3 α// − α⊥

45 1/3 α// + 2α⊥

i2

2 

+ 4 α// − α⊥

2

(6)

being α// and α⊥ the parallel and ortogonal component of the polarizability, respectively, for an axially-symmetric molecule.

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565

Figure 3. Dependence of the absorption (upper plot) and scattering (lower plot) crosssection on the number of constituting monomers (curve a: N=1, curve b: N=2, curve c: N=5, curve d: N=10, curve e: N=20), calculated by using the RLS theory[51] for a J-type aggregate of H2 T P P S4 .

In practice, the depolarization ratio depends on the slip angle φ between adjacent porphyrin planes (see inset of figure 4) and can give information on the geometry of the excited state of an aggregated species. By assuming a parallel arrangement of the transition moments of the exciton-coupled chromophores, in the case of J-aggregates confined in a water pool of a microemulsion the slip angle decreases on increasing the droplet size, as reported in Figure 4. On considering that aggregates grow in the water pool and the dimensions are limited by the confined environment, it is possible to combine the calculated angles with simple geometrical analysis and estimate the aggregation number, Naggr , and the length of the porphyrin clusters in the inner water compartments. Assuming the model recently proposed in the literature[53], in which the porphyrins form monodimensional arrays, we can calculate the coherence length as L = (Ndel + 1)R (where R is the radius of a single molecule). The ratio r = (Ndel + 1)/(Naggr + 1) indicates the quality of the exciton delocalization, being close to unity for an ideal J-aggregate. The maximum value for this parameter reported so far has been 0.25, for J-aggregates obtained in acidic aqueous solutions, using ammonium chloride as the nucleating agent.[53] For J-aggregates in a microemulsion the r values obtained range between about 1 and 0.35 for the smallest core and the largest one, respectively, suggesting a high level of coherence for these aggregates.[48]

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Valentina Villari, Norberto Micali and Luigi Mons´u Scolaro

Figure 4. Slip angle between porphyrins as a function of the droplet radius of the water/AOT/decane microemulsion (experimental conditions: volume fraction φ = 0.05, [H2 O/AOT ]=5 ÷ 65), [acetate buffer]=25mM, pH=2.7, TPPS stock solution at 80µM ). The inset reports the depolarization ratio dependence on the slip angle, calculated from the RLS theory for a TPPS4 solution with N=50.

3.

Kinetics of Self-assembly

Aggregation of porphyrins (particularly in water solutions) can be induced acting on many parameters, so that the mesoscopic and local arrangement of the final structures vary significantly depending on the induced kinetic mechanism and rate. The rate of aggregation of particles depends on frequency of collision and probability of sticking after collision. These factors are related to the solution parameters such as temperature, pH or other chemical properties that affect particle correlations, and on the presence of an external field. In 1917 Marian Smoluchowski introduced his kinetic expression for the time evolution of particle aggregation; it relates the concentration changes of the species i and j to form species i+ j due to binary collisions.[54] This equation is based on a mean field theory of identical particles (namely it assumes that the probability of two particles meeting is simply proportional to the product of their concentration) and includes a kernel, which depends on the physical properties of the system like shape and size of particles, composition, and interaction potential. The kernel can take many forms, depending on collision interactions among the particles, and in recent years many different forms of the kinetic profile have been obtained analytically and numerically for different systems[55, 56]. For a finite system the Smoluchowski equation takes the form of a system of coupled ordinary differential equations: K(i,j)

Pi + Pj → Pi+j P∞ dck 1P j=1 K(k, j)cj i+j=k ci K(i, j)cj − ck dt = 2

(7)

where Pj denotes an aggregate containing j monomers, K(i, j) is the reaction kernel of creation and destruction of the aggregate and ck is the concentration of the species k. By

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567

considering dilute solutions, in which only binary collisions are relevant, Kernels are configurational and orientational averages of the reaction rates between two colliding species. Smolukowski equation is analytically solvable only for few kernels like the constant kernel K(i, j) = K(1, 1) (i.e. same probability of collision for all clusters and monomers), the sum kernel K(i, j) ≈ i + j, the product kernel K(i, j) ≈ ij and their linear combinations. Another particular case was suggested by Smoluchowski, who obtained an explicit form for the Kernel K(i, j) ∝ 4π(Ri + Rj )(Di + Dj ), by solving the diffusion equation, Ri and Di being the radius and the diffusion coefficient of the aggregate constituted by i monomers, respectively. No general solution exists and for real systems the reaction kernel depends on many parameters, making it difficult to know its form. However, in the case of kinetic mechanism of aggregation driving to clusters with specific properties, like scaling, self-similarity and universality, some considerations and simplifications can be done. For self-similar structures, for example, namely those structures which are space scaleinvariant (like fractals for which Ri = Rmon i1/Df , with Df defined as fractal dimension), the homogeneity relation for the kernel holds:[57, 58] K(ai, aj) = aλ K(i, j)

(8)

with the condition that the kernel for the collision between the species i and j is K(i, j) = iµ j ν , with j >> i and ν ≤ 1. The quantity λ = µ + ν ≤ 2 describes the tendency of a large cluster to join with another large cluster, so determining the overall rate of aggregation; µ, on the other hand, governs the aggregation rate between big clusters and small clusters. The occurrence λ > 1 describes a gelation process which gives rise to an infinite-size cluster in finite time. In this frame, for different values of the quantities λ and ν, different fractal structure of the aggregate is obtained. Among the numerous collision mechanisms which can be considered, particular attention has to be devoted to the diffusion driven collisions. If the sticking probability between smaller and bigger clusters is predominant ( λ = 0 or, more generally λ < ν) the aggregation process gives rise to monodisperse Diffusion Limited Aggregates (DLA) with Df = 2.5, whereas higher sticking probability between bigger clusters ( µ > 0) generates polydisperse Diffusion Limited Cluster-Cluster Aggregates (DLCCA) characterized by Df = 1.75. In the case of the Smoluchowski kernel, for example, the self-preserving shape implies: K(i, j) ≈ (i1/Df + j 1/Df )(i−1/Df + j −1/Df ).

(9)

where the Einstein-Stokes relation D = kB T /(6πηR) was used (however, a more general dependence Di ∝ i−φ can be also considered[59]). The kernel used by Smoluchowski represents an example of DLA mechanism in which λ = 0. From the mean field and Smoluchowski treatment it is clear that in colloidal systems the kinetics of growth and the relative mechanism drive the morphology of the resulting aggregate. Examples of fractal structures built through different aggregation mechanisms in a lattice can be obtained by numerical calculation and molecular dynamics, giving to each lattice site a probability p to be occupied. In a real system p is represented by the probability that,

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Valentina Villari, Norberto Micali and Luigi Mons´u Scolaro

Figure 5. Numerical calculation for fractal structure originating from DLA (at left) and percolation (at right) mechanism. under collision, two species stick together. In figure 5a) the fractal structures originating by the diffusion-limited aggregation mechanisms is sketched, as an example. A quantitative description of the aggregation process can be done by considering that the distribution of fractal clusters with different mass (or size) is an homogeneous function evolving in a self-preserving form[60], independently of the initial distribution:[61] ¯ ci (t) = i−2 f (i/S(t))

(10)

¯ ∝ tz is the where ci is the concentration of clusters constituted by i monomers and S(t) ¯ average aggregation number of clusters (or in terms of mean cluster radius R(t) ∝ tz/Df ), with z = (1 − λ)−1 . For the Smolukowski kernel given in eq. 9 it is z = 1 which corresponds to the result otained for a DLA growth mechanism. Because it is easy from an experimental point of view to monitor the time dependence of the monomer concentration (an example is reported in the following subsection), let us consider the kernel for the reaction involving only collisions between clusters and monomers: K(1, j) = j ν . The Smoluchowski equation 7 becomes: ∞ X dc1 ¯ ν−1 j ν cj ∝ S(t) = −c1 dt j=1

(11)

the last proportionality being obtained by using eq. 10 for cj (t). By integrating previous equation, the time dependence of the monomer concentration results in a stretched exponential: h i (12) c1 (t) = Aexp −t−(λ−ν)/(1−λ) According to the mean-field theory[54] and to Molecular Dynamics simulations[57, 62], a reaction-limited (RLA) early stage of the fractal aggregation (with Df = 2.1) is expected to precede the diffusion-limited regime described above. In this early stage the clustering process is hindered by short-range high repulsive barrier, so that many collisions are required before aggregation occurs (low sticking probability). For reaction limited aggregation the

Self-assembled Porphyrins at the Mesoscopic Scale

569

dynamic scaling law of the mean cluster radius is described better by an exponential in¯ ∝ exp(bt), where b depends on the experimental conditions. crease S(t) Due to its transient regime, it is difficult to observe the RLA early stage experimentally, but in the subsection 4.2 and 6.2 it will be shown how to obtain the required information in the case of porphyrin solutions. Another mechanism originating fractal structures is percolation, a model which represents a powerful tool for the study of many physical processes like second order phase transitions or gelation in complex fluids, as well as of porous materials, biological structures, polymers, just to cite some examples. This model assigns a critical threshold concentration (called site percolation probability) at which the largest cluster includes almost all the molecules and spatially extends to all the space[63], as sketched in Figure 5b).

3.1.

Experimental Observation of Porphyrin Fractal Structures

The fractal structure of colloidal aggregates can be studied experimentally by means of Elastic Light Scattering experiment, which measures the zero-time autocorrelation function of the scattered field, ES :[65] G1(Q, 0) = hES∗ (Q, 0)ES (Q, 0)i = I(Q) = KMw cP (Q)S(Q)

(13)

in which P(Q) and S(Q) are the normalized form factor and the structure factor, respectively, Mw the weight average molecular weight of the particle, c the mass concentration, Q = (4πn/λ0)sin(θ/2) and K = [4π 2n2 /(NA λ40)](dn/dc)2 the optical constant (λ0 being the incident wavelength, θ the scattering angle, NA the Avogadro number, n the refractive index and dn/dc the refractive index increment)[64]. The factorization in the product between P(Q) and S(Q) is performed under the hypothesis of independence of intermolecular and intramolecular averages. The form factor, P(Q), represents the intra-particle interference and is related to the spatial Fourier transform of the particle mass distribution, w(r), all over the volume of the particle: R 2 (iQ · r) dr V w(r)exp R P (Q) = V w(r)dr

(14)

For particles smaller than the wavelength (tipically smaller than λ0/20) P (Q) ≈ 1. The structure factor, S(Q), is defined as the Fourier transform of g(r1, r2) − 1, where g(r1, r2) = hρ(r1)ρ(r2)i is the number density correlation function. In the case of selfsimilar systems this function is homogeneous, that is: 0

hρ(σr1)ρ(σr2)i = σ −A hρ(r1)ρ(r2)i

(15)

where ρ(r1) is the number density of monomers at the position r1. For an homogeneous and isotropic system the number density correlation function depends only on the distance between two particles g(r1, r2) = g(|r2 − r1| , 0), simply indicated as g(r). Therefore, the structure factor can be written as: S(Q) = 1 + ρ

Z

[g(r) − 1]exp (iQ · r) dr

(16)

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Valentina Villari, Norberto Micali and Luigi Mons´u Scolaro 0

For self-similar structures eq.15 implies that hρ(r)ρ(0)i ≈ r−A , A0 being related to the space dimension d and to the fractal dimension Df as A0 = d − Df . As already seen in the previous section, the fractal dimension is defined through a scaling law between the number of monomers constituting the cluster with radius R and the radius itself: N (R) = RDf . By using the power-dependence of the number density correlation function on r, the structure factor is easily calculated:[60] S(Q) = Q−Df (17) Previous equation is valid for describing ideal fractals, i.e. clusters extending all over the space scale; real fractal systems, however, have a finite size and at Q values low enough they must obey a Gaussian law (Guinier limit)[66]. Chen and Teixeira[67] proposed a structure factor suitable for describing a finite fractal aggregate constituted by m0 monomers, encompassing Gaussian and fractal behavior. The finite extension of the aggregate is taken into account by an exponential cut-off exp(−r/ξ), with ξ a cutoff correlation length: S(Q) = m0

sin[(Df − 1)arctan(Qξ)] (Df − 1)Qξ(1 + Q2 ξ 2)(Df −1)/2

(18)

Previous equation reduces to eq. 17 if Qξ >> 1 and ξ/Rmon > 50. The theoretical approach described above, along with the experimental results obtained by light scattering, turns out to be extremely powerful in showing that the kernel of the kinetic equation can be easily modified by means of pH, ionic strength and concentration to give rise to aggregates with different fractal morphology. In figure 6, as examples, the intensity profiles of porphyrin solutions at different pH values are reported. By considering that the smallness of porphyrin monomers makes the form factor be approximately unity, the measured intensity profile is directly related to the structure factor. The porphyrin aggregation process is driven by the interparticle potential which can be described through the well-known Derjaguin-Landau-Verwey-Overbeek potential (DLVO),[68] related to the presence of a diffuse double layer surrounding colloidal particles. For H2 T P P S44− solutions containing [HCl] < 0.01 M the protonation of the core (giving rise to the diacid form H4T P P S42− ) originates the zwitterionc form. Under these conditions the electrostatic repulsion between negative charges of the SO3− perypheral groups is still high and is responsible for the high kinetic barrier to the approach of monomers and the low probability of aggregation, so that light scattering experiments do not reveal any detectable amount of aggregates. On increasing the acid concentration (0.1M < [HCl] < 4M ) negative charges are progressively screened by H+ ions and the sticking probability increases accordingly. At [HCl] ≈ 0.1M the interaction between diacid porphyrins is still repulsive (charge 2-) with a long screening length and the weak attractive van der Waals forces ( ∝ r−6 ) are scarcely effective, leading to a small probability of interaction. Therefore, after an induction period due to the small sticking probability, the DLVO potential, whose attractive component depends on the third inverse power of the interparticle distance, favors the adhesion between clusters (µ > 0). As shown in Figure 6, the scattered intensity profile follows a power law with Df = 1.75, indicating that small clusters are able to self-interact

Self-assembled Porphyrins at the Mesoscopic Scale

571

Figure 6. Absolute scattered intensity profile of aggregated H2T P P S4 in water solution under strongly acidic conditions (at the end of kinetics). Experimental conditions: [H2 T P P S4 ]=3µM and [HCl]=0.1 M (squares), [HCl]=1 M (circles), [HCl]=2 M (triangles), λ0 = 532nm. The continuous line has slope equal to -1.75 and the dashed one slope equal to -2.5. leading to larger aggregates with a DLCCA mechanism. For [HCl] ≥ 1M the net charge of the molecules is small enough to lower the screening length and to increase the sticking probability between clusters and monomers ( µ < 0); aggregation takes place according to a DLA mechanism (Df = 2.5). At [HCl] ≈ 4M , the net charge on the monomer becomes very small and the ionic strength is large enough to almost reduce to zero the screening length. The aggregation process is driven almost entirely by attractive forces and it takes place by nucleation, without any fractal arrangement. Any further increase of the acid concentration leads to the full protonation of the sulfonate end groups and to the disruption of the aggregates. Because of the presence of the two positive charges, the H8 T P P S42+ porphyrin does not self-interact even under these extreme ionic strength conditions.

3.2.

Monitoring of the Fractal Growth

As the structural information on the formed fractal cluster can be obtained by Static Light Scattering, the size distribution and dynamic scaling properties can be studied by Quasi-elastic Light Scattering. In fact, by considering that the form factor of monomers is unity, the autocorrelation function of the scattered field contains information on the dynamics of the growing aggregates: G1 (Q, τ ) = hES∗ (Q, 0)ES (Q, τ )i = a

X

Mi2Ni Si (Q)exp(−Γi τ )

(19)

i

where a is a factor depending on the experiment geometry, Mi , Ni and Si represent the mass, number and structure factor of the i-th cluster with radius Ri and the decay rate Γi =

572

Valentina Villari, Norberto Micali and Luigi Mons´u Scolaro

Figure 7. Diffusion coefficient of a fractal aggregate of H2T P P S4 at 3µM with [HCl] = 1M as a function of the exchanged wave vector (at the end of kinetics). The dashed line indicates the Q2 dependence. The incident wavelength is λ0 = 532nm. Di Q2 + A(QR)Θi (Di and Θi being its translational and rotational diffusion, respectively, and A(QR) being related to the anisotropy of the cluster). The initial decay rate of the correlation function is given by: Γ=

1 G1 (Q, 0)

Z

Mi2Ni (t)Si (Q)(DiQ2 + A(QR)Θi)dMi

(20)

In the QR << 1 limit, the amplitude A(QR) is small and the translational term is domiP 0 0 ¯ nant, so that Γi ∝ R−1 i Ri (t)N (Ri, t) (N (Ri , t) being the i . The mean radius R(t) = size distribution), obtained from the mean decay rate of the correlation function, allows for the check of the dynamic scaling of the aggregate size.[69] More generally the relaxation rate of the correlation function can be written as Γ = DQ2 F (QR), where the function F (Q, R) takes into account aggregate internal motions. In the limiting case of QR << 1, F (QR) = 1 and Γ contains information only on the translational diffusion, as previously seen; in the opposite case, when QR >> 1, F (QR) ∝ Q3 and no diffusion coefficient can be extracted. For fractal systems, for which the condition QR >> 1 is fullfilled, the diffusion coefficient (and hence the size) are obtained only for rigid structures (absence of internal motions). Thanks to the strong inter-porphyrin interactions inside the aggregate, fractals constituted of porphyrins are rigid, as put in evidence by the dynamic light scattering measurements (see Figure 7). From a practical point of view, although dynamic light scattering is able to measure the mean radius of the growing clusters for the cheking of the dynamic scaling, in the case of porphyrin solutions the rate of the kinetic process is faster than the time required for data aquisition. Alternatively, UV-vis measurements guarantee a higher dynamic range and allow for the determination of the time-dependence of the monomer concentration and an estimation of that of the mean cluster mass. Figure 8a) reports the time evolution of the spectral contribution of monomeric t − H2 Pagg

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573

Figure 8. Integrated area of the absorption band of the monomeric t − H2Pagg , centered at 420 nm (plot a), and of the extinction band of the aggregated form, centered at 452 nm (plot b). Experimental conditions: [t − H2Pagg ]=5µM , [N aCl] = 0.1M . The continuous line in plot a) represents the fit according to eq. 21 and that in plot b) a power law with exponent z ≈ 1. porphyrins (see figure 9 for the structure) which are depleted from solution to originate clusters. The theoretical prediction of a stretched exponential (eq. 12) is valid in the late aggregation stage and it cannot describe the whole kinetic. As already pointed out a complete description of the measured monomer concentration must take into account the RLA early stage, during which the depletion of monomers follows an exponential form[70]. So that the time dependence of the monomer concentration can be represented by:[71, 72] 



c1 (t) = A0 + A1exp −(k1 t)−γ + A2 exp(−k2t)

(21)

where k1 and k2 are two observed time constants, γ = (λ − ν)/(1 − λ) and each process is weighted by a proper amplitude factor. The continuous line in Figure 8a) represents the fit

Figure 9. Scheme of the trans-bis(N-methylpyridinium-4-yl)diphenylporphyine ( t − H2 Pagg ) structure.

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Valentina Villari, Norberto Micali and Luigi Mons´u Scolaro

with eq. 21 and shows a very good agreement with the experimental behaviour. On the other hand, by following the time evolution of the increasing spectral contribution ascribed to porphyrins in the aggregated form, it appears that it is obeying (at least before the levelling off due to finite concentration constraints) the power law of the mean cluster mass, with z ∼ = 1 (see Figure 8b). This implies a value λ ≈ 0 and suggests that, after the very early stage driven by RLA, a DLA growth mechanism occurs. The dynamic scaling of the cluster mass is fullfilled likely due to the predominance of resonant scattering, which depends on the weight average molecular mass of the aggregate, on the absorption at the characteristic wavelength of the J aggregate.

4.

Tuning and Control of the Aggregate Mesoscopic Structure

A careful survey of the literature about aggregation of the anionic TPPS4 porphyrin shows that different authors found different structural arrangements for the final aggregate[73, 74, 75, 76], especially regarding structure and size. In this section it will be shown that structural changes of porphyrin aggregates can occur in aqueous solutions by simply changing porphyrin concentration and ionic strength. These parameters, acting on the interaction potential and on the aggregation kinetic rate, lead to final aggregates with different structure and size. The resonant scattering allows for collecting experimental data with a good signal to noise ratio even in the depolarized configuration and for determining the reorientational motions of the porphyrin aggregates. To this aim all the components of the polarizability tensor, in principle, must be measured, but, for sake of simplicity, let us assume a cylindrical symmetry of the polarizability tensor. For a dilute solution of monodisperse noninteracting N particles, resulting from the aggregation of small molecules (small with respect to the wavelength of radiation), the Rayleigh-Debye-Gans approximation can be used. Thus, the normalized polarized and depolarized field autocorrelation functions in the absence of correlation between translational and rotational motions are:[64, 66, 77, 78] 

GV V (Q, t) = N S(Q) α2iso +



  4 2 β exp (−6Θt) exp −DQ2 t 45

(22)

  β2 exp (−6Θt) exp −DQ2 t (23) 15 where αiso is the isotropic excess polarizability of the particle ( 1/3T r(α)) with respect to the solvent, β the anisotropy of the particle polarizability and D and Θ the translational and rotational diffusion coefficient, respectively. The latter two coefficients can be obtained experimentally from the initial decay rate of the field correlation functions, as follows:

GV H (Q, t) = N S(Q)

ΓV V (Q) = DQ2 +

6Θ 45α2iso 4β 2

(24)

+1

ΓV H (Q) = DQ2 + 6Θ

(25)

From equations 22 and 23 taken at t=0, the depolarization ratio is written as ρdep = β 2 /(15α2iso + 4/3β 2), indicating that the intensity of the depolarized scattering becomes

Self-assembled Porphyrins at the Mesoscopic Scale

575

relatively more important the larger the optical anisotropy of the scattering object. The hypothesis of independence of rotational and translational motions leading to eqs. from 22 to 25 remains valid for dilute solutions until the coupling parameter Q2 ∆D/Θ < 5 (where ∆D = D// − D⊥ is the anisotropy in the translational diffusion motion).

4.1.

From Fractal to Rod-Like Structures

As it was already shown, in dilute solution of water soluble porphyrin aggregation is triggered by lowering pH enough to shield the charged sulfonate groups and allow a closer approach of molecules to occur; the same effect is obtained by increasing ionic strength. In both cases extended fractal structures are formed.[20, 21] In the presence of salt the autocorrelation function relaxation rate in polarized and depolarized configuration does not indicate any difference, as shown in Figure 10a) and b). This finding indicates that rotational motions do not contribute to the width of the quasielastic optical spectrum and eqs from 22 to 25 are not distinguishable one from another. Indeed, although aggregates are optically anisotropic, the rotational diffusion of the whole aggregate is too slow to be detected. Moreover, the Q2-dependence of the relaxation rate proves that, despite their large size (RH ≈ 0.85µm), the strong intermolecular interaction gives rise to internally rigid structures. The intensity profile shown in Figure 10c) clearly indicates a fractal arrangement of porphyrins. The Q-range investigated by a combination of data at wide and small angles is wide enough to display the bending at small Q values, representing the effect of the cutoff of the density correlation function described in eq. 18. From the fit it results that this cutoff is ξ ≈ 0.7µm and the fractal dimension is Df ≈ 2.2 (consistent with that found for an analogous system[79]). It has been shown recently[80] that the fractal dimension of a DLCCA and an RLCA aggregate is an increasing function of the monomer aspect ratio, so that aggregation of rod-shaped building blocks leads to a loss of distinction between the two kinetic mechanisms of growth as the axial asymmetry increases. By considering that the building blocks of the fractal are the strongly exciton coupled porphyrins, their nonspherical shape can be responsible of the underestimated value of the fractal dimension deriving from a DLA growth mechanism. √ ∼ The gyration radius of the aggregates, Rg = 3ξ = 1.2µm, along with the value of the hydrodynamic radius, allows for an estimation of the fractal porosity through their ratio RH /Rg ∼ = 0.7; this value is in good agreement with that expected for fractals with an exponential cutoff in the density correlation function.[81, 82] When salt is not added to the solution, aggregation of T P P S4 is fostered at mild acidic conditions by increasing porphyrin concentration[21, 75]; under this condition porphyrins aggregate in rod-like mesoscopic structures. In the concentration range from 40 up to 500 µM , polarized and depolarized autocorrelation functions have significatively different relaxation rates one from another, as shown with some examples in Figure 11. From eqs.24 and 25 it is clear that the zero-Q value of the depolarized relaxation rate gives the rotational diffusion coefficient, whereas the slope of the Q2-dependence of both polarized and depolarized relaxation rate furnishes the translational diffusion coefficient. In order to adopt a reasonably corrected model for calculating the aggregate size, let us consider that the structure factor of molecules in the aggregate can be regarded as the form

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Valentina Villari, Norberto Micali and Luigi Mons´u Scolaro

Figure 10. Relaxation rate of the polarized (plot a) and depolarized (plot b) correlation function; the continuous lines represent the Q2 -dependence with the same slope. Plot c) reports the whole intensity profile along with the fit with eq. 18, typical of finite fractals. Experimental conditions: [H2 T P P S4 ]=3µM , pH=2.7, [NaCl]=1.5M, λ0 = 532nm factor P(Q) of the aggregate itself; moreover, if solution is diluted enough to neglect interactions between aggregates the structure factor can be approximated to unity. In this frame the form factor is simply proportional to the measured intensity profile through eq.13. Figure 12 displays that the scattered intensity of the concentrated porhyrin solutions (in the absence of added salt) can be well fitted by a rod form factor with length L, described by: P (Q) =

2 QL

Z 0

QL





sin(x) 2 QL dx − sin x QL 2

2

(26)

Then, for a diffusion coefficient of a rod the following Broesma’s equations hold:[83, 84] Θ= D=


3kB T 0 δ −ζ 3 πηL

  kB T  0 δ − 1/2 γ// + γ⊥ 3πηL

(27) (28)

with D=

D// + 2D⊥ 3

(29)

Self-assembled Porphyrins at the Mesoscopic Scale

577

Figure 11. Relaxation rate of the polarized (plot a) and depolarized (plot b) correlation functions at pH=2.7 for two concentration values: [ H2 T P P S4 ]=40µM (squares), λ0 = 532nm and [H2 T P P S4 ]=500µM (circles), λ0 = 780nm. The continuous lines represent the Q2 -dependence with the same slope for the same symbols and the dashed lines indicate the zero-Q value of the depolarized relaxation rate.

Figure 12. Absolute scattered intensity profile for the H2 T P P S4 solutions for different concentration values: 40µM (squares) at λ0 = 532nm, 250µM (diamonds) at λ0 = 780nm, 500µM (circles) at λ0 = 780nm, all at pH=2.7. The continuous lines are the fit according to the rod form factor (eq. 26). kB T 0 (δ − γ//) 2πηL kB T 0 (δ − γ⊥ ) D⊥ = 4πηL   2 0 δ = ln ψ

D// =

(30) (31) (32)

578

Valentina Villari, Norberto Micali and Luigi Mons´u Scolaro

ψ being the ratio between diameter and length of the rod and ζ, γ//, γ⊥ functions of the parameter δ 0 [83, 84]. By inserting the measured values of the translational and rotational diffusion coefficients in the Broesma’s equations, the solution of the non-linear systems of equations gives the geometrical parameter of the rod (L and ψ). Both static and dynamic measurements agree in indicating that the length of the rod depends on porphyrin concentration and in particular from 40µM to 250µM it is about 0.5 µm and then decreases to about 0.2 µm at 500µM ; moreover, the dynamic measurements add information on the rod diameter, ψL ∼ = 10nm. The decrease of the length with increasing concentation can be attributed to the lower steric hindrance of a large number of shorter rods with respect to a small number of longer rods. The increase of absorbance, and hence of the local heating of the sample upon irradiation, makes it more difficult to investigate higher concentration values by means of light scattering; however, X-ray measurements on the same porphyrin by Gandini et al.[75] showed that, at concentration values ten times higher and above, porphyrins are arranged in even shorter rods (about 35 nm). These authors proposed two plausible structures for the rods: one consisting in planar association (sheets) into a ringlike configuration (hollow cylinders) and the other deriving by the formation of a continuous shallow helix without disruption into separate layers. This tubular structure has been found also for other self-assembled bio-systems like bacterial antenna chlorophyll[85], guanosine nucleoside[86] and melanin particles[87]. Also porphyrin derivatives under proper conditions were proved to form fiber-like aggregates[88, 89] with only one molecule per cross section. In section 7.1 it will be shown that, in the presence of a chiral templating agent, rod-like H2 T P P S4 J-aggregates are more consistent with an helical arrangement for which the projection of the transition dipole moment is mostly perpendicular to the long axis of the rod,[90] as also observed for other J-aggregate wires.

4.2.

Effects of Initial Conditions in the Diffusion-Limited Aggregation Kinetics

The kinetic profile of a real diffusion controlled aggregation process is strongly dependent not only on the initial concentration of the reactant but also on their order of mixing (i.e. initial spatial distribution of reactants).[91] Another important factor influencing the aggregation kinetic profile is the presence of nucleation centres, for example small aggregates (seeds) already formed in an aged stock reactant, which represent a divergence in the spatial distribution of reactant itself. As far as self-assembly of porphyrins in solution are concerned, dramatic effects on kinetics (and also on the structure) are observed varying the initial spatial distribution of reactants (or using aged stock concentrated porphyrin solution). It was suggested that a protocol to mix the reagents must to be set up in order to get reproducible results, besides a careful control of the thermodynamic conditions.[92, 93] Pasternack et al., for instance, put in evidence the importance of the mixing order in studying the kinetic of supramolecular assembling of the t − H2Pagg porphyrin onto the surface of DNA[94] and the formation of J-aggregates from H2 T P P S4 [93]. Also the relative concentrations of the stock solution of reagents play an important role, as well as the lag time between injecting a reagent and mixing the solu-

Self-assembled Porphyrins at the Mesoscopic Scale

579

tion. This peculiar behaviour is also characteristic of most simple reactions as bimolecular diffusion limited reaction (A + B → C) were an ”anomalous” rate law is observed[95, 96, 97]. The anomalous effects result from the preservation of a ”memory” of the initial spatial reactant distribution, which can be minimized by thorough continuous randomization (stirring). In many diffusion limited reactions, however, the reaction progresses very quickly (at finite reactant concentrations), so that it is difficult to randomize the reactant distribution. In such a case it is difficult to fix a well defined initial time for the reaction since the mixing itself takes time. Also for this reason in realistic case of diffusion limited reaction, the inadequate mixing and the initial reactant distribution effects are dramatic. The effects of the initial distribution of reactants on the kinetics is explicity taken into account in the differential equation describing the time evolution of the density of a pair of species in the simpler case of the bimolecular diffusion limited reaction (e.g. recombination):[98] ∂ρ (33) = −Lρ = − (L0 + kr S) ρ ∂t where kr is the rate of reaction of pairs, S is a term containing information on the efficiency of reaction at different distances of the two species and their mutual orientation, and, for diffusive processes it can be written L0 = −D∇2 ρ. In the absence of reaction the density of pairs in the volume V is at the equilibrium ρ = ρeq ≈ g(r)/V and L0 ρeq = 0. By transferring the space-time dependence of the density of pairs to the probability that the R pair survive at time t[98], P (t) = drρ (r, t), equation 33 becomes: dP =− dt

Z

t

V −1 kr (τ )P (t − τ )dτ − keq I(t)

(34)

0

with I(t) a term depending on the rate constant at equilibrium, on the efficiency of reaction and on the ratio ρ(0)/ρeq (ρ(0) being the initial distribution). Equation 34, therefore, takes into account the initial distribution of reactants through the term I(t),which is equal to zero only when ρ(0) = ρeq . In the next subsections some examples of the effects of zero-time conditions on the aggregation of porphyrins are reported. 4.2.1. Effects of Mixing Order and Aging in Porphyrin Aggregation Two different mixing methods can be adopted in preparing H2 T P P S4 aqueous solutions in the presence of other reagents required for triggering aggregation: adding a stock solution of concentrated porphyrin as last reagent (porphyrin-last mixing, PL) or as first reagent at the desired concentration in water (porphyrin-first mixing, PF). Although UV-vis measurements do not distinguish between the two final solutions, being sensitive only to the local molecular arrangement, light scattering puts in evidence two significatively different mesoscopic structures. In fact, as previously seen, the kinetics of growth and the relative mechanism drive the morphology of the resulting aggregate. The intensity profiles reported in Figure 13 put in evidence that at higher pH aggregates exhibit a fractal structure when the porphyrin is added in an acidic environment (PL mixing), while they form statistical isotropic objects using the PF mixing. The structure of such

580

Valentina Villari, Norberto Micali and Luigi Mons´u Scolaro

Figure 13. Absolute scattered intensity profile for the H2 T P P S4 solutions ([H2 T P P S4 ]=3µM ) at different concentration of hydrochloric acid ([HCl]=0.1 M (plot a) and [HCl]=2 M (plot b)) for PL (squares) and PF (circles) mixing order (λ0 = 532nm). isotropic objects can be described by the Ornestein-Zernike behavior: 

I(Q) ∝ 1 + ξc2Q2

−1

(35)

ξc being the correlation length. In the PF mixing case there is no fractal arrangement in the whole investigated Q range. At lower pH, but still high enough not to break the aggregates (see section 4.1), the distinction between the structures formed with the two mixing methods is progressively lost. All these occurrences can be ascribed to the inevitable concentration gradient of the added last reagent, which causes different nucleation probability in some parts of the system. In the case of the acidic porphyrin solutions this effect becomes more important when the final concentration of acid is not high and is mainly due to the large volumetric ratio between the reagent solutions to be mixed. Another example is given by the porphyrin aggregation in the presence of a polyamine, spermine; also in this case the mixing order is a key parameter. In figure 14 the intensity profiles shows that, in the PF mixing, aggregates take a fractal structure with a fractal dimension of Df ≈ 2.5 and radius of 3µm (as obtained by the fitting with eq. 18), whereas in the PL mixing no fractal arrangement is observed (aggregate radius ≈ 0.6µm). The presence of small aggregates pre-formed in the porphyrin solution before adding the reactant and inducing aggregation can act as nucleation centers. The number and the size of these pre-existent aggregates affect, in a uncontrolled way, the features of the aggregation

Self-assembled Porphyrins at the Mesoscopic Scale

581

Figure 14. Absolute scattered intensity profile for the H2 T P P S4 solutions ([H2 T P P S4 ]=3µM , λ0 = 532nm) in the presence of 30-fold higher spermine concentration for PF (plot a) and PL (plot b) mixing order. kinetics, e.g. the induction time and rate. The induction time is the apparent initial reaction inactivity, often observed experimentally in porphyrin aggregation[94], which makes sigmoidal-shaped the kinetic profile. An example of the aging effect on induction time and rate of aggregation is displayed in figure 15.

5.

Formation of Organic Fractal Composites

The resonance properties of porphyrin J-aggregates have proved to be extremely useful for the study of their structure and dynamics, as well as for insights in the aggregation kinetic mechanisms of fractals. But porphyrin fractals reserve other interesting electromagnetic phenomena, like the scattering enhancement with scaling properties analogous to fractal metal composites. In composites (thin films, colloidal aggregates, etc...), which are formed by non-linear material embedded in a host medium (either linear or non-linear), the enhancement of the linear and non-linear optical response originates from the strongly fluctuating local fields, Ei , which exceed the applied one, E0.[99, 100, 101, 102] The local field enhancement is described by the factor G=

N X 1 |Ei|2 . N |E0 |2 i=1

(36)

N being the number of monomers constituting the cluster. By considering the composite constituted by N polarizable small (with respect to incident wavelength) monomers with short range dipole-dipole coupling, the local field, acting on the i-th monomer is a superposition of the incident and all the scattered waves:[103, 104] Eis = α−1 mon dis = E0s exp(ik0 · ri ) −

N X j6=i=1

Wsl (ri − rj )djl

(37)

582

Valentina Villari, Norberto Micali and Luigi Mons´u Scolaro

Figure 15. Kinetic profiles of t − H2 − Pagg aggregation ([t − H2 − Pagg ]=5µM with [NaCl]=0.1M) obtained by starting from a freshly prepared (within a week) stock solution (continuous curve) and from the same stock aged more than one month (dashed curve). with αmon the isolated monomer polarizability, and s and l indicate the different orientation in the space. Previous coupled-dipole equation means that the amplitude of the transition dipole moment di at the i-th monomer within the composite is related to the external field E0 through a light-induced dipolar interaction W . The condition of short range dipole-dipole coupling is not so restrictive for fractal composites because of the strong localization which makes the interaction of monomers at distances greater than the excitation wavelength negligible. The dipole-dipole interaction, expressed in eq. 37, can be synthetically written as: (Z + W )|d >= |E0 >

(38)

where Z = α−1 mon and |d > is the state of the light-induced dipolar moments. Because polarizability is, in general, complex, let us define X = −Re (1/αmon ) and δ = −Im (1/αmon ) its real and imaginary part, respectively. The parameter X has the meaning of a relative frequency detuning and δ determines the resonance width (related to the dissipation within a cluster). The optical properties are easily related to the polarizability α of the system, α = P s and l for sake of brevity) being obtained through the 1/N N i=1 αi , αi (omitting indeces P solution of equation 38 as αi = nj [(< is|n >< n|jl)/(Z + wn )]. The eigenvalues wn of W (with eigenvectors |n > traditionally called surface plasmons) contribute to the width of the resonant band. With these definitions the enhancement described by equation 36 becomes:[104] h

i

G = δ 1 + (X/δ)2 Im(α).

(39)

The ratio |X| /δ is called quality factor, because, when |X| >> δ, the enhancement factor can be very large. Dynamic simulations showed that enhancements and fluctuation of local fields in non-fractal composites are significantly lower than for fractals.[105]

Self-assembled Porphyrins at the Mesoscopic Scale

583

The enhancement in a composite concerns both non-linear (e.g. harmonic generation phenomenon and Kerr effect) and linear (e.g. scattering, absorption and fluorescence) optical response and is defined as the ratio between the optical response from monomers in a cluster and that of the same number of isolated monomers. As far as linear properties are con  (0) cerned, Rayleigh scattering enhancement factor, for instance, is GR = σs / N σs , where (0)

σs is the scattering cross-section of the composite and σs that of the single monomer. For Raman scattering, on the other hand, the enhancement factor depends also on the Raman (S) shift. By considering that the dipole moment di induced at the Stokes-shifted frequency (S) νS on an isolated monomer (with Raman polarizability χi ) interacts, besides with the local field, with other dipoles, the enhancement of the Raman scattering is:

GRS =

  P (S) 2 i di

(40)

N |χ|2 |E0|2

where |χ| is the Raman polarizability of the monomer (due to the incoherent nature of the Raman scattering hχ∗i χj i = χ2 δij ) and averages are over orientations. For very large Stokes shift, outside the aborption band of the cluster the enhancement factor is due only to the local field and GRS reduces to eq. 39. But when the Stokes shift is small also the Raman amplitudes are enhanced according to: GRS ≈

*

|Ei |4 |E0 |4

+

E

D

4

≈ |αi |4 α−1 0 .

(41)

More in general for a non-linear optical process ∝ E n the enhancement can be written as: G=

N X 1 |Ei |n ≈ |X|n δ 1−n Im(α) n N |E0| i=1

(42)

where the right-hand of the previous equation follows from eq.39 under the condition |X| >> δ. The enhancement of the optical response is expecially high in composites with fractal morphology for which the breaking of the translational invariance causes the localization of the excitations (eigenmodes) in subwavelength regions; in these ’hot zones’ absorption by monomers is much higher than by other monomers in a fractal composite. Moreover, the scale invariance gives to the optical response of these composites interesting scaling properties provided that the eigenmodes delocalization (i.e. coherence length) occurs over a distance in between the characteristic spacing between the nearest monomers and the radius of the cluster (R0 << L << R, where for fractals it has to be recalled that (R/R0)Df = N ). In a self similar composite any physical quantity obeys a scaling law in Z with some scaling index and, under the condition |X| >> δ, the following scaling law was derived:[100, 104] 

d0 −1

Im(α) ≈ R30 R30 |X|

(43)

with 0 ≤ d0 ≤ 1 the optical spectral dimension, which is analogous to the spectral dimension appearing in the scaling law for the vibrational density of states. Also the coherence

584

Valentina Villari, Norberto Micali and Luigi Mons´u Scolaro

length of dipolar excitations obeys a scaling law involving the same optical spectral dimen
(d −1)/(3−Df ) . sion L ≈ R0 R30 |X| 0 From eqs. 42 for Rayleigh scattering (n=2) and 43 it follows[103] that the Rayleigh scattering enhancement of a fractal composite obeys the scaling law GR ≈

d0 +1 N  3 R |X| 0 R30 δ

(44)

for clusters smaller than the excitation wavelength ( kR << 1, with k = 2π/λ0). For larger clusters eq. 44 becomes:[103] GR ≈

d0 +1 CR  3 R |X| . 0 R30 δ

(45)

Previous equation indicates that the enhancement comes from the resonance of the scattering by dipolar eigenmodes (i.e. from δ −1 ) and from the prefactor CR , representing the coherence due to fractality, which takes different form depending on the fractal dimension Df .[100] It is worth to note that when a fractal tends to an homogeneous system ( Df =3) the |X| range in which the enhancement occurs reduces to the frequency of isolated monomer (for metal particles it is the surface plasmon frequency). Numerical simulations of optical properties in some fractal composites showed that the scaling behaviour of the quantities discussed is obeyed.[100, 103, 106] On the other hand some experimental investigations have been carried out on clusters of metal particles.[107, 108, 109] Recent experimental works[110, 111] proved that porphyrin fractals embedded in a polyamine matrix (spermine) give rise to scattering enhancement phenomena, which find a natural framework, as organic fractal composite, in the theory described above. In the next subsections it will be shown that absorption and elastic scattering of this organic composite obey the scaling laws of fractal composites and that the enhancement of the Raman scattering is a powerful tool for the study of the very early stage of porphyrin aggregation.

5.1.

Rayleigh Enhancement in a Porphyrin Composite

In typical J aggregates the exciton dipoles are delocalized over about a few tens of molecules; under this condition there is a resonant scattering at the J-aggregate absorption band and optical properties are well described by the Frenkel theory. Local field enhancement has been obtained only in the presence of metal tip in near-field optical microscopy (e.g. ref. [112]). Unlike H2T P P S4 aqueous solution in acidic conditions, for which the width of the Jaggregate extinction band is narrow, when a ligand molecule (i.e. a polyamine) is added in excess concentration a significant broadening occurs, as displayed in Figure 16. By applying the Frenkel model, the spectral width of the J-aggregate resonant peak would correspond to Ndel < 3, which is inconsistent with the measured resonance wavelength; in fact, the expected value from the resonant light scattering theory for a small number of coupled excitons is close to Emon (for Ndel =3 it should be λJ =470 nm[93]).

Self-assembled Porphyrins at the Mesoscopic Scale

585

Figure 16. Extinction spectra of H2T P P S4 in the presence of hydrochloric acid (curve a: Experimental conditions [H2T P P S4 ]=3 µ M and [HCl]=0.1M) or of spermine (curve b: Experimental conditions [ H2 T P P S4 ]=3 µ M, [spermine]=100 muM, pH=2.7, [citrate buffer]=10 mM). Rather, the broadening extending in the long wavelength region well outside the absorption frequency of the J aggregates suggests that porphyrins can interact through dipolar coupling, giving rise to scattering enhancement. Moreover the enhancement is particularly high since, as seen at the end of section 5, the aggregates formed in the presence of polyamine possesses a fractal structure (when porphyrin is added as first reagent). The proposed model for the structure of the present system[110] can be sketched by considering the formation of an intermolecular network of porphyrins (in an edge-to-edge arrangement); in such a network the ligand induces the branching through interactions between the protonated nitrogen atoms of the polyamine and the negatively charged sulfonate end groups of the porphyrin not involved in the porphyrin-porphyrin contacts. In this heterogeneous system the ligand is directly involved in the formation of the fractal and, besides the exciton coupling inside the J aggregate, the porphyrin molecules (embedded in the polyamine matrix and not involved in the J aggregate) interact through dipolar coupling. In order to discriminate between absorption and scattering contribution in the extinction spectra, two independent measurements of extiction and absorption (see Figures 16 and 17) were performed using a homemade spectrophotometer.[110] Absorption can be related to the imaginary part of the polarizability (see eq. 5) as kIm(α) and plotted against |X| = |νmon − ν|, νmon being the frequency of the porphyrin monomer absorption band. In figure 17 data fullfill the scaling law of eq. 43 with slope -0.7. This value is in a very good agreement with the d0 = 0.3 found through numerical simulation in DLCCA metal fractal composites.[103] The enhanced scattering contribution can be calculated as the difference between the extintion and absorption spectra.[113] Figure 18 clearly shows the monotonic increase of the enhanced scattering with the wavelength for the solution in the presence of polyamine; for the acidic solution in the absence of polyamine the resonance[94] is limited in the wavelength range close to the J-aggregate absorption band.

586

Valentina Villari, Norberto Micali and Luigi Mons´u Scolaro

Figure 17. Scaling law for the absorption, A ≈ kIm(α) (see eq. 43), obeyed with d0 = 0.3, both for positive and negative values of X. In the inset the absorption spectrum of H2 T P P S4 aggregates is reported in the presence of spermine. In the inset of figure 18 the Rayleigh scattering enhancement is reported against X; it is well represented by eq. 45, valid for large fractals (kR ≈ 45). The optical spectral dimension obtained by the slope shown in the inset of figure 18 is d0 = 0.3. The explicit form used for the prefactor CR is CR = Qsc πR2 (Qsc being the scattering efficiency factor of nonabsorbing spherelike aggregates). These results suggests that, in order to explain the broadening of the absorption band, the energy shift observed experimentally, and the wavelength dependence of the scattering, the porphyrin-spermine system should be regarded as a fractal nanoparticle composite.

5.2.

Identification of the RLA Early Stage in Porphyrin Aggregation

In the H2T P P S4 -spermine composite the Raman scattering enhancement aids the identification of the early stage of the porphyrin aggregation.[111] We have shown that the diffusion-limited kinetics describe well the late stage of the aggregation leading to the final fractal structure, but it should be preceeded by a reaction limited early stage during which the seeds triggering the extended aggregation are formed. This occurrence was suggested experimentally (in section 4.2) by the presence of an additional exponential decay in the time dependence of the depleting monomer concentration. The frequency shifts in the low frequency Raman spectrum, shown in Figure 19, are very similar to those reported in the literature for the aggregated diacid H4 T P P S4 . In particular, the two bands at 242 and 316 cm−1 were previously assigned to vibrational motions of the aggregated porphyrin[114]. Because both the excitation wavelength and the low-frequency Stokes shifts belong to the scattering enhancement region, the Raman scattering is strongly enhanced. The Raman intensity increase with elapsing time is directly related to the concentration increase of porphyrins involved in the growing aggregate. As it is shown in the main plot of figure 19, the time evolution of the two characteristic bands is complete in about 60 s.

Self-assembled Porphyrins at the Mesoscopic Scale

587

Figure 18. Scattering enhancement for the aggregate formed in the presence of spermine. In the inset the scaling law (see eq. 45) is shown for positive values of X.

Figure 19. The main figure reports the time evolution of the integrated area ( ≈ [c1(0) − c1(t)]) of the low frequency Raman spectrum (from 180 to 340 cm −1 ); the continuous line is the fit according to Eq. 47 with a2 = 0.01s−2 . Figures from a) to d) show the characteristic Raman shifts of the aggregated diacid porphyrin ( H4T P P S4 ) at different elapsed time starting from mixing of reagents. Note that the Y-axis scale are the same for all the plots a) to d). The used excitation wavelength is 514 nm.

588

Valentina Villari, Norberto Micali and Luigi Mons´u Scolaro

During this time the Rayleigh scattering profile is flat and no resonance light scattering appears, due to the smallness of the aggregates (initial seeds). Afterward, the crossover toward an almost constant Raman scattered intensity and the appearance of an increasing scattering signal suggest the beginning of the typical slower kinetics (tens of hours) leading to the larger DLA aggregates. According to the von Smoluchowski theory we can identify this early stage of the aggregation with the predicted RLA mechanism. In order to make this discussion more quantitative let us use the form of Di Biaso et al.[62] for the reaction kernel to be introduced in the von Smoluchowski equation: K(i, j) =

(

ps f or i = j = 1 ps f or i, j 6= 1

(46)

where pS is the sticking probability and 0 <  < 1. The analytical solution of Eq. 9 obtained for monomers becomes: h

c1(t) ∝ 4 + 4at + (at)2

i−1

h

≈ 4 + (at)2

i−1

(47)

with at a scaled adimensional time depending on ps and on the initial monomer concentration; the last equality derives from the negligibility of the linear term, that in the present case is at least three orders of magnitude smaller than the quadratic term. This approach takes into account the low sticking probability of monomers and the crossover from RLA to DLA mechanisms. As already observed, the Raman intensity is proportional to c1(0) − c1(t), c1 (0) being the known initial porphyrin concentration. In the main plot of Figure 19 the continuous line represents the result of the fitting procedure with eq. 47. Complementarily, monomer depletion can be monitored in real time by the fluorescence quenching of the typical two-band emission of the diacid porphyrin upon aggregation (see the inset of Figure 20). In fact, although the enhancement is expected also in the fluorescence signal, it cannot be observed experimentally because porphyrin aggregates do not emit, as shown in the spectra of figure 20. The kinetic trace of fluorescence quenching is also well fitted by eq. 47, with the same values of the parameter a2 . Trying to fit data with an exponential behaviour does not give satisfactory result. The very good agreement between the two independent data sets and the theoretical approach of Di Biaso et al.[62] gives an experimental prove of the existence of the theoretically expected crossover between the two aggregation regimes. The induction time often observed in the UV-visible and light scattering kinetic profiles of porphyrin solutions[94] could be due to the nondetectable RLA initial stage. All the peculiar features of the porphyrin-spermine composite are retained when samples are dried after the DLA aggregation process is completed. Aggregates are embedded in a polyamine matrix (halo in Figure 21A) and their fractal morphology is visible in the optical image. A micro-Raman map, collected for a generic cluster and reported in Figure 21B), clearly puts in evidence the spatial distribution of the Raman intensity. The brighter zones of the image, i.e. those zones with higher Raman intensity, enclose the single inhomogeneities of the localized field which are averaged out due to the low resolution of the micro-Raman spectroscopy.

Self-assembled Porphyrins at the Mesoscopic Scale

589

Figure 20. Fluorescence bands of the diacid H4 T P P S4 in the monomeric form (continuous line) and upon aggregation (dashed line). The inset reports the time-dependence of the fluorescence intensity (≈ c1 (t)) at 700nm, with excitation at 488 nm. The continuous line is the fit obtained with eq. 47, with the same values of the parameter a2  as in figure 19.

Figure 21. A) Optical microscopy image of a sample of TPPS4-spermine composite deposited on a glass cover slide after evaporating the solvent (the bar is 10 µm). B) MicroRaman image collected as intensity at 242 cm −1 (the explored surface is 7 × 7µm2). (Reprinted figures with permission from N. Micali, V. Villari, A. Romeo, M.A. Castriciano, and L. Mons´u Scolaro, Physical Review E 76, 011404 (2007). Copyright (2007) by the American Physical Society. http://link.aps.org/abstract/PRE/v76/e011404).

590

6.

Valentina Villari, Norberto Micali and Luigi Mons´u Scolaro

Supramolecular Chirality

Chirality is the properties of optically active particles to display different optical response to the handedness of the incident circularly polarized light and is widely exploited, through the circular dichroism (CD) spectroscopy, to investigate structure and conformation of molecules, inclusion complexes and macromolecules of biological relevance. Chirality can also occur at supramolecular level (i.e. the nonsymmetric arrangement of molecules in a noncovalent assembly) because chiral assemblies can be obtained through chiral units, through interaction with chiral matrixes or be originated from macroscopic chiral field (vortex motion). In this sense, many chiral templates have been used to produce the induced supramolecular chirality. In the literature beautiful examples of noncovalent systems have been described in which the process of self-assembly is fully controlled and intriguing concepts such as chiral memory[37] and chiral amplification[115] have emerged with very promising consequences. Many studies have been carried out on porphyrins binding to both condensed and noncondensed DNA or helical polypeptides[23, 34, 40]. Through the study of the supramolecular chirality it is possible to investigate the biopolymer structure, as well as to identify and control the formation of long-range, organized porphyrin assemblies on the biopolymer, which serves as a template. Other species of chiral templates was proved to give the chiral ’imprinting’ in the very first step of the formation of chiral aggregates of porphyrins, but, once formed, these aggregates seem to retain ’memory’ of the chiral arrangement despite the removal of the template[36, 37]. It was also proved that the stored chiral information can be removed with increasing temperature or just with time, so that it is possible to accurately design supramolecular systems for which a complete cycle of imprinting, storing, releasing, and restoring of the memorized chirality can be performed[116]. Symmetry breaking and evidence of the role of asymmetric forces in the process of chiral selection were also obtained by promoting aggregation under gentle rotary evaporation of dilute porphyrin solutions.[32, 33] Much interest is recently devoted to the supramolecular chirality and the recent experimental results open the way towards the design of chiral soft materials that could be used for symmetry amplification, for chemical sensing and enantioselective catalysis. In the next subsections, by using self-assembled porphyrins, a brief outline of the differential scattering in chiral supramolecular assemblies will be reported; moreover, a relation between mesoscopic structure and supramolecular chirality will be shown.

6.1.

Scaling of the Asymmetry Factor in Porphyrin J-aggregates

Beyond the investigation of the different structures of porphyrin aggregates obtained under different experimental conditions, it is important to know if correlation there exists between the fractal-to-rod change of the mesoscopic arrangement and changes in the local porphyrin arrangement. The supramolecular chirality induced by a chiral acidic template during the aggregation process can be exploited for this aim, since the aggregate mesoscopic final structure is not affected by the chiral acid used. The induced circular dichroism (ICD) bands are displayed in Figure 22 for both fractal and

Self-assembled Porphyrins at the Mesoscopic Scale

591

Figure 22. Circular dichroism spectra of fractal (dashed line) and rod-like (continuous line) H2T P P S4 aggregates in solution (Experimental conditions: [ H2T P P S4 ]=3µM , [NaCl]=1.5M, [L-tartaric acid]=25mM, pH=2.7 for fractals and [H2 T P P S4 ]=40µM , [Ltartaric acid]=25mM, pH=2.7 for rods.

rod-like aggregates and appear in correspondence with the monomer and J-aggregate absorption bands (see also Figure 1). As it will be better explained in the next subsection, the strong scattering contribution causes a broadening of the ICD bands. The occurrence that linear dichroism can affect data interpretation can be excluded due to the exactly specular behaviour obtained in the presence of the two opposite enantiomers of the templating agent. The spectra of figure 22 have opposite sign, indicating the unexpected opposite supramolecular chirality for the fractal and the rod-like aggregates. Such a result can be rationalized by considering that the presence of added salt in solution has manifold effects: screening the high electrostatic repulsion between porphyrins favouring their approach, increasing consequently the kinetic rate driving towards the fractal morphology, allowing an unconstrained local arrangement. In the absence of salt, on the other hand, aggregation is forced by increasing concentration but the repulsive barrier between monomers is still high. This implies that the local arrangement is strongly dependent on those configurations which minimize the repulsion. Analogous considerations can be done for porphyrin aggregates formed in the aqueous interior of microemulsions; despite the low concentration of the microemulsion droplets in the continuous medium, the confinement of porphyrins in the water pool simulates a high concentration enviroment.[48] The sign of the circular dichroism signal obtained for the confined rods is the same as that of the rod-like aggregates in bulk and the strength increases with increasing the rod length. This occurrence indicates an effect of chirality amplification due to the long-range chiral order of the strongly exciton-coupled porphyrins (high coherence length) It is worthnoting that the choice of the circular dichroism band of the monomer (instead of that of the aggregate) was done in order to avoid the contribution of the resonance scattering. As it will be shown in the following, the contribution from scattering can affect

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Valentina Villari, Norberto Micali and Luigi Mons´u Scolaro

Figure 23. Circular dichroism spectra of an aggregated t − Cu − P agg solution obtained by placing the sample near (continuous curve) and far (dashed curve) from the detector. Experimental conditions: [t − Cu − P agg]=5µM , [polyGlutammate]=500µM , [NaCl]=0.15M, [buffer]=2 mM, pH=5.5. significantly the shape and the intensity of the circular dichroism spectrum.

6.2.

An Outline of Differential Scattering in Porphyrin Aggregates

The circular dichroism spectrum of a chiral particle is essentially dominated by differential absorption when particle size is smaller than λ0/20, but in the case of large aggregates the differential scattering, namely the ability of long-range chiral macrodomains to scatter preferentially one or the other circular polarizations of light, can affect significantly the shape of CD bands. In particular, contributions in the CD spectrum at wavelengths outside the absorption bands appear[117, 118, 119] and the magnitude of the signals inside the absorption band can be up to 1000 times larger than those observed for small systems.[120] These effects, originally referred to as ’anomalous behaviour’, allows for obtaining relevant information on the structure of a chiral system. In figure 23 the reported CD spectra have been obtained by placing the sample near and far from the detector; with these two different experimental geometries it is possible to collect transmitted light under different angular acceptance, putting in evidence the contribution of differential scattering (see eq. 48 reported in the following). The dependence of the measured spectrum on the sample position clearly shows how the differential scattering affects significantly both shape and intensity of CD bands in an aggregated porphyrin solution. A thorough quantum mechanical treatment of the differential scattering was done by some authors[121, 122], but in the following the main points of the classical approach by Bustamante et al.[123] are recalled. The signal measured in a standard CD spectrophotometer along the transmission beam contains both the transmitted and the zero-degree scattered light: −2.303(L − R )cl σL (0) − σR(0) IL − IR = + 2 IL + IR 2 2r + σL (0) + σR (0)

(48)

Self-assembled Porphyrins at the Mesoscopic Scale

593

where IL and IR indicate the intensities measured when the incident light is left and right circularly polarized, respectively, c is the concentration, l the optical path length, σ(0) the scattering cross section in the forward direction and r the distance between the scattering volume and the detector. In this equation the attenuation of the incident beam before reaching the scattering volume and that of the scattered beam before leaving the cell was taken into account. From a practical point of view the second term of equation 48 can be made small enough to be neglected (for instance by placing the detector far from the sample and by decreasing the angular acceptance for the detector). In the absence of the scattering contribution, the first term is the usual circular dichroism, but in the more general case it must be written as L − R = (aL − aR ) + (sL − sR ), with aL − aR the differential absorption coefficient and N sL − sR = 2.303

Z

2π

dφ

Z

0

π

[σL (θ) − σR (θ)] sinθdθ

(49)

0

the differential scattering coefficient, which is dependent on the scattering all over the angles. The sign of sL − sR , which is not simply related to the sign of the circular intensity differential scattering, depends on the relative orientations of the individual scattering elements constituting the particle and on the particle size. It has been shown[123] that the relative contribution between the circular differential scattering and absorption is related to the third power of the ratio between the distance among the individual scattering elements (and hence the particle size) and the wavelength of the incident light. At any other angle, except zero, the detected signal contains only the circular intensity differential scattering (CIDS): "

#

IL − IR −2.303(L − R )cl (1 + cosθ)2 σL (θ) − σR (θ) = + 2 IL + IR 2 2(1 + cos θ) σL (θ) + σR (θ)

(50)

In the first concentration-dependent term the differential attenuation of the scattered beam, due to the circular differential absorption and scattering, is taken into account. The square brakets contains the dependence of the polarization on the scattering angle, written for the case of point-like scatterers; in most cases this dependence can be also used for an arbitrary scatterer since the deviation from the point-like case is smaller than 5 percent. More generally, however, this dependence, if known, can be explicity considered in eq. 50. The second term characterizes the dependence of the differential scattering cross section on the scattering angle. The two terms of eq. 50 can be separated by changing concentration or by measuring the scattering in the backward direction; indeed, the first term goes to zero at zero concentration or at θ = 180o because the right circularly polarized light becomes left circularly polarized when it is scattered backwards. The CIDS profiles are expecially useful in the determination of the supramolecular structure of self-assembled or organized systems[124, 125, 126]. For helical structures, for instance, according to this approach, the differential scattering displays intense lobes with alternating signs as a function of scattering angle, as also confirmed by experimental observation in different helically organized biological structures[127, 128]. Position, sign, and number of lobes are sensitive indicators of the sense of handedness of the chiral structure for wavelength outside the absorption bands.[129]

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Valentina Villari, Norberto Micali and Luigi Mons´u Scolaro

Figure 24. a) Intensity profile of a t − H2Pagg aggregate in the presence of poly-Lglutammate (PLGA) with Mw = 13600 (λ0 = 632.8nm). The straight line is a power law fit with a slope Df ≈ 2.5. b) Circular differential absorption spectrum in the forward direction and circular differential scattering spectrum at 0.5 ≤ θ ≤ 5o. Experimental conditions: [PLGA]=20µM, [t − H2 Pagg ]=5µM, [NaCl]=0.1M, [acetate buffer]=5mM, pH=4.5.

It is worthnoting that differential scattering effects can also appear for non chiral aggregates and are due, for instance, to the alignement of non isotropic particles or to a not exactly circular polarization of the incident light. Therefore, in differential scattering experiments care must be taken in order to avoid any unwanted artifacts not originating from the supramolecular chirality. As it was discussed at the beginning of the section, aggregation of porphyrins in the presence of chiral templates is a powerful tool for the investigation of the memory effects of chiral ’imprinting’ on the aggregate structure. Fractal aggregates of t − H2 Pagg , whose formation is induced by the ionic strength, take a chiral arrangement when the aggregation process occurs in the presence of a polypeptide as templating chiral agent. Figure 24a) displays the intensity profile of the t − H2 Pagg aggregated according to a DLA mechanism of growth, whereas figure 24b) displays the circular differential extinction coefficient (i.e. the first term of equation 48) and the circular differential scattering coefficient of eq. 49 integrated over 0.5 ≤ θ ≤ 5o and 0 ≤ φ ≤ 2π. The latter coefficient can be approximated with the theoretical one (integrated all over the θ range) if considering that almost all the intensity is scattered at small angle, as indicated by the intensity profile of Figure 24a). In order to obtain the spectra in Figure 24b) a home-made experimental apparatus was

Self-assembled Porphyrins at the Mesoscopic Scale

595

Figure 25. Circular intensity differential scattering against the scattering angle of a t − H2 Pagg aggregate in the presence of poly-L-glutammate (PLGA) with Mw = 1000 (λ0 = 633nm). The inset reports the intensity profile at the same excitation wavelength; the curve represents the fit according to eq. 18 with Df ≈ 2.2. Experimental conditions: [PLGA]=20µM, [t − H2Pagg ]=5µM, [NaCl]=0.1M, [acetate buffer]=5mM, pH=4.5.

designed: as light source and electronic processing a JASCO 500A spectropolarimeter was adopted and a spatial Fourier transform optics was interposed between sample and detector. Just before the detector a spatial filter was placed in the conjugate Fourier plane of the sample, so that circular differential extinction and scattering coefficients were obtained when a pinhole or a beamstop were used as spatial filter, respectively. As far as the differential scattering profile is concerned, it was shown that for helix structure smaller than the wavelength of light CIDS possesses one lobe; in this case it is easier to relate the change in its sign with the change in the long-range chiral arrangement of constituting entities. In the case of aggregated t − H2Pagg in the presence of polypeptide the measured CIDS angular profile (eq. 50) is reported in Figure 25 for excitation wavelength outside the absorption band. It is obtained by replacing the Fourier transform optics with a relay optics in which the image of the sample is formed onto the detector; both optics and detector are placed in a rotating goniometer. In the inset of figure 25 the intensity profile, showing a DLA mechanism of growth, is reported; the bending at low Q values is to be attributed to the finite size of the aggregates. Independently from structural information, however, measurements of the differential scattering are extremely useful as sensitive probe of supramolecular chirality. The higher sensitivity of CIDS with respect to CD for chiral aggregates originates from the fact that the former is a relative measurement of the polarized to the total scattering (see the second term of eq. 50) and the latter an absolute difference of extinction (see the first term of eq. 48). Therefore, small differences in the extinction and scattering cause a small CD signal, but a large CIDS signal, when also the total scattering is small. In this respect porphyrins interacting with chiral biopolymer or self-assembled in the presence of chiral templates constitute a valid system for the detection of supramolecular chi-

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rality due to the high extinction coefficient in the visible range and the resonance scattering effects.

7.

Conclusion

The non-covalent nature of self-assembly, and of porphyrin self-assembly in particular, is the key of the huge variety of structural and optical properties of the formed aggregates. As it has been shown, these properties crucially depend on the kind of intermolecular interaction potential. This review aimed to show that the phenomenology observed in porphyrin solutions and some apparently ’anomalous’ experimental results, like different shape of the final aggregates, mixing order effects, broadening of the absorption bands or unexpected chirality inversion, can be well framed theoretically. The knowledge of the molecular mechanisms underlying the aggregation process helps to drive the system towards the desired target: through the control and tuning of the aggregates shape and size, as well as of the optical response, porphyrin aggregates represent a promising and versatile tool for applications in material sciences, sensor fields and nanotechnology.

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[112] S´anchez, E. J.; Novotny, L.; Xie, X. S. Phys Rev Lett 1999, 82, 4014-4017. [113] Bohren, C.F.; Huffman D.R. Absorption and Scattering of Light by Small Particles; Wiley and Sons, Inc: New York NY, 1998. [114] Akins, D. L.; Ozcelik, S.; Zhu, H. R.; Guo, C. J Phys Chem 1996, 100, 14390-14396; Akins, D. L.; Zhu, H. R.; Guo, C. J Phys Chem 1996, 100, 5420-5425. [115] Palmans, A. R. A.; Meijer, E.W. Angew Chem Int Ed 2007, 46, 8948-8968. [116] Mammana, A.; DUrso, .; Lauceri, R.; Purrello, R. J Am Chem Soc 2007, 129, 80628063. [117] Bustamante, C.; Maestre, M. F.; Tinoco, I. Jr. J Chem Phys 1980, 73, 4273-4281. [118] Bustamante, C.; Maestre, M. F.; Tinoco, I. Jr. J Chem Phys 1980, 73, 6046-6055. [119] Bustamante, C.; Maestre, M. F.; Tinoco, I. Jr. J Chem Phys 1981, 74, 4839-4850. [120] Keller, D.; Bustamante, C. J Chem Phys 1986, 84, 2972-2980. [121] Barron, L. Molecular light scattering and optical activity, Cambridge University press: Cambridge UK, 2004. [122] Harris, R.A.; McClain, W.M. J Chem Phys 1977, 67, 265-268. [123] Bustamante, C.; Tinoco, I. Jr.; Maestre, M. F. Proc Natl Acad Sci 1983, 80, 35683572. [124] Bustamante, C.; Samof, B. S.; Builest, E. Biochemistry 1991, 30, 5661-5666. [125] Phillips, C. L.; Mickols, W. E.; Maestre, M. F.; Tinoco, I Jr Biochemistry 1986, 25, 7803-7811. [126] Lakhwani, G.; Meskers, S. C. J.; Janssen, R. A. J. J Phys Chem B 2007, 111, 51245131. [127] Tinoco, I. Jr.; Mickols, W.; Maestre, M. F.; Bustamante, C. Annu Rev Biophys Biophys Chem 1987, 16, 319-349. [128] Garab, G.; Wells, S.; Finzi, L.; Bustamante, C. Biochemistry 1988, 27, 5839-5843. [129] Bustamante, C.; Maestre, M. F.; Keller, D. Biopolymers 1985, 24, 1595-1612.

SHORT COMMUNICATIONS

In: Nanotechnology… Editors: C.J. Dixon and O.W. Curtines, pp. 605-613

ISBN: 978-1-60692-162-3 © 2010 Nova Science Publishers, Inc.

Short Communication A

THE INFLUENCE OF THIOPHENE ADDITION ON CATALYTIC PYROLYSIS OF POLY (DIMETHYL SILOXANE) 1

K.F. Cai*,1, C.W. Zhou1, A.X. Zhang1 and J.L. Yin2

Tongji University, Functional Materials Research Laboratory, 1239 Siping Road, Shanghai 200092, China 2 Tongji University, Shanghai Key Laboratory of Development and Application for Metal-Functional Materials, 1239 Siping Road, Shanghai 200092, China

Abstract The influence of thiophene addition on the pyrolysis of poly(dimethyl siloxane) catalyzed by ferrocene at ~1050 oC in Ar was studied. The as-synthesized product was characterized by Xray diffraction, scanning electron microscopy, transmission electron microscopy and highresolution transmission electron microscopy. The thiophene addition caused several changes. Firstly, the yield of the product was increased by several times and the diameters of the product were somewhat increased. Secondly, the product was changed from only SiC/SiO2 nanocables to a mixture of SiC/SiO2 nanocables and SiC-SiO2 side-by-side nanowires. Thirdly, more “Y” type nanostructures were found. Finally, the growth process of the product was altered as the nanostructures each had a polyhedral FeS nanoparticle rather than spherical Fe nanoparticle. However, lengths of the product were still on the millimeter scale. The promotion mechanism of thiophene addition was also analyzed.

Keywords: synthesis; nanocables; nanowires; pyrolysis; silicon carbide; silicon oxide.

1. Introduction Because nanocables have one-dimensional (1D) features of both nanowires and nanotubes in axial direction and are ideal heterostructures in radial direction, they are *

E-mail address: [email protected]. Corresponding author, Tel.:/Fax: +86-21-65980255.

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excellent candidates for fundamental scientific research and potential applications in nanodevices. Recently, nanocables have attracted great attention. Until now, many kinds of nanocables, including Zn/ZnS [1], Zn/ZnO [2,3], SnO2/Sn [4], In2O3/Sb [5], C/SiO2 [6], ZnSe/SiO2 [7], Ge/SiO [8], Ge/SiCNx [9] ZnO/ZnS [10], CdS/TiO2 [11], ZnS/SiC [12], CdS/Si [13], Ag/peptide/Au [14], SiC/SiO2/C [15], Si3N4/Si/SiO2 [16], and so on have been synthesized. SiC/SiO2 nanocable system is one of the most extensively studied ones. Up to now, many methods, such as laser ablation [15, 17], arc discharge [18, 19], carbon thermal reduction [20-23], chemical vapor reaction [24], and others [25-27] have been developed for synthesis of SiC/SiO2 nanocables. Recently, we developed a simple route to SiC/SiO2 nanocables of a few millimeters in length by pyrolysis of poly(dimethyl siloxane) (PDMS) at 1050oC in a flowing Ar atmosphere [28]. Moreover, we [29] investigated the pyrolysis of PDMS catalyzed by ferrocene at 1050oC in a flowing Ar atmosphere. The yield was increased by several times and the nanocables synthesized became very uniform with a spherical iron nanoparticle at their tips. Most nanocables were about 5 to 10 nm in diameter. This improvement owes to the iron from ferrocene. When ferrocene is used as a catalyst precursor to synthesize carbon nanotubes, thiophene is usually added for promoting the growth of carbon nanotubes [30-32]. However, the exact promotion mechanism of thiophene has not been reported yet. In this work, the pyrolysis of PDMS catalyzed by ferrocene and thiophene has been investigated in order to know the influence of thiophene addition.

2. Experimental Procedure The synthesis was carried out in the same system as that for the synthesis of SiC/SiO2 nanocables described in our previous work [28, 29] (see schematic diagram of the system in reference [33]). Briefly, the system contained a quartz tube (inner diameter ~85 mm) with a sealed end. The sealed end was put in the hot zone of a muffle furnace. In a typical synthesis, about 6 ml analytical pure PDMS and 0.1g ferrocene were placed in a 10-ml corundum crucible and the crucible was put on a mullite substrate (~150 × 80 × 12 mm3). A 5-ml corundum crucible loaded with ~2 ml thiophene (99% purity, Alfa Aeser) was placed on the mullite substrate beside the 10-ml crucible at the downstream side. The substrate was carefully pushed into the hot zone of the quartz tube. The open end of the big quartz tube was sealed with a rubber stopper for gas inlet and outlet. After completely eliminating the air in the tube furnace, the furnace was heated up to ~1050 oC at a rate of 10 oC/min, and held at 1050 oC for 2h, in flowing Ar at a rate of ~8 ml/min. The furnace was cooled from 1050 oC to 700 oC in 10 h, and then cooled naturally to room temperature. White wool-like product was observed on the surface of the substrate and the thickness of the product was about 2 mm. The size, structure, and composition of the product were examined by powder X-ray diffraction (XRD, Rigaku, D/max2550), scanning electron microscopy (SEM, JSM5510), transmission electron microscopy (TEM, H-800), and high-resolution TEM (HRTEM, JEM2100F, JEOL) equipped with energy dispersive X-ray spectroscopy (EDS), respectively.

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3. Results and Discussion With the help of thiophene, the yield of the product was about ten times higher than that of the product from pyrolysis of PDMS catalyzed by only ferrocene, which indicates that thiophene is a good promoter.

Figure 1. (a) XRD pattern of the as-synthesized product, (b) a typical SEM image of the product.

Figures 1(a) and (b) are XRD pattern and typical SEM image of the product, respectively. The XRD pattern (figure 1(a)) of the as-grown product has a hillside background and the peaks corresponding to SiC (JCPDS card file, No.29-1129), which is very similar to that for SiC/SiO2 nanocables shown in reference [29]. However, several additional peaks indicated by asterisks appear in figure 1(a). These peaks can be indexed to FeS (JCPDS card file No.231120), which should be resulted from the addition of the ferrocene and thiophene. SEM image (figure 1(b)) shows that the product mainly consists of nanowires. Unlike the entangled

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SiC/SiO2 nanocables from pure PDMS as shown in reference [28], the nanowires are mainly straight, which is similar to the SiC/SiO2 nanocables from pyrolysis of PDMS catalyzed by ferrocene as shown in reference [29]. The nanowires are dense and fine. The lengths of the nanowires are too long to be measured under SEM. They are estimated to be on the millimeter scale with an optical microscope. TEM observations at low magnifications reveal that most of the nanowires have diameters in the range of 5 to 25 nm and that a small amount of the nanowires are thick with diameters up to 70 nm (see figure 2(a)), which are somewhat thicker than those from pyrolysis of PDMS catalyzed by only ferrocene. Some of the nanowires have “Y” junction. The Y-junction usually accompanies twin structure whose twin plane is along the central line of one branch of the junction, shown clearly as in the inset of figure 2(b).

Figure 2. Continued on next page.

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Figure 2. (a) A typical TEM image at low magnification of the product, small black arrow indicates a faceted nanoparticle, (b) TEM image showing a Y-junction, the inset is an HRTEM image of the zone marked by a white square in (b), showing one branch of the Y-junction contains a twin and the twin plane is in the central of the branch indicated by arrows, (c) HRTEM image of a SiC/SiO2 nanocable, clearly showing the lattice spacing of about 0.25 nm, which corresponds to the distance between (111) planes of SiC, and the growth direction is indicated by the long arrow, (d) typical EDS spectrum recorded on the rod part of the nanocable, (e) TEM image of the product at low magnification, the thicker 1D nanostructure looks like a nanocable, (f) HRTEM image of the thicker 1D nanostructure, revealing that it is a side-by-side biaxial nanowires, an array of small arrows indicate the crystal defects.

TEM observations at high magnifications indicate that all the nanowires have a core-shell structure. Usually a thicker core has a thicker shell. The thickest shell is about 5 nm. The shells are much thinner than those synthesized from pure PDMS and are similar to those from pyrolysis of PDMS catalyzed by only ferrocene. HRTEM observation reveals that the cores are single crystalline and that the shells are amorphous, shown as in figure 2(c). The hillside background of XRD pattern in figure 1(a) should be related to the amorphous shell. The distance between two adjacent lattice planes is about 0.25 nm, which corresponds to the spacing between (111) lattice planes of cubic SiC (JCPDS card file No. 29-1129), suggesting that the core is β-SiC with <111> growth direction. EDS analysis reveals that the nanocables are composed of C, O and Si (see figure 2(d)). The Cu peaks are from the copper grid of TEM sample holder. Besides the nanocables, side-by-side SiC-SiO2 composite nanowires are occasionally observed under HRTEM. Figure 2(e) shows a low-magnification TEM image of the product. The thickest 1D nanostructure in figure 2(e) looks like a nanocable with a 5-nm thick shell. However, it is observed under HRTEM (figure 2(f)) that there does have an amorphous layer (~5 nm) at its right side, while no amorphous layer at its left side. That looks like (but in fact not) an amorphous layer at its left side under TEM results from many lattice defects. The defects form into a line that is parallel to the growth direction of the 1D nanostructure. Namely, it is a side-by-side SiC-SiO2 biaxial nanowire. Wang et al. [21] have also synthesized such nanowires as well as SiC/SiO2 nanocables using amorphous SiO and carbon/graphite as starting materials at ~1500oC for 12h. Like nanocables, the side-by-side nanowires could be potentially useful for high-strength composites [21].

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Therefore, the product is composed of SiC/SiO2 nanocables and side-by-side SiC-SiO2 biaxial nanowires. For simplicity, we call them composite nanowires hereinafter.

Figure 3. (a) another TEM image at low magnification of the product, clearly showing a faceted nanoparticle at the tip of a nanocable, the inset is HRTEM image of the zone marked by a white square on the nanoparticle, showing that the nanoparticle is single crystalline coated with a thin amorphous layer, (b) EDS spectrum recorded on the nanoparticle, (c) TEM image of a separated nanocable with a faceted nanoparticle at its tip, the inset is SAED pattern taken on the nanoparticle, indicating the nanoparticle is cubic FeS.

Figures 3(a) to (c) show TEM images and EDS spectrum of the composite nanowires. Figure 3(a) shows a faceted nanoparticle at a tip of a thick nanowire (see faceted nanoparticle also in figure 2(a)). The HRTEM image (inset in figure 3(a)) of the zone marked by a white square in figure 3(a) reveals that the faceted nanoparticle is single crystalline coated with a thin amorphous layer. EDS analysis (figure 3(b)) shows that the nanoparticle contains Fe and

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S. The C and Cu peaks in figure 3(b) originate from the TEM sample holder, and the Si and O peaks are from the thin amorphous SiO2 layer at the surface of the nanoparticle (see the inset of figure 3(a)). Figure 3(c) shows a TEM image of a separated nanowire with a faceted nanoparticle at its tip. SAED pattern (inset in figure 3(c)) taken on the nanoparticle indicates that the nanoparticle is cubic FeS (JCPDS card, No. 23-1123). The 1D nanostructures have faceted FeS tips, indicating that the growth of the composite nanowires is governed by the conventional vapor-liquid-solid (VLS) mechanism [34]. The growth process of the composite nanowires is obviously different from that of the nanocables from pure PDMS, and it is also somewhat different from that of the nanocables from pyrolysis of PDMS catalyzed by only ferrocene since the nanoparticle at the tips of the nanostructures is faceted FeS rather than spherical Fe. Based on the above results, we propose the growth process of the composite nanowires as follows. Ferrocene begins to vaporize at about 185oC and decomposes above 400oC [31]. PDMS vaporizes above 200oC and as the temperature is high enough the cleavage of the Si-C bond occurs to produce methyl (-CH3) groups, and the methyl groups decompose to form CH4, C2H6, C and H2 below 600oC [35]. As all the methyl groups are separated from the PDMS and the temperature is high enough, SiO vapor forms [28]. The decomposed ferrocene is reduced by the H2 to form atomic iron, and thiophene decomposes to produce atomic sulfur around 600oC. The iron atoms and sulfur atoms, carried by Ar gas mixed with hydrocarbon gases to the downstream site, react with each other to produce FeS species that agglomerate into FeS clusters. At high temperatures the clusters become FeS droplets and deposit on the substrate. The droplets serve as preferential sites for absorption of the hydrocarbon gases and SiO vapor. Hydrocarbon gases are catalytically decomposed into carbon atom and the carbon atom diffuses into the FeS droplet and then FeS-Si-O-C alloy forms. As the FeS-Si-O-C alloy becomes supersaturated, SiC and SiO2 concurrently segregate from the alloy droplet by the reaction: 2SiO+C→ SiC+SiO2. Because the melting point of SiC is much higher than that of SiO2, SiC solidifies earlier. A SiC nanorod is formed along <111> direction, and then amorphous SiO2 is nucleated on the outer surface of the newly formed SiC nanorod. The growth of the composite nanowires continues as long as the catalyst alloys remain in a liquid state and the reactants are available. Finally, the composite nanowires with cubic FeS (JCPDS card, No. 23-1123) nanoparticles at their tips form. Note that the phase structure of the faceted FeS nanoparticles is different from that of the FeS revealed by XRD. This is probably because the amount of both ferrocene and thiophene is overmuch used. Earlier deposited FeS droplets are covered by the later deposited ones; therefore, the former droplets have no chance to absorb the reactant gases to become FeS (JCPDS card, No. 23-1123) nanoparticles but to form FeS (JCPDS card, No. 23-1120) nanoparticles. The SEM and TEM samples could be contaminated by such nanoparticles. This could be the reason that some flocky dots and some agglomerated nanoparticles appear in the SEM image (see figure 1(b)) and TEM images (see the nanoparticles indicated by the white arrows in figures 2(a), 2(b), 2(e) and 3(a)), respectively. On the other hand, the FeS in the FeS-Si-C-O alloy finally crystallizes into cubic phase (JCPDS card, No. 23-1123) due to the droplet contained Si, C and O. What is the promotion mechanism of thiophene addition? It can be described as follows. The Fe from ferrocene and S from thiophene react with each other at high temperature to form FeS. The melting point of FeS (~1190oC) is much lower than that of Fe (1534 oC); therefore, the FeS-Si-C-O alloy solidifies at lower temperature than Fe-Si-C-O alloy. Namely,

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the nucleation and growth process of the composite nanowires will last longer, and the assynthesized product has a higher yield than the product from pyrolysis of PDMS catalyzed by only ferrocene. More 1D nanostructures with Y-junction are found in the present product (see figure 2(a)) than in the product from pyrolysis of PDMS catalyzed by only ferrocene. This implies that the formation of the Y-junctions is related to the thiophene (sulfur). Deepak et al. [36] have also found that thiophene (sulfur) plays an important role in the formation of carbon nanotubes with Y-junction. The semiconductive Y-junctions could render them useful for exploitation in nanoelectronics. The composite nanowires are not as uniform as the nanocables from pyrolysis of PDMS catalyzed by only ferrocene. This is probably related to the amount of catalyst precursors used. The influence of adjusting the amount of catalysts on the product will be further studied.

4. Conclusion Pyrolysis of PDMS catalyzed by ferrocene and thiophene is investigated. Compared with the case of pyrolysis of PDMS catalyzed by only ferrocene, the present case has much higher yield. The product consists of SiC/SiO2 nanocables and side-by-side SiC-SiO2 biaxial nanowires. The product has more Y-junctions and the nanoparticles at the tips of the 1D nanostructures are cubic FeS. Thiophene is a good promoter for the growth of SiC/SiO2 nanocables and side-by-side SiC-SiO2 biaxial nanowires.

Acknowledgment This work was partly supported by Shanghai Pujiang Program and Natural Science Foundation of China (50371062, 50872095).

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[9] Mathur, S.; Shen, H.; Donia, N.; Ruegamer, T.; Sivakov, V.; Werner, U. J. Am. Chem. Soc. 2007, 129, 9746-9752. [10] Sun, C. W.; Jeong, J. S.; Lee, J. Y. J. Crystal Growth 2006, 294, 162-167. [11] Hsu, M. C.; Leu, I. C.; Sun, Y. M.; Hon, M. H. J. Crystal Growth 2005, 285, 642-648. [12] Hu, J. Q.; Bando, Y.; Zhan, J. H.; Golberg, D. Appl. Phys. Lett. 2004, 85, 2932-2934. [13] Fu, X. L.; Ma, Y. J.; Li, P. G.; Chen, L. M.; Tang, W. H.; Wang, X.; Li, L. H. Appl. Phys. Lett. 2005, 86, 143102. [14] Carny, O.; Shalev, D. E.; Gazit, E. Nano Lett. 2006, 6, 1594-1597. [15] Zhang, Y.; Suenaga, K.; Colliex, C.; Iijima, S. Science 1998, 281, 973-975. [16] Wu, X. C.; Song, W. H.; Zhao, B.; Huang, W. D.; Pu, M. H.; Sun, Y. P.; Du, J. J. Solid State Comm. 2000, 115, 683-686. [17] Shi, W. S.; Zheng, Y. F.; Peng, H. Y.; Wang, N.; Lee, C. S.; Lee, S. T. J. Am. Ceram. Soc. 2000, 83, 3228-3230. [18] Liu, X. M.; Yao, K. F. Nanotechnology 2005, 16, 2932-2935. [19] Li, Y. B.; Xie, S. S.; Zou, X. P.; Tang, D. S.; Liu, Z. Q.; Zhou, W. Y.; Wang, G. J. Cryst. Growth. 2001, 223, 125-128. [20] Meng, G. W.; Zhang, L. D.; Mo, C. M.; Zhang, S. Y.; Qin, Y.; Feng, S. P.; Li, H. J. J. Mater. Res. 1998, 13, 2533-2538. [21] Wang, Z. L.; Dai, Z. R.; Gao, R. P.; Bai, Z. G.; Gole, J. L. Appl. Phys. Lett. 2000, 77, 3349-3351. [22] Yang, Z. X.; Zhou, W. M.; Zhu, F.; Zhang, Y. F. Mater. Chem. Phys. 2006, 96, 439-441. [23] Ryu, Y.; Tak, Y.; Yong, K. Nanotechnology 2005, 16, S370-S374. [24] Wang, C. S.; Zhang, J. L.; Meng, A. L.; Zhang, M.; Li, Z. J. Physica E 2007, 39, 128132. [25] Xing, Y. J.; Hang, Q. L.; Yan, H. F.; Pan, H. Y.; Xu, J.; Yu, D. P.; Xi, Z. H.; Xue, Z. Q.; Feng, S. Q. Chem. Phys. Lett. 2001, 345, 29-32. [26] Liu, D. F.; Xie, S. S.; Yan, X. Q.; Ci, L. J.; Shen, F.; Wang, J. X.; Zhou, Z. P.; Yuan, H. J.; Gao, Y.; Song, L.; Liu, L. F.; Zhou, W. Y.; Wang, G. Chem. Phys. Lett. 2003, 375, 269-272. [27] Zhu, Y. Q.; Hu, W. B.; Hsu, W. K.; Terrones, M.; Grobert, N.; Hare, J. P.; Kroto, H. W.; Walton, D. R. M.; Terrones, H. J. Mater. Chem. 1999, 9, 3173-3178. [28] Cai, K. F.; Lei, Q.; Zhang, L. C. J. Nanosci. Nanotechnol. 2005, 5, 1925-1928. [29] Cai, K. F.; Zhang, A. X.; Yin, J. L. Nanotechnology 2007, 18, 485601. [30] Sen, R.; Govindaraj, A.; Rao, C. N. R. Chem. Phys. Lett. 1997, 267, 276-280. [31] Cheng, H. M.; Li, F.; Su, G.; Pan, H. Y.; He, L. L.; Sun, X.; Dresselhaus, M. S. Appl. Phys. Lett. 1998, 72, 3282-3284. [32] Zhu, H. W.; Xu, C. L.; Wu, D. H.; Wei, B. Q.; Vajtai, R.; Ajayan, P. M. Science 2002, 296, 884-886. [33] Cai, K. F.; Lei, Q.; Zhang, A. X. J. Nanosci. Nanotechnol. 2007, 7, 580-583. [34] Wagner, R. S.; Ellis, W. C. Trans. Metal. Soc. AIME 1965, 233, 1053-1064. [35] Plawsky, J. L.; Wang, F.; Gill, W. N. AIChE J. 2002, 48, 2315-2323. [36] Deepak, F. L.; Govindaraj, A.; Rao, C. N. R. Chem. Phys. Lett. 2001, 345, 5-10.

In: Nanotechnology… Editors: C.J. Dixon and O.W. Curtines, pp. 615-620

ISBN: 978-1-60692-162-3 © 2010 Nova Science Publishers, Inc.

Short Communication B

NANOFINISHING OF COTTON TEXTILES N. Vigneshwaran* and Virendra Prasad Chemical and Biochemical Processing Division, Central Institute for Research on Cotton Technology, Adenwala Road, Matunga, Mumbai – 400 019, India

Abstract Nanotechnology revolutionized every field in science and technology. Recently, its usefulness in nanofinishing of cotton fabrics by imparting functional properties like antimicrobial, UVresistance, self-cleaning and drug-delivery is well documented. In addition, enhancement in comfort properties of cotton textiles is also being evaluated with the help of nanofinishing. With judicial use of nanomaterials, keeping in view their bio-safety and environmental impact issues, nanofinishing will be a great boon to the users of cotton textiles.

With the advancement in science and technology, a new era has emerged in the realm of textile processing. Apart from the regular finishing of textile materials for colour and handle, functional properties like antimicrobial, UV-resistant, stain and water repellent, wrinkle resistant, flame resistant, moisture control and drug delivery need to be imparted through novel finishes for high-end applications. Nanofinishing or nanocoating the surfaces of textiles and clothing is one approach to the production of highly active surfaces to have UV-blocking, antimicrobial, and self-cleaning properties [1,2]. Fabrics are an excellent medium for the growth of microorganisms when the basic requirements such as nutrients, moisture, oxygen and appropriate temperature are present. Compared to synthetic fibres, natural fibres like cotton are more susceptible to microbial attack resulting in yellowing, bad odour and strength loss of the fabric. Also, cotton textiles are not resistant to ultraviolet (UV) radiation and highly susceptible to fire hazard. Hence, there is a need for functional finishing in cotton textiles to overcome the above mentioned problems. Among the various metal salts, silver nitrate is the commonly used antibacterial agent [3,4]. Silver nitrate has been used historically as an antiseptic agent. They bind to protein molecules inhibiting the cellular metabolism, and eventually the microorganisms die. A *

E-mail address: [email protected] (Corresponding author)

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major disadvantage of silver nitrate is that it stains everything it touches to brown-black when exposed to light [5]. Silver nanoparticles, in contrast, does not produce stains while retaining their antibacterial property. In our earlier article, [6] we have reported a novel, in situ synthetic route to prepare silver nanoparticles on the surface of cotton fabrics. These silver nanoparticles impregnated cotton fabrics showed excellent antibacterial activity against Staphylococcus aureus and bacteriostasis activity against Klebsiella pneumoniae (Figure 1).

Figure 1. Cotton fabrics impregnated with silver nanoparticles. The increasing colour intensity corresponds to increasing concentration of nano-silver.

Scanning electron microscopic analysis revealed the presence of silver nanoparticles on the surface of cotton fibre (Figure 2). In addition, nano-impregnated fabrics were yellow to pink coloured due to surface plasmon resonance of silver nanoparticles and provided better protection against UV radiation due to its absorption in the near-UV region. Another research group from the Hanyang University, Korea [7] padded the colloidal silver solution onto the textile fabrics, including cotton, polyester, cotton/polyester and cotton/spandex blended fabrics. They have also reported efficient antibacterial activity against both Staphylococcus aureus and Klebsiella pneumoniae.with good laundering durability. Duran et. al. [8] demonstrated the dip-coating of silver nanoparticles produced by fungal process on textile fabrics and their effluent treatment. The silver ions were reduced extracellularly by Fusarium oxysporium to generate stable silver nanoparticles in water. These nanoparticles were coated onto cotton fabrics by dip coating and the coated fabrics showed excellent antibacterial activity. Simultaneously, the effluent generated was treated with Chromobacterium violaceum to reduce the silver nanoparticles

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concentration. Nano-silver particles are widely applied to socks in order to avoid growth of bacteria. In addition, nano-silver can be applied to a range of other healthcare products such as dressings for burns, scald, skin donor and recipient sites. In silver containing wound care devices, the silver in contact with wound enters it and becomes absorbed by undesirable bacteria and fungi, so that those organisms get killed [9].

Figure 2. Scanning electron microscopic image of cotton fibre surface deposited with silver nanoparticles.

ZnO nanoparticles score over nano-silver in cost-effectiveness, whiteness, and UVblocking property. The UV-blocking property of a fabric is enhanced when a dye, pigment, delustrant, or UV absorber finish is present that absorbs UV radiation and blocks its transmission through the fabric to the skin [10]. Metal oxides like ZnO as a UV-blocker are more stable when compared to organic UV-blocking agents. Hence, the nanoform ZnO will really enhance the UV-blocking property due to the increased surface area and intense absorption in the UV region. Our research finding [11] proved the excellent antibacterial activity against two representative bacteria, Staphylococcus aureus and Klebsiella pneumoniae and promising protection against UV radiation by the nano-ZnO impregnated cotton textiles. Figure 3 shows the process for coating of cotton fabrics with nano-ZnO prepared in our laboratory.

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Figure 3. Schematic representation of nano-coating process for cotton fabrics.

The Ultraviolet Protection Factor (UPF) was calculated using the following equation as per the AATCC test method number 183-2004 [12]: 400

∑

UPF =

E λ × S λ × Δλ ----------------------------------λ = 280 400

∑

Eλ × S λ × Tλ × Δλ

λ = 280

where, E λ is relative erythermal spectral effectiveness, S λ is solar spectral irradiance, Tλ is average spectral transmission of the specimen, Δλ is measured wavelength interval (nm). The UPF equation weighs the UV-B radiation more heavily than UV-A. Apart from nano-ZnO, nano-silver coating onto cotton fabrics was found to increase the UPF factor due to their absorption in the near-UV-region. [6,13] Nanocrystalline titanium dioxide coatings have received much attention as photocatalysts in practical applications such as environmental purification, deodorization, sterilization, antifouling and self-cleaning glass due to their high oxidizing ability, nontoxicity, long term stability and low cost. Among the different crystalline phases of titania, anatase is reported to have the best performance. Daoud and Xin [14] successfully grown anatase nanocrystallites on cotton fabrics and these fabrics could be made into self-cleaning clothes that tackle dirt, environmental pollutants and harmful microorganisms. Also, they have reported [15] a transparent thin layer of nanocrystalline titania coating on cotton textiles by a dip-pad–drycure process. These titania coated cotton textiles possess significant photocatalytic selfcleaning properties, such as bactericidal activity, colorant stain decomposition and degradation of red wine and coffee stains. Figure 4 shows the photo catalytic activity of titania nanoparticle.

Nanofinishing of Cotton Textiles

619

Figure 4. Photo catalytic activity of titania nanoparticles.

New generation medical textiles are important growing field that require functional properties like bacteriostatic, anti-viral, fungistatic, non-toxic, high absorbent, non-allergic, breathable, haemostatic and biocompatible. So, apart from metal and metal oxide nanomaterials, nanoscale biological materials like enzymes and drugs are necessary to add specific functionality to medical textiles [16]. Specialized nanomaterials functionalized with ligands can be introduced on the surface of cotton textiles with the aim to absorb odours, provide strong and durable antibacterial property, smooth pain and relieve irritation. Also, such value added textiles could be of immense use in tissue engineering, drug delivery and protective clothing.

References [1] [2] [3] [4] [5] [6] [7] [8]

Bozzi, A; uranova, TY; Kiwi, J. J. Photochem. Photobiol. A, 2005, 172, 27. Qian, L. AATCC Review, 2004, 4, 14. Purwar, R; Joshi, M. AATCC Review, 2004, 4, 22. Ramachandran, Y; Rajendrakumar, K; Rajendran, R. Institution of Engineers (India) Journal, 2004, 84, 42. Parikh, DV; Fink, D; Rajasekharan, K; Sachinvala, ND; Sawhney, APS; Calamari, TA. Textile Res. J., 2005, 75, 134. Vigneshwaran, N; Kathe, AA; Varadarajan, PV; Nachane, RP; Balasubramanya, RH. Functional Finishing of Cotton Fabrics Using Silver Nanoparticles. J. Nanosci. Nanotechnol., 2007, 7, 1-5. Lee, HJ; Yeo, SY; Jeong, SH. Antibacterial effect of nanosized silver colloidal solution on textile fabrics. J. Mater. Sci., 2003, 38, 2199-2204. Duran, N; Marcato, PD; Souza, GIHD; Alves, OL; Esposito, E. Antibacterial Effect of Silver Nanoparticles Produced by Fungal Process on Textile Fabrics and Their Effluent Treatment. Journal of Biomedical Nanotechnology, 2007, 3, 203-208.

620

N. Vigneshwaran and Virendra Prasad

[9] Forner-Cordero, I; Navarro-Monsolu, R; Munoz-Langa, J; Alcober-Fuster, P; RelMonso, P. Use of nanocrystalline silver dressing in lymphatic ulcers in patients with chronic lymphoedema. Journal of Wound Care, 2007, 16, 235-238. [10] Hustvedt G; Crews PC. J. Cot. Sci., 2005, 9, 47. [11] Vigneshwaran, N; Sampath kumar; Kathe, AA; Varadarajan PV; Virendra Prasad. Nanotechnology, 2006, 17, 5087-5095. [12] AATCC Test Method 183-2004, AATCC Technical Manual, 2005, 338. [13] Gorensek, M; and Recelj, P. Nanosilver Functionalized Cotton Fabric. Textile Res. J., 2007, 77(3), 138-141. [14] Daoud, WA; and Xin, JH. J. Am. Ceram. Soc., 2004, 87, 953. [15] Qi, K; Daoud, WA; Xin, JH; Mak, CL; Tang, W; Cheung, WP. Self-cleaning cotton. Journal of Materials Chemistry, 2006, 16, 4567-4574. [16] Petrulyte, S. Advanced textile materials and biopolymers in wound management. Danish medical bulletin, 2008, 55(1), 72-77.

INDEX adhesion, xv, 108, 137, 147, 148, 149, 509, 512, 514, 570 adlayers, 38 3,4-ethylenedioxythiophene, 31 adsorption, 9, 103, 122, 144, 302, 328, 395 aerosols, 52 A aerospace, 311 Ag, ix, 8, 10, 20, 24, 33, 53, 71, 72, 73, 74, 75, 76, Aβ, 348, 349, 572 77, 79, 82, 84, 85, 86, 87, 88, 90, 93, 95, 96, 98, 99, 100, 194, 195, 217, 279, 280, 281, 282, 283, AAS, 139 284, 285, 286, 287, 288, 290, 320, 321, 355, 361, absorption coefficient, 117, 118, 123 541, 606 absorption spectra, 145, 484, 538, 539, 540, 541, 546, age, 425, 535 547, 548, 563, 585 agent, 8, 112, 115, 119, 127, 148, 335, 561, 578, 591, accelerator, 392 594, 615 acceptor, 291, 504 agents, x, 107, 114, 115, 119, 120, 123, 125, 127, access, xi, 18, 27, 293, 295 134, 140, 142, 561, 617 accessibility, 85 aggregates, 393, 395, 410, 418, 422, 430, 431, 484, accommodation, 389, 390, 391 489, 497, 505, 560, 561, 562, 563, 565, 569, 570, accuracy, 19, 77, 108, 171, 177, 189, 299, 376 571, 574, 575, 576, 578, 579, 580, 581, 584, 585, acetate, viii, 51, 53, 67, 129, 246, 566, 594, 595 586, 588, 590, 591, 592, 594, 595, 596 acetic acid, 140 aggregation, 121, 131, 312, 315, 376, 415, 421, 425, acetone, 129, 439, 511 561, 565, 566, 567, 568, 570, 571, 573, 574, 575, acid, viii, 12, 16, 17, 30, 32, 33, 38, 51, 53, 59, 113, 578, 579, 580, 581, 582, 584, 586, 588, 589, 590, 120, 121, 129, 140, 143, 247, 318, 320, 323, 328, 591, 594, 596 412, 492, 560, 563, 570, 571, 580, 585, 590, 591 aggregation process, 561, 567, 568, 570, 571, 578, acidic, 33, 111, 114, 115, 126, 140, 245, 312, 565, 588, 590, 596 571, 575, 579, 580, 584, 585, 590 aging, 117, 119, 145, 146, 149, 246, 335, 412, 581 acoustic, 436, 453, 514 aging process, 246 acoustic signals, 514 aid, 95, 108, 261 acquisitions, 74, 99 AIP, xv, 461, 462, 464, 466, 467, 469, 470, 509, 510, acrylate, 310, 316, 318 516, 517, 518, 522, 523 acrylic acid, 53, 120, 318 air, 7, 11, 14, 32, 33, 56, 91, 92, 93, 121, 247, 282, ACS, 108, 140 283, 284, 335, 350, 370, 412, 413, 414, 415, 416, activation, 148, 172, 246, 331, 332, 333, 340, 343, 419, 420, 422, 423, 425, 426, 427, 428, 429, 492, 345, 347, 348, 349, 350, 352, 353, 354, 362, 363, 512, 531, 606 364, 365, 366, 367, 368 albumin, 130 activation energy, 172, 246 alcohol, viii, 51, 53, 130, 131, 493 active transport, 114 alkaline, 16, 17, 135 actuators, 77, 332, 339, 342, 347, 349 alkane, 38, 39 acute, 126, 127, 147 alloys, 7, 39, 102, 137, 152, 332, 354, 554, 611 acute myeloid leukemia, 127 alternatives, 294, 295, 297, 298 additives, 8, 245, 347, 355 aluminum, viii, 4, 6, 8, 9, 24, 25, 27, 31, 247, 248 adenine, 38, 103 aluminum oxide, viii, 4, 31, 247 adenocarcinoma, 123 amine, 245, 320

3

622 amino, 321, 327, 560 amino acid, 560 amino groups, 321 ammonia, viii, ix, 51, 53, 64, 66, 68, 411, 412, 413, 415, 418, 421, 422, 423, 424, 425, 426, 427, 428 ammonium, 130, 321, 565 ammonium chloride, 565 amorphization, 332, 343, 351, 354, 369 amorphous, xiv, 5, 27, 35, 126, 277, 282, 337, 343, 344, 351, 354, 355, 356, 363, 364, 368, 397, 461, 465, 479, 480, 481, 482, 488, 489, 493, 496, 505, 510, 520, 534, 538, 540, 548, 550, 609, 610, 611 amorphous carbon, 35, 126, 397 amplitude, 20, 21, 91, 93, 333, 375, 401, 572, 573, 582 analytical models, 386 anatase, xiv, xv, 31, 338, 479, 480, 481, 482, 489, 490, 492, 493, 494, 496, 497, 498, 499, 500, 502, 505, 506, 618 angiogenesis, 110, 111, 143, 144 Angiogenesis, 110, 141 angiopoietin, 111 aniline, 13, 24, 31, 33, 34, 40 animals, 128, 133 anisotropy, 153, 161, 165, 167, 168, 207, 212, 255, 277, 278, 291, 335, 543, 544, 554, 555, 572, 574, 575 annealing, xvi, 7, 161, 167, 224, 344, 345, 357, 361, 362, 373, 463, 488, 489, 490, 493, 495, 496, 497, 500, 505, 523, 525, 526, 528, 530, 531, 532, 534, 535, 536, 537, 539, 540, 541, 547, 549, 551 annihilation, 171, 541, 562 anode, 17, 28, 249, 252, 267 anomalous, 261, 579, 592, 596 antenna, 481, 561, 578 antenna systems, 561 antibacterial, 615, 616, 617, 619 antibiotic, 128 antibody, 119, 120 anticancer, 139, 142, 147, 560 anti-cancer, x, 107, 108, 109, 111, 112, 127, 133, 134, 138 anticancer drug, 147 antiferromagnetic, 153, 345 antigen, 119, 141, 143 anti-HER2, 123 antioxidant, 133 antitumor, 142 apoptosis, 126, 147, 148 apoptotic, 142 aqueous solution, 23, 28, 30, 35, 36, 42, 144, 245, 246, 249, 252, 490, 565, 574, 579, 584 aqueous suspension, 420 ARC, 435 arc plasma jet, 455, 456 argon, 511 aromatic rings, 323 arsenic, 148 arsenide, 161, 276

Index ASI, 599 aspect ratio, 18, 20, 238, 245, 246, 269, 381, 382, 385, 387, 388, 394, 404, 575 ASTM, 171 astronomy, 404 asymmetry, 258, 287, 575 atmosphere, 11, 84, 287, 302, 422, 511, 512, 528, 529, 530, 536, 537, 546, 547, 549, 606 atomic force, 9, 73, 101, 105, 197, 282 atomic force microscope (AFM), 9, 11, 12, 13, 14, 15, 73, 74, 78, 79, 80, 81, 84, 87, 104, 105, 194, 195, 196, 197, 222, 253, 254, 282, 286, 315, 448, 449, 450 ATRP, 309, 310, 311, 312, 313, 314, 315, 316, 317, 318, 319, 320, 323, 324, 325, 326, 328 attachment, 119, 181, 218, 372, 416 Au nanoparticles, 33 Au substrate, 10 autocorrelation, 569, 571, 574, 575 Avogadro number, 569

B backscattering, 74, 94, 256 bacteria, 617 bacterial, 578 bacteriostatic, 619 band gap, xiv, 5, 7, 29, 35, 258, 260, 261, 263, 269, 294, 301, 460, 474, 475, 479, 481, 482, 484, 486, 489, 490, 491, 501, 502, 503, 506, 539, 540 bandwidth, 460 barium, 156, 361, 367, 370 barrier, 25, 246, 265, 568, 570, 591 barriers, 11, 165, 263, 510, 518 basement membrane, 110 basic research, 144 batteries, 23, 31 beams, 104, 286, 392 behavior, xv, 148, 260, 263, 277, 278, 283, 285, 286, 288, 289, 312, 324, 346, 351, 355, 373, 393, 394, 395, 403, 456, 465, 474, 480, 493, 525, 526, 540, 543, 544, 570, 580 Beijing, 304, 309, 435, 453, 459, 509 bending, 25, 186, 194, 417, 455, 519, 575, 595 benefits, 97, 376, 450 benign, 42 bias, 11, 12, 14, 29, 75, 373, 395, 511, 512, 513, 514, 518, 522, 523 biaxial, 609, 610, 612 binding, 112, 128, 239, 240, 261, 376, 398, 528, 546, 560, 590 binding energies, 376, 398 binding energy, 239, 240, 261, 528 bioactive compounds, 129 biocompatibility, 109, 115, 119, 123, 128, 140, 328 biocompatible, 109, 114, 123, 128, 129, 144, 239, 319, 619 biodegradability, 128, 132 biodegradable, 109, 114, 129, 133, 142, 148, 149

Index biological processes, xvi, 559 biomacromolecules, 310 biomaterials, xii, 148, 309 biomedical applications, vii, 239 biomolecules, 120, 143 biopolymer, 590, 595 biopolymers, 602, 620 biosensors, 22, 29, 31, 40 biotechnology, 146, 310, 311, 312, 314 bipolar, viii, 4, 41, 42 bismuth, xii, 331, 333, 351, 352, 367, 368 blends, 53 blocks, xiv, 238, 416, 459, 617 blood, 109, 110, 111, 112, 113, 115, 116, 118, 119, 123, 126, 128, 133, 140, 145, 147, 321 blood flow, 111, 126 blood stream, 111, 119, 126 blood supply, 111 blood vessels, 109, 110, 111, 112, 113, 118, 119, 128 bloodstream, 115 blueshift, 263, 264 body weight, 113 Bohr, 117, 476 boiling, 54, 152, 412, 460 Boltzmann constant, 287, 438, 543, 544 bonding, 55, 61, 63, 87, 404, 411, 416, 417, 424, 455, 560 bonds, 74, 165, 282, 310, 394, 399, 400, 401, 402, 412, 417, 560 bone cancer, x, 107, 108, 110 bone graft, 134 bone marrow, 110 boron-doped, 454 bottom-up, vii, viii, xi, 3, 4, 293, 300 bovine, 130 brain, 116 branching, 25, 111, 422, 499, 585 branes, 24 brass, 283 breast cancer, 120, 123, 127, 133 bromine, 280 buffer, ix, 38, 71, 72, 73, 74, 75, 76, 80, 81, 82, 86, 87, 88, 100, 175, 259, 322, 566, 585, 592, 594, 595 building blocks, xiii, 239, 409, 410, 411, 417, 425, 428, 430, 431, 560, 575 bulk crystal, 89, 97, 101 bulk materials, ix, xv, 72, 73, 184, 228, 361, 525, 526 by-products, 156

C C++, 376 cables, 152, 605 cadherin, 110 cadmium, 4, 5, 117, 121 calcination temperature, 335, 336, 344, 345, 349 calcium, 108, 137, 248, 352, 368

623

calibration, 376 cancer, ix, x, 107, 108, 109, 110, 111, 114, 116, 117, 119, 120, 121, 123, 124, 126, 127, 128, 132, 133, 134, 138, 140, 141, 142, 143, 144, 145, 146, 147, 148, 149 cancer cells, ix, 107, 110, 116, 119, 123, 127, 132, 148 cancer treatment, 109, 114, 126, 127, 140 cancerous cells, 108, 110, 126 candidates, xi, xii, 38, 52, 111, 117, 119, 124, 291, 293, 299, 303, 331, 335, 342, 344, 357, 460, 526, 561, 606 capacitance, 19, 23, 33, 348 capacity, 52, 68, 130, 178, 320, 424, 431, 521 capillary, viii, 14, 36, 41, 42, 51, 52, 54, 56, 57, 67, 414, 429, 430 caps, 391 carbide, 173, 333, 334, 335, 338, 339, 343, 353, 361, 605 carbon atoms, 372, 373, 376, 377, 388, 402, 403 carbon film, 397 carbon materials, 379, 390, 392 carbon nanotubes, viii, 4, 9, 22, 24, 34, 41, 102, 238, 267, 298, 299, 315, 319, 320, 372, 381, 390, 391, 393, 397, 404, 606 Carbon nanotubes (CNTs), 33, 34, 298, 315, 316, 317, 318, 319, 320, 321, 322, 323, 324, 372, 373, 397, 400, 403 carbonates, 159, 332 carbonization, 126 carboxyl, 314, 320 carboxyl groups, 314, 320 carboxylic, 320 carcinogenesis, 133 carcinoma, 127, 141, 144, 146 carcinomas, 116, 132 cardiovascular system, 145 carefulness, 347 carrier, xi, 65, 114, 150, 240, 263, 275, 460, 465 Cartesian coordinates, 386 CAS, 506 case study, 365 casting, viii, 4, 355 catalysis, 239, 309, 590 catalyst, viii, xi, 51, 237, 239, 241, 242, 243, 269, 310, 326, 606, 611, 612 catalytic activity, 618, 619 catalytic system, 313 cathode, xv, 17, 31, 42, 247, 250, 252, 253, 267, 397, 509, 510, 512 cation, 342, 349, 393, 484 cavities, 37, 438 cell, 5, 10, 14, 15, 17, 22, 110, 112, 115, 116, 119, 124, 126, 127, 130, 133, 137, 139, 142, 143, 145, 147, 148, 153, 160, 161, 247, 255, 280, 302, 342, 412, 450, 468, 494, 593 cell culture, 139 cell cycle, 126 cell death, 110

624 cell growth, 110 cell membranes, 115 cell surface, 119, 147 cellular phones, 276 centrifugal forces, 334 CH4, xiii, 435, 438, 440, 441, 442, 443, 444, 445, 446, 453, 611 chain transfer, 310 channels, 23, 25, 40, 238, 247, 248, 417 charge coupled device, 73 charged particle, 131 chemical approach, 480 chemical bonds, 121, 380 chemical composition, 238, 412 chemical kinetics, 600, 601 chemical properties, 566 chemical reactions, 245, 362, 369 chemical reactivity, 11, 374 chemical sensing, xi, 294, 590 chemical structures, 312 chemical vapor deposition, 241, 454, 456, 457 chemical vapour deposition, 410 chemicals, 239 chemisorption, 310, 381 chemoprevention, 134 Chemopreventive agents, 142 chemotherapeutic agent, 142 chemotherapy, ix, 107, 123, 124, 127, 128 Chemotherapy, 148 China, 51, 237, 309, 317, 331, 435, 459, 475, 479, 506, 509, 524, 525, 605, 612 chiral, 560, 561, 578, 590, 591, 592, 593, 594, 595 chirality, 561, 590, 591, 594, 595, 596 chitosan, 129, 130, 131 chloride, 27, 31, 245, 248, 313, 323 chloroform, 323 chlorophyll, 560, 563, 578 chopping, 85 chromatography, 311 circular dichroism, 562, 590, 591, 592, 593 circular dichroism (CD), 590 circularly polarized light, 590, 593 circulation, 111, 115, 119, 123, 144 cisplatin, 132, 140 classes, 132, 238, 412, 413 classical, 161, 189, 398, 402, 592 cleaning, 29, 374, 380, 511, 618 cleavage, 145, 450, 611 clinical trials, 128 clustering, 419, 489, 568 clusters, xv, xvi, 10, 11, 140, 253, 277, 375, 376, 377, 378, 388, 390, 397, 399, 400, 401, 404, 411, 505, 525, 540, 551, 559, 565, 567, 568, 570, 571, 572, 573, 584, 611 C-N, 416 CNS, 148 CO2, 140, 343 coagulation, 124, 126

Index coatings, xv, 29, 109, 111, 115, 121, 125, 127, 131, 456, 510, 518, 525, 618 cobalt, 36 coding, 143 coherence, 123, 153, 349, 564, 565, 583, 584, 591 collaboration, 291, 404 collisions, 334, 389, 397, 566, 567, 568 colloidal particles, 144, 429 colors, 461, 464 combustion, 332, 343, 354, 369, 498 commercialization, 18 communication, 108, 140, 562 communities, 372 community, 295, 333 compatibility, 191, 460, 506 compensation, 500 competition, 164, 263, 562 complement, 115, 119 complex systems, vii, 3 complexity, xii, 108, 258, 372, 409, 410, 425, 431, 511 compliance, 172 complications, 126, 174 components, xii, 108, 145, 173, 261, 294, 299, 331, 333, 341, 342, 343, 352, 361, 372, 410, 413, 418, 424, 430, 431, 510, 528, 530, 560, 574 composites, x, 34, 151, 161, 165, 207, 321, 358, 361, 526, 561, 581, 582, 583, 584, 585, 610 composition, xiii, 6, 103, 114, 117, 120, 132, 238, 247, 250, 263, 300, 340, 344, 350, 354, 358, 412, 435, 436, 444, 446, 485, 552, 560, 566, 606 compound semiconductors, 4, 5, 7 compounds, xii, 6, 7, 8, 115, 129, 135, 138, 154, 155, 161, 210, 213, 247, 276, 309, 326, 328, 331, 335, 338, 343, 347, 350, 352, 353, 354, 362, 369, 411, 546, 560 computer simulations, 380, 397 condensation, 245, 492, 505 condensed matter, 560 conductance, 301, 372, 373, 404 conducting polymers, viii, 4, 10, 13, 22, 23, 31, 34, 40 conduction, 263, 264, 265, 277, 278, 279, 280, 281, 282, 287, 301 conductive, 9, 15, 39, 122, 265, 277, 297 conductivity, viii, xi, 15, 30, 41, 51, 53, 66, 67, 169, 170, 240, 241, 275, 278, 279, 437, 510 conductor, 265, 279, 381 configuration, ix, 72, 73, 77, 89, 96, 97, 101, 266, 297, 373, 383, 384, 393, 397, 501, 502, 503, 560, 574, 575, 578 confinement, 18, 117, 257, 258, 466, 532, 591 conjecture, xii, 371, 372, 397 connectivity, 424 conservation, 74 constant load, 455 constant rate, 162 constraints, 239, 574 construction, viii, 4, 21, 152, 238

625

Index consumer electronics, 294 consumption, 313, 344 contamination, 190, 298, 338, 339, 528 contrast agent, 119, 125, 149 control group, 123, 127 convective, 429 conversion, 27, 181, 313, 365 cooling, 159, 161, 165, 169, 224, 228, 230, 335, 339, 376, 390, 391, 542 cooling process, 376 coordination, 8, 282, 399, 410, 480, 560 copolymer, 40, 119, 129, 132, 311, 317 copolymers, viii, 4, 24, 40, 132, 147, 310, 313, 314 copper, vii, x, 3, 9, 10, 20, 39, 53, 126, 127, 151, 153, 168, 250, 251, 262, 278, 414, 609 copper oxide, 168 core-shell, 119, 301, 310, 313, 314, 319, 320, 326, 409, 414, 417, 418, 431, 609 correlation, 143, 372, 522, 569, 570, 572, 574, 575, 576, 577, 580, 590 correlation function, 569, 570, 572, 574, 575 correlations, 566 corrosion, 11, 510 cost-effective, 299, 300, 333, 339 costs, 23, 303 cotton, xvi, 615, 616, 617, 618, 619, 620 coupling, xvi, 14, 127, 258, 335, 340, 345, 347, 352, 559, 560, 562, 564, 575, 581, 582, 585 covalent, 115, 121, 165, 310, 510 covalent bond, 115, 310 covalent bonding, 310 coverage, 9, 28, 35, 134, 135 covering, 283, 377, 412 CPD, 412 CPU, vii, 3 crack, 161, 170, 184, 185, 186, 201, 203, 205, 216, 217, 335, 439, 447, 448, 449, 450 cracking, 168, 184, 185, 186, 335 CRC, 46, 270, 507 critical current density, 154, 159 critical temperature, 554 critical value, 515 crops, 142 cross-linking, 314 cross-sectional, 154, 155, 375, 403, 462, 472, 473, 474, 511, 521, 533, 534, 535, 536, 538, 548, 549, 552 crosstalk, 460 cryogenic, 194, 374, 375, 377, 390 crystal growth, 238, 245, 246, 251, 253, 438, 506, 518 crystal lattice, 494, 517, 520 crystal structure, 153, 161, 277, 278, 280, 288, 290, 348, 349, 416, 460 crystal structures, 277 crystalline solids, 191 crystallinity, 415, 428, 429, 489, 493, 496, 552 crystallites, 41, 352, 353, 413, 414, 415, 416, 494, 518

crystallization, xiv, 238, 246, 255, 351, 354, 356, 479, 480, 488, 493, 497, 505, 506 crystals, xiii, 37, 74, 90, 98, 153, 154, 158, 159, 161, 168, 181, 239, 245, 246, 255, 278, 282, 353, 354, 409, 416, 422, 438, 440, 499, 526, 528, 529, 530, 531, 534, 538, 539, 540, 552 Cu cluster, 10 cubic boron nitride, 510 culture, 137, 139 culture media, 139 current limit, 152, 170 current ratio, 277 cutting tools, 456 CVD, xiii, 241, 244, 435, 436, 437, 438, 439, 440, 441, 443, 445, 447, 449, 451, 452, 453, 454, 455, 456, 457 cycles, 5, 6, 34, 38 cyclic voltammetry, 34, 35, 37, 38 cycling, 170 cyclodextrin, viii, 4, 40 cyclodextrins, 40 cytokine, 126 cytometry, 143 cytoskeleton, 126 cytostatic drugs, 132 cytotoxic, 121, 132 cytotoxicity, 121, 142

D data set, 176, 588 database, 292 death, 123, 124, 126, 149 decane, 566 decay, 467, 498, 504, 571, 572, 574, 586 decomposition, 9, 14, 156, 165, 245, 282, 283, 369, 388, 437, 489, 493, 618 defect formation, 56 defects, xv, 9, 41, 72, 80, 81, 154, 158, 161, 166, 210, 211, 224, 227, 228, 229, 250, 255, 258, 288, 290, 354, 362, 391, 403, 416, 417, 438, 460, 471, 503, 504, 509, 609 defense, 133 definition, 4, 16, 178, 345, 410 deformation, x, 56, 151, 161, 165, 171, 173, 174, 177, 178, 181, 184, 187, 188, 189, 190, 191, 192, 193, 211, 216, 217, 218, 221, 373, 397, 400, 401, 402, 518 degenerate, 90, 373 degradation, 12, 33, 114, 132, 140, 378, 618 degrading, 85 dehydration, 245 delivery, x, xvi, 85, 107, 114, 115, 116, 120, 140, 141, 142, 143, 144, 145, 146, 147, 148, 149, 437 delocalization, xvi, 559, 560, 561, 562, 563, 565, 583 demand, 173, 510 dendrites, 414 density functional theory, 103 density values, 154

626 depolarization, ix, 72, 73, 89, 95, 96, 97, 98, 99, 100, 101, 564, 565, 566, 574 deposition rate, 6, 438, 441, 510 deposits, viii, 4, 5, 6, 7, 8, 14, 27, 42 derivatives, 13, 30, 560, 578 desorption, 66, 128, 282, 302, 378, 380, 395, 438 destruction, 126, 566 detection, 22, 24, 31, 40, 53, 65, 73, 76, 77, 82, 89, 93, 98, 103, 108, 110, 124, 300, 327, 370, 439, 441, 512, 595 deviation, 58, 62, 63, 80, 181, 522, 593 DFT, 103 diamond films, xiii, 435, 439, 443, 446, 447, 450, 451, 454, 455, 456 diamond grains, 440, 443 diamonds, 174, 186, 455, 456, 577 dielectric constant, 99, 335, 340, 342, 344, 345, 346, 347, 348, 357, 360, 363 differential scanning, 282 differential scanning calorimetry, 282 diffraction, ix, xi, xiv, 71, 72, 73, 76, 77, 80, 81, 87, 102, 237, 287, 288, 290, 327, 413, 417, 442, 468, 488, 496, 517, 518, 520, 531, 532, 533, 537, 552, 605, 606 diffusion, 17, 110, 114, 129, 145, 153, 156, 160, 161, 164, 167, 226, 228, 242, 246, 302, 345, 353, 417, 430, 431, 437, 465, 466, 510, 567, 572, 574, 575, 576, 578, 579 diffusivity, 438 dimensionality, 238, 239, 277 dimerization, 381 dimethacrylate, 326 diodes, 6, 7, 27, 239, 258, 301, 546 dipole, 99, 146, 240, 386, 504, 564, 578, 582, 583 dipole moment, 240, 564, 578, 582, 583 diseases, 108 dislocation, 157, 161, 162, 163, 165, 166, 167, 190, 191, 207, 211, 213, 451, 460, 518, 519 dislocations, 154, 157, 161, 162, 165, 167, 171, 190, 191, 210, 221, 354 disorder, 255, 474, 562 dispersion, 126, 147, 297, 310, 312, 352 displacement, 8, 171, 173, 175, 176, 179, 181, 189, 197, 218, 342 disposition, 560 dissociation, 162, 377, 388, 399, 400 dissolved oxygen, 27, 28 distilled water, 53 distortions, 354 disulfide, 314, 315 divergence, 578 DNA, 40, 110, 142, 147, 312, 560, 561, 578, 590 domain structure, 285, 544 dominance, 190 donor, 262, 263, 504, 617 donors, 262, 263 dopant, 31, 32, 74, 277, 341, 484, 493, 502 dopants, 6, 38, 240, 277, 335, 489

Index doped, xiv, 74, 81, 239, 262, 277, 303, 323, 341, 479, 480, 481, 482, 483, 488, 489, 490, 491, 493, 494, 499, 501, 502, 503, 505, 506 doping, xi, xiv, 6, 31, 73, 153, 238, 241, 265, 275, 277, 281, 291, 372, 394, 460, 479, 493 dosage, 526 drainage, 111, 112 drug carriers, 112, 120, 145, 149 drug delivery, vii, 108, 111, 114, 115, 122, 124, 128, 141, 143, 145, 146, 147, 148, 312, 313, 314, 560, 615, 619 drug delivery systems, 108 drug discovery, vii, 146 drug release, 128, 132, 133, 140 drug targets, 113 drug use, 132, 133 drugs, ix, x, 107, 109, 111, 112, 113, 114, 115, 116, 120, 122, 124, 127, 128, 129, 132, 133, 142, 143, 145, 313, 314, 560, 619 drying, 148, 412, 413, 414, 415, 417, 418, 419, 421, 422, 423, 425, 426, 427, 428, 429, 430, 431, 439 drying time, 412, 414, 415, 417, 421, 422, 423, 425, 428, 431 DSC, 282, 287 DTA curve, 338, 339 durability, 616 duration, 18, 19, 20, 21, 282, 287, 291, 342, 352, 375, 412, 461, 462, 463, 471 dyes, 31, 117, 118, 119 dynamic scaling, 569, 571, 572, 574 dynamical properties, 561

E elastic constants, 179 elastic deformation, x, 151, 171, 191, 211 elasticity, 404 electric arc, 372, 437 electric charge, 287 electric conductivity, 53 electric current, 42, 170, 246, 278 electric potential, 131 electric power, 170 electrical characterization, 295 electrical properties, x, xi, xv, 117, 152, 277, 278, 281, 283, 285, 293, 294, 299, 300, 335, 342, 345, 347, 348, 361, 362, 364, 371 electrical resistance, 300 Electroanalysis, 49 electrochemical deposition, 15, 23, 27, 28, 36, 37, 38, 40, 245, 247, 249 electrochemical reaction, 10, 14, 18, 20, 42, 248 electrochemistry, viii, 4, 15, 16, 18, 37, 40, 41, 42, 248 electrocrystallization, 8 electrodeposition, vii, 3, 5, 6, 7, 8, 9, 17, 25, 26, 27, 31, 34, 37, 39, 42, 246, 247, 248

627

Index electrodes, 4, 8, 18, 19, 20, 23, 24, 33, 34, 40, 41, 42, 53, 127, 247, 250, 268, 296, 299, 344, 361, 372, 384, 388, 389 electroluminescence, 29, 276 electrolysis, 9 electrolyte, 17, 18, 19, 20, 21, 22, 23, 24, 27, 28, 29, 32, 33, 41, 42, 246, 247, 248, 249, 250, 252 electromagnetic, xiv, 97, 122, 140, 404, 459, 460, 475, 581 electromagnetic waves, 140 electromigration, 373 electron beam, 9, 265, 286, 296, 298, 465 electron beam lithography, 9, 265, 465 electron density, 460 electron diffraction, 250, 254, 520, 549 electron microscopy, xi, xiv, xv, 32, 114, 135, 215, 237, 242, 321, 372, 373, 412, 413, 459, 526, 527, 533, 535, 605, 606 electron paramagnetic resonance, 261 electronic circuits, 372 electronic materials, 291 electronic structure, 5, 142, 280, 286, 326, 372, 394, 403, 456, 460 electronic systems, 294, 300 electrons, 38, 72, 122, 153, 260, 263, 264, 265, 266, 278, 279, 282, 288, 294, 297, 298, 373, 394, 540, 564 electrospinning, viii, 51, 52, 53, 54, 55, 57, 58, 61, 66, 67, 68 electrostatic force, 56, 255, 315 electrostatic interactions, 410 ELS, 552 emission source, 264 emitters, xii, 9, 267, 269, 371, 372, 373, 383, 384, 385, 388, 394, 397 Empedocles, 270 emulsification, 130, 145 emulsifier, 130 enantiomers, 591 encapsulated, 114, 115, 116, 119, 121, 142 encapsulation, 109, 116, 121, 130, 133, 141, 144 endoplasmic reticulum, 148 endothelial cells, 110, 111, 112, 119 endothermic, 339 energy transfer, 141, 480, 505, 561 enlargement, 515 entertainment, 16 entrapment, 130, 132, 156, 157 environment, 20, 114, 126, 133, 175, 312, 335, 350, 411, 425, 429, 431, 481, 489, 498, 499, 500, 501, 503, 506, 523, 565, 579 environmental conditions, 121, 332 environmental factors, 143 environmental impact, xvi, 615 environmental stimuli, 315 enzyme immobilization, 31 enzymes, 114, 115, 619 epitaxial growth, 243 epitaxy, viii, 4, 241

epithelial cell, 119 epithelial cells, 119 epoxy, 392, 402, 403 EPR, 111, 112, 115, 125, 261, 262, 263, 264 equality, 588 equilibrium, 52, 88, 160, 394, 399, 400, 454, 579 equilibrium state, 400 equipment, 42, 172, 173, 276, 295, 296, 334, 363, 410, 512 Erk, 292 erosion, 510 esterification, 316, 323 estrogen, 145 etching, 11, 16, 17, 25, 28, 30, 104, 134, 226, 311, 375, 390, 438, 443, 461, 471, 511, 512, 528, 530 ethanol, 11, 17, 29, 53, 54, 55, 56, 59, 439, 492, 493 ethyl alcohol, 226 ethylene, 30, 31, 111, 145, 148, 314, 326 ethylene glycol, 30, 31, 111, 132, 145, 326 ethylene oxide, 148 europium, 481, 482 evolution, 145, 198, 225, 252, 337, 338, 344, 347, 350, 351, 362, 367, 369, 422, 423, 424, 431, 566, 572, 574, 579, 586, 587 exciton, xi, 117, 237, 239, 258, 261, 499, 546, 547, 562, 563, 564, 565, 575, 584, 585 exothermic peaks, 493 experimental condition, 28, 39, 77, 310, 497, 566, 569, 590 exposure, 93, 113, 121, 124, 147, 297, 300, 301 extinction, 109, 118, 563, 573, 584, 585, 594, 595, 596 extraction, xiv, 459, 465, 467, 469, 470, 474, 475

F fabricate, xiv, xvi, 9, 14, 22, 23, 24, 25, 26, 27, 42, 73, 158, 265, 277, 297, 299, 303, 348, 355, 359, 360, 410, 438, 440, 451, 453, 459, 460, 525, 546 face-to-face interaction, 560 failure, 399, 402, 403, 461 family, 42, 110, 299, 333, 347, 350, 351, 362, 386 fault current limit, 152, 170 faults, 154, 161, 162, 163, 402 FCC, 463 FCL, 170 feedback, 15 feeding, 20 FEM, 88 Fermi, 465 ferroelectrics, 331, 332, 333, 342, 343, 347, 348, 352, 362 ferromagnetic, xv, 525, 526, 543, 544, 555 Ferromagnetic, 526 ferromagnetism, xvi, 278, 525 FFT, 415, 416 FIB, viii, 4, 20, 42, 85, 295, 297, 298, 303, 375 fiber, 41, 42, 52, 55, 56, 61, 62, 65, 66, 77, 78, 84, 283, 377, 390, 391, 439

628 fibers, viii, xii, 51, 52, 53, 54, 55, 56, 57, 58, 59, 60, 61, 62, 63, 64, 66, 67, 68, 371, 373, 375, 377, 390, 401 fibronectin, 137, 144 field emission scanning electron microscopy, 412 field-emission, 9, 391, 392 filament, 265, 381, 393 fillers, 311 film thickness, 7, 38, 444, 489, 512, 513 filters, viii, xv, 51, 525 filtration, 52, 119 finite element method, 88 fire hazard, 615 flame, 455, 615 flat-panel, xi, 237, 264 flexibility, 27, 299, 312 flexural strength, 453 flight, viii, 51, 54, 55, 56, 57, 375, 388 floating, 429 flow, 41, 65, 111, 126, 143, 153, 156, 188, 190, 191, 221, 224, 228, 278, 425, 437, 440, 454, 461, 492, 511 flow rate, 65, 425, 440, 461 fluctuations, ix, 71, 395, 397 fluid, 144, 425 fluorescence, 104, 113, 117, 119, 145, 326, 493, 497, 504, 583, 588, 589 fluorescence decay, 504 fluoride, 29, 30, 31 fluorinated, 30 fluorophores, 117 flux pinning, 154, 158 focused ion beam, viii, 4, 20, 42, 85, 297 focusing, ix, xiv, 107, 140, 374, 383, 404, 411, 459 foils, 30 folate, 113, 116, 145, 146, 147, 148 folic acid, 113, 143 Ford, 316, 317, 320, 329, 330 Fourier, 416, 569, 595 fractal cluster, 568, 571 fractal dimension, 567, 570, 575, 580, 584 fractal structure, 567, 568, 569, 575, 579, 580, 585, 586 fractality, 584 fractal-like, 73 fractals, 567, 570, 572, 575, 576, 581, 582, 583, 584, 586, 591 fracture, 170, 173, 174, 184, 185, 186, 190, 194, 195, 196, 197, 199, 201, 203, 204, 205, 213, 215, 216, 217, 224, 228, 373, 397, 403, 436, 439, 443, 447, 450, 453, 454, 455, 456 fracture processes, 190 fragility, 276 fragmentation, 142, 353, 377, 399 free energy, 4, 246, 251, 353, 354, 356, 429 free radical, 323 freedom, 17, 297 friction, xv, 172, 175, 190, 334, 509, 514, 515, 516 FTIR, 140, 412, 416

Index fullerene, 50, 391 functionalization, 34, 310, 316, 318, 324, 328, 372 fungal, 616 fungi, 617 Fusarium, 616 fusion, 145, 421, 425, 518 FWHM, 250, 258, 261, 264, 448, 450, 498, 532

G GaAs, 7, 103 gadolinium, 125 gallium, 161, 239, 243, 276 gas, 7, 9, 27, 53, 64, 65, 66, 68, 238, 240, 247, 294, 300, 302, 303, 374, 375, 379, 383, 389, 390, 391, 392, 394, 395, 437, 438, 439, 443, 454, 455, 456, 461, 511, 512, 514, 522, 523, 606, 611 gas diffusion, 302 gas phase, 238, 394, 437, 439, 443, 454 gas sensors, 240, 300, 303 gases, 24, 65, 66, 239, 300, 302, 303, 374, 390, 511, 512, 611 gasification, 438 gastric, 133, 149 Gaussian, 87, 259, 260, 469, 528, 530, 570 gelatin, 129 gelation, 31, 567, 569 gels, 336, 363 gene, vii, 143, 144, 145, 147 gene therapy, vii, 143, 147 generation, xi, 7, 10, 40, 118, 171, 184, 216, 227, 293, 294, 299, 410, 522, 583, 619 Georgia, 140 germanium, 72, 73, 239, 244, 460 Germany, 293, 334, 433, 476 GFP, 118 Gibbs, 596, 597, 598, 601 glass, 6, 7, 28, 32, 34, 37, 41, 42, 246, 461, 480, 488, 489, 497, 546, 589, 618 glass transition, 461 glass transition temperature, 461 glasses, 354, 355 glow discharge, 522 glutathione, 135 glycerol, 318 gold, 9, 11, 14, 27, 33, 35, 36, 37, 38, 39, 40, 41, 42, 79, 104, 109, 121, 122, 123, 124, 242, 244, 247, 249, 250, 251, 256, 259, 261, 310, 314, 315, 372 gold nanoparticles, 121, 242, 250, 314, 315 grades, 454 grafting, 311, 324 grain boundaries, 154, 158, 161, 277, 336, 352 grains, xiii, xv, 73, 158, 159, 190, 208, 250, 253, 277, 332, 335, 336, 338, 353, 354, 355, 435, 440, 443, 444, 446, 448, 453, 509, 510, 517, 518, 520, 521, 549 graphene sheet, 381, 398, 399, 401, 402

Index graphite, 7, 8, 9, 10, 13, 41, 139, 181, 372, 373, 376, 377, 378, 379, 380, 390, 391, 392, 395, 398, 399, 400, 401, 402, 403, 404, 411, 412, 418, 610 gravity, 56, 334 groups, xi, 11, 12, 24, 27, 112, 121, 123, 129, 173, 241, 245, 275, 278, 279, 311, 312, 314, 315, 316, 317, 320, 321, 337, 349, 402, 429, 499, 559, 570, 571, 575, 585, 611 growth factor, 110 growth mechanism, xi, 24, 153, 237, 238, 241, 247, 253, 269, 480, 505, 510, 568, 574, 575 growth rate, xiii, 110, 153, 156, 157, 168, 251, 435, 436, 437, 438, 441, 443, 446, 451, 453, 457 growth temperature, 242, 447, 448, 450 growth time, 252 guidance, 143 guidelines, 294 gyration radius, 575

H H2, xiii, 21, 25, 376, 435, 438, 439, 440, 441, 442, 443, 444, 445, 446, 453, 530, 536, 537, 611 half-life, 115 halogen, 278, 279, 538 haloperidol, 141 Halothane, 147 Hamiltonian, 467, 562 handedness, 590, 593 handling, 158, 294 harvesting, 27, 239, 294, 301, 561 health, 108 healthcare, 617 heart disease, 149 heat, 124, 126, 127, 159, 284, 332, 354, 401, 453, 482, 486, 497, 505, 541 heating, 31, 41, 127, 143, 156, 161, 162, 224, 283, 284, 339, 363, 373, 402, 404, 437, 461, 471, 472, 492, 578 heating rate, 492 heavy metal, 120, 121 heavy metals, 120 height, 10, 13, 15, 27, 73, 80, 211, 224, 229, 230, 265, 266, 375, 377, 381, 438, 498 helium, 374, 375, 376, 378, 379, 380, 383, 390, 391, 392, 393, 394 helix, 578, 595 hematological, 133 hemisphere, xiv, 374, 381, 384, 386, 387, 459, 460, 461, 463, 470, 471, 473 hemoglobin, 560 hepatic injury, 149 hepatocyte, 121 hepatocytes, 148 hepatotoxicity, 147 HER2, 112, 119, 123, 124 hES, 569, 571 heterogeneity, 80, 81, 150, 190, 342 heterogeneous, x, 152, 156, 181, 311, 422, 585

629

heterostructures, 102, 320, 510, 605 hexafluorophosphate, 32 hexane, 126 high pressure, 153, 353, 411 high resolution, xii, 38, 39, 168, 371, 472, 498, 499, 537, 548, 549, 550, 552 high tech, 510 high temperature, xii, 152, 156, 159, 319, 331, 332, 335, 343, 344, 345, 346, 347, 348, 352, 353, 354, 355, 358, 373, 401, 411, 464, 471, 480, 500, 505, 526, 611 high-speed, 510, 512, 515, 516, 523 high-tech, 294 Hoechst, 120 Holland, 49, 273 homogeneity, ix, 71, 126, 210, 277, 333, 348, 392, 567 homogenous, ix, 71, 72, 159, 190, 349, 428 homopolymers, 314 Honda, 555 Hong Kong, 509, 524 hormone, 133 host, xiv, xv, 73, 479, 480, 481, 484, 489, 490, 491, 498, 501, 502, 505, 506, 525, 526, 581 House, 273 HRP, 31 HRTEM, 168, 250, 321, 323, 327, 416, 425, 426, 428, 429, 482, 483, 496, 512, 517, 518, 520, 521, 538, 606, 609, 610 human, vii, 108, 109, 119, 124, 142, 146, 150 humans, 128, 133 humidity, viii, 11, 13, 15, 51, 53, 54, 55, 56, 57, 59, 60, 61, 62, 63, 65, 66, 68, 486, 487 hybrid, xii, 31, 33, 299, 309, 310, 311, 312, 314, 318, 321, 326, 328, 363 hybridization, 122, 147, 355 hybrids, 314, 319, 321, 322, 326 hydro, 12, 109, 111, 114, 115, 119, 120, 129, 133, 318, 326, 390, 392, 460 hydrocarbons, 390, 392, 443, 611 hydrochloric acid, 580, 585 hydrodynamic, 127, 312, 313, 575 hydrofluoric acid, 16 hydrogen, vii, xiii, 9, 11, 12, 31, 280, 376, 383, 410, 435, 437, 443, 446, 453, 456, 560 hydrogen bonds, 410 hydrogen gas, 9 hydrogen peroxide, 12, 31 hydrolysis, 245, 246, 319, 323, 492, 505 hydrolyzed, 245 hydrophilic, 12, 109, 111, 114, 115, 119, 120, 129, 133, 318, 326, 460 hydrophilicity, 132 hydrophobic, 36, 111, 114, 115, 120, 313, 318, 326, 410, 560 hydrophobic interactions, 115, 560 hydrophobic properties, 115 hydrophobicity, 111

630

Index

hydrothermal, 244, 245, 332, 339, 350, 367, 410, 490, 491, 496 hydrothermal process, 367 hydroxide, 7, 8, 245, 310, 326, 327, 332, 505 hydroxyl, 40, 129, 320, 402, 505 hydroxyl groups, 320, 402 hyperbranched polymers, 328 hypersensitive, 503 hyperthermia, 109, 124, 126, 127, 140, 141, 142, 143, 144, 146, 312 hypothesis, 288, 569, 575 hypoxic, 111 hysteresis, 339, 359, 360, 542, 544, 545 hysteresis loop, 359, 360, 542, 544, 545

I IBM, vii, 3, 43, 476 ICD, 590, 591 identification, 79, 118, 143, 176, 560, 586 IgG, 120 illumination, 17, 77, 85, 91, 93, 94, 95, 96, 97, 99, 100, 102, 119, 285, 287, 465 imaging techniques, 124 immersion, 73 immobilization, 143 immunodeficiency, 124 immunoglobulin, 115 immunopathology, 143 immunotherapy, 142 implants, 108, 146 imprinting, 590, 594 impurities, 211, 338, 339 in situ, xii, 7, 115, 148, 315, 316, 371, 373 in vitro, 113, 137, 138, 140 in vivo, 117, 119, 124, 127, 140, 143, 144, 145, 146, 147, 149 inactivation, 121 incidence, 90, 91, 92, 93, 94, 95, 96, 100, 133, 134, 142 inclusion, 153, 590 independence, 569, 575 indexing, 535 indication, 131, 346, 394, 523 indicators, 593 indium, 239 induction, 125, 154, 561, 570, 581, 588 induction period, 570 induction time, 581, 588 industrial, xi, 109, 116, 267, 270, 275, 276, 410, 510 industry, 190 inert, 7, 282, 372, 391, 392, 394 inert liquid, 372 infancy, 411 infinite, 178 inflammation, 148 inflammatory, 145 information and communication technology, 108 infrared, 117, 121, 122, 255, 282, 412, 416

infrared spectroscopy, 412, 416 infrastructure, 28 inherited, 325, 440, 441 inhibitor, 245 inhomogeneities, 263, 588 inhomogeneity, 211 initiation, 40, 315 injection, 28, 119, 124 Innovation, 506 inorganic, viii, xi, xii, 4, 9, 25, 33, 42, 117, 120, 275, 276, 277, 278, 295, 309, 310, 311, 312, 315, 326, 328, 411 InP, 7, 103 insertion, 98, 108 insight, xvi, 378, 506, 559 instabilities, 372, 394, 395, 396, 438, 454 instability, xiv, 52, 54, 57, 58, 67, 121, 283, 429, 436, 437, 438, 451, 453 instinct, 436 instruments, 173, 247, 300, 467 insulation, 18, 279 insulators, xv, xvi, 291, 525, 526 integrated circuits, vii, 294 integration, vii, 3, 23, 79, 301, 302, 303 integrity, 115 Intel, vii, 3, 72, 304, 476 intercalation, 560 interdependence, 394 interface, 40, 73, 74, 122, 153, 156, 157, 159, 161, 162, 163, 165, 168, 175, 185, 210, 217, 229, 238, 247, 250, 269, 286, 290, 299, 300, 301, 303, 314, 429, 437, 438, 439, 465, 466, 468, 489, 520 interface energy, 250 interference, 176, 190, 527, 542, 569 intermetallic nanoparticles, xvi, 525, 526 intermetallics, 332 intermolecular, xvi, 410, 559, 561, 562, 569, 575, 585, 596 intermolecular interactions, xvi, 559 internal environment, 114 internalization, 114 internalizing, 112, 114 interpretation, 288, 404, 591 interstitial, 36, 269 interstitials, 258, 551 interval, 118, 258, 283, 375, 618 intestine, 133 intravascular, 148 intrinsic, xi, 152, 167, 211, 237, 238, 239, 242, 258, 262, 277, 293, 295, 297, 303, 332, 388, 411, 450, 460, 480, 499, 504, 510, 523, 561 inversion, 498, 596 Investigations, 146, 213 ion beam, 24, 276, 297, 298, 305, 378, 389, 392 ion bombardment, xv, 297, 395, 396, 510, 512, 514, 521, 523 ion implantation, xvi, 525, 526, 527, 528, 532, 533, 534, 536, 537, 538, 539, 540, 541, 546, 548, 549, 550, 551, 552

Index ionic, xi, xiv, 8, 17, 32, 34, 129, 153, 165, 275, 443, 479, 480, 494, 560, 561, 570, 571, 574, 575, 594 ionic liquids, 8, 32, 34 ionization, 282, 374, 376, 378, 389, 392, 394, 510, 511, 522 IOP, 326, 327, 415, 416 iron, 17, 18, 36, 125, 127, 128, 149, 323, 343, 344, 345, 362, 606, 611 irradiation, xi, 123, 275, 281, 282, 283, 284, 286, 287, 288, 289, 290, 291, 296, 397, 413, 549, 578 irritation, 619 island, 21, 74, 75, 76, 77, 380 isomers, 373 isostatic pressing, 161 isothermal, 158 isotropic, 212, 564, 569, 574, 579, 580, 594 Italy, 559 ITO, 6, 7, 28, 32, 34 ITT, x, 151, 152, 171, 172, 173, 174, 175, 176 I-V curves, 287 IVH, 564

J Japanese, 68, 102, 103 Jc, 154 Jordan, 109, 124, 126, 127, 144 Jung, 45, 235, 271, 507, 556

K KBr, 412 kinases, 126 kinetic energy, 376, 378, 383, 389, 400, 510 kinetics, x, 41, 152, 156, 171, 242, 247, 313, 351, 400, 411, 561, 567, 571, 572, 578, 579, 581, 586, 588, 600, 601 KOH, 17, 25, 27, 323, 375 Korean, 102, 145

L labeling, 118, 149 lactic acid, 129 lamellar, 484 Langmuir, 45, 47, 48, 49, 271, 433, 600 Langmuir-Blodgett, 239 lanthanide, xiv, 326, 479 lanthanum, 310, 326, 327, 341, 342 large-scale, 28, 247, 276, 295, 303, 460 laser, 73, 75, 77, 78, 79, 83, 84, 85, 87, 89, 91, 93, 105, 239, 256, 257, 258, 259, 282, 372, 376, 378, 409, 411, 412, 413, 414, 415, 418, 421, 422, 423, 424, 426, 427, 428, 448, 454, 461, 464, 465, 480, 489, 493, 503, 504, 546, 550, 606 laser ablation, xiii, 409, 410, 411, 413, 431, 606 lasers, 7, 127, 239, 240, 546 latency, 460 latex, viii, 4, 24, 35, 36, 37 lattice parameters, 494

631

lattices, 416, 426, 506, 510 laundering, 616 law, 6, 92, 100, 179, 190, 191, 396, 569, 570, 573, 574, 579, 583, 584, 585, 586, 587, 594 layer-by-layer growth, 3, 5 lead, xii, 130, 167, 184, 226, 247, 253, 254, 258, 277, 278, 302, 326, 331, 332, 333, 338, 342, 343, 344, 345, 348, 352, 353, 354, 362, 363, 364, 365, 366, 369, 403, 494, 515, 574 leakage, 109, 116, 133 LED, xiv, 28, 29, 239, 258, 459, 465, 470, 474, 475 Leibniz, 293, 303 lens, 73, 77, 78, 85, 93, 95, 97, 98, 412, 439, 464 LEO, 439 lesions, 110 leukemia, 127 lifespan, 116 lifetime, 110, 117, 467, 475, 504 ligand, 112, 113, 116, 120, 489, 498, 584, 585 ligands, 116, 119, 619 light beam, 291 light emitting diode, 7, 27 light scattering, 562, 570, 572, 578, 579, 584, 588, 599, 602 limitation, 128, 129, 133, 224, 295, 303, 460 limitations, xi, 7, 116, 124, 243, 293 linear, xii, 157, 179, 191, 218, 283, 324, 371, 372, 373, 383, 388, 389, 394, 395, 396, 397, 398, 399, 400, 401, 402, 403, 404, 438, 563, 567, 581, 583, 588, 591 linear regression, 563 linkage, 417 links, 402 lipid, 114, 115, 147 lipids, 114, 125, 143 lipopolysaccharide, 149 liposome, x, 107, 114, 115, 116, 140, 143 liposomes, 109, 113, 114, 115, 116, 140, 141, 142, 143, 144, 145, 147, 148, 149 liquid film, 429 liquid helium, 390 liquid nitrogen, 152, 171, 283 liquid phase, xiii, 153, 156, 164, 409 liquids, 8, 32, 34, 52, 411 lithium, xii, 31, 278, 331, 333, 368, 453 Lithium, 352 lithography, vii, viii, xiv, 3, 4, 10, 11, 13, 14, 15, 16, 17, 18, 24, 239, 242, 276, 277, 295, 296, 297, 298, 303, 431, 459 liver, 111, 113, 121, 133, 146 liver cells, 146 liver disease, 146 living radical polymerization, xii, 309, 310 L-lactide, 40 localization, 79, 132, 263, 269, 582, 583 location, 9, 10, 15, 127, 176, 379, 413, 482, 500, 520 London, 46, 270, 273, 306, 405, 406, 556, 599 long-term, 300, 302, 395, 411 low temperatures, 350, 361, 362, 375, 377

632

Index

285, 339, 373, 378, 454, 455, 464, 467, 473, 527, 531, 532, 534, 540, 542, 595 measures, 172, 569 mechanical behavior, 519 mechanical properties, x, 134, 151, 152, 170, 171, 172, 173, 175, 176, 177, 181, 189, 194, 196, 201, 203, 206, 212, 225, 238, 276, 277, 340, 372, 398, 403, 404, 455, 456, 510, 514, 523 mechanical stress, 167, 169, 402 mechanical testing, 171, 402, 403 media, 24, 92, 139, 245, 314, 333, 334, 335, 339, 351, 352, 361, 411 median, 184, 185 medicine, 108, 147, 310 melanin, 578 M melanoma, 148 melt, 52, 158, 159, 161, 165, 168, 207, 224, 225, 353, machines, vii, 3, 333, 402 354 macromolecules, 144, 324, 590 melting, 127, 156, 159, 161, 265, 344, 355, 461, 510, macrophages, 111, 119, 146 611 maghemite, 125 melting temperature, 344 magnesium, 129, 343, 365, 366, 541 melts, 52 magnetic field, 124, 125, 127, 128, 144, 152, 159, membranes, viii, 4, 24, 27, 32, 40, 65, 66, 141, 311 322, 335, 345, 542, 544, 553 memory, 33, 345, 351, 561, 579, 590, 594 magnetic fluids, 313 MEMS, 21, 46, 190, 360, 369, 370, 436 magnetic materials, 37 men, 108 magnetic moment, 125, 544 mercury, 117, 493 magnetic particles, 128 mesoporous materials, 143 magnetic properties, xi, 109, 124, 127, 152, 239, 275, mesoscopic, 224, 277, 380, 385, 388, 422, 561, 562, 321, 326, 344, 345, 542 566, 575, 579, 590 magnetic resonance, 124, 312 metabolic, 111 magnetic resonance imaging, 124 metabolism, 133, 148, 615 magnetism, xi, 275, 278 metal ions, 7, 8, 15, 245, 247, 315, 320, 560 magnetite, 125, 127, 313, 314 metal nanoparticles, 8, 33, 320, 321 magnetization, 127, 542, 543, 544, 545, 553, 554 metal oxide, xi, 4, 8, 9, 293, 294, 295, 298, 300, 302, magnetizations, 543 315, 411, 619 magnetron, 514 metal oxides, 4, 9, 411 magnetron sputtering, 514 metal salts, 14, 615 magnets, 109, 124, 127, 128, 170 metalloids, 133 mainstream, vii metalloporphyrins, 560 management, 108, 117, 453, 620 metallurgy, 391 manifold, 300, 591 metal-oxide-semiconductor, 29, 101, 460 manipulation, 5, 10, 172, 294 metals, viii, 4, 7, 8, 9, 10, 11, 14, 27, 35, 42, 73, 85, manufacturer, 183 109, 121, 133, 135, 147, 149, 172, 176, 190, 191, manufacturing, vii, 3, 16, 17, 291 211, 247, 278, 291, 316, 355, 376, 380, 411, 560 mapping, 81, 144, 552, 553 metal-semiconductor, 301 market, 16, 294 metastases, 116 Markov, 598 metastasis, 110, 127 mask, 9, 286, 296, 297, 303, 474 metastasize, 110 mass transfer, 17, 153 methane, xiii, 435, 443, 446, 453 material sciences, 596 methyl group, 611 material surface, 10 methyl groups, 611 materials science, 5, 310, 560 methyl methacrylate, 311, 313, 318, 326 MBE, 241, 244 Mg2+, 342 MBP, 347 mica, 181 mean-field theory, 568 micelles, 121, 142 measurement, ix, 66, 71, 72, 77, 80, 95, 97, 101, 172, microcracking, 167, 335 176, 179, 190, 198, 250, 253, 266, 282, 283, 284, microelectrodes, 20, 295, 297, 298 microelectronics, 29, 303, 305, 355 low-temperature, viii, xii, 4, 42, 245, 362, 371, 373, 376, 377, 378, 379, 380, 390, 404 luminescence, xiv, 143, 258, 264, 461, 465, 474, 479, 480, 482, 489, 497, 498, 503, 504, 506 lung, 108, 116, 133, 150 lung cancer, 133, 150 lungs, 108, 111, 133 lying, 379, 422, 468 lymph node, 144 lymphatic, 111, 112, 620 lymphatic system, 112 lymphoma, 133 lymphomas, 132 lysosomes, 114

633

Index microemulsion, 332, 350, 564, 565, 566, 591 microemulsions, 591 microenvironment, 118, 144 microfabrication, 17 micrometer, 108 microorganism, 29 microorganisms, 615, 618 microparticles, 141 microphotographs, 395 microscope, xii, xiv, xv, 10, 32, 41, 73, 84, 101, 102, 113, 197, 201, 226, 297, 371, 373, 374, 375, 378, 381, 383, 387, 390, 391, 404, 448, 451, 462, 464, 479, 509, 512, 608 microscopy, xi, xii, xiii, xiv, xv, 9, 41, 54, 72, 74, 77, 102, 103, 104, 105, 118, 131, 135, 164, 167, 226, 237, 321, 371, 372, 373, 374, 380, 383, 390, 404, 413, 435, 459, 527, 533, 534, 535, 584, 589, 605 microspheres, 140, 148, 254, 411, 490, 491 microstructure, xv, 158, 161, 207, 345, 359, 360, 362, 363, 364, 369, 460, 465, 467, 509, 516, 518, 521, 522, 523, 539 microstructures, xii, 16, 17, 18, 20, 181, 252, 331, 332, 453, 454, 466 microwave, 127, 438, 440, 443, 451, 455 microwave radiation, 127 migration, 396, 438, 549 minerals, 108 miniaturization, 17, 294, 355, 371 mining, 403 misfit dislocations, 157 misleading, 277, 563 mitochondrial, 121 mixing, 135, 276, 277, 333, 354, 578, 579, 580, 581, 587, 596 MLC, 349 MMA, 311, 313 mobility, 72, 165, 240, 425, 429, 431, 521 MOCVD, 244 modeling, 41, 303, 398 models, xvi, 153, 180, 212, 362, 372, 386, 468, 559 modulation, 24, 37, 460 modulus, x, 151, 165, 171, 172, 173, 174, 175, 176, 177, 178, 179, 181, 184, 186, 189, 194, 196, 197, 198, 199, 206, 207, 208, 209, 210, 212, 403, 404 moieties, 311, 560 moisture, 32, 300, 302, 303, 332, 615 molar ratios, 132, 337 molar volume, 438 molecular beam, 241 molecular beam epitaxy, 241 molecular biology, 410 molecular dynamics, 374, 398, 400, 567 molecular mass, 574 molecular mechanisms, 596 molecular structure, 278, 430 molecular weight, 130, 132, 313, 569 molecular weight distribution, 313 molybdenum, 9, 438 momentum, 74, 389

monochromator, 493 monoclonal, 119, 120 monoclonal antibody, 119 monolayers, viii, 4, 5, 6, 8, 11, 12, 14, 24, 38, 124, 311, 391, 462 monolithic, 171, 186 monomer, 21, 32, 33, 40, 129, 132, 313, 314, 316, 562, 563, 568, 571, 572, 573, 575, 581, 582, 583, 584, 585, 586, 588, 591 monomeric, 572, 573, 589 Monte-Carlo, 10 Moon, 69, 231, 369, 507 morphological, 144, 431, 437, 438, 454, 456 mortality, 133, 142 mortality rate, 133 motion, 171, 191, 333, 374, 384, 389, 425, 430, 431, 561, 575, 590 motives, 372 motors, xii, 170, 331 moulding, 17 mouse, 119 movement, 211, 213 MOVPE, 241, 244 MRI, 124, 125, 141, 145 MRS, 270 MTS, 218 multidisciplinary, 108 multiple factors, 121 multiples, 95, 258, 501 multiplier, 375 myeloid, 127

N NA, 73, 77, 93, 569 NaCl, 24, 517, 520, 573, 576, 582, 591, 592, 594, 595 NADH, 38 nanobelts, 238, 315 nanobiology, 315 nanoclusters, 5, 109, 135, 137, 139, 526 nanocomposites, 144, 315, 321, 357, 363, 510, 526 nanocrystal, 117, 140, 412, 416, 417, 427, 429, 467, 480, 499, 501, 502 nanocrystalline, xiii, 5, 8, 9, 117, 344, 345, 347, 367, 369, 435, 454, 455, 490, 491, 518, 519, 521, 522, 523, 618, 620 nanocrystals, xiii, xiv, 6, 8, 140, 142, 145, 147, 246, 300, 409, 411, 422, 427, 429, 430, 431, 433, 465, 479, 480, 481, 489, 490, 492, 494, 497, 500, 502, 503, 504, 505, 506, 510, 526 nanodevices, xvi, 238, 239, 240, 248, 297, 299, 315, 372, 559, 561, 606 nanodots, viii, 4, 25, 37 nanoelectronics, 23, 371, 612 nanofabrication, vii, viii, xi, 3, 4, 9, 10, 13, 15, 21, 24, 38, 40, 42, 43, 293, 295, 298, 299, 315 nanofibers, viii, ix, xii, 31, 51, 53, 64, 65, 66, 67, 68, 371, 377, 403

634 nanofibrous membranes, 65 nanoindentation, x, 151, 171, 173, 174, 175, 181, 186, 190, 191, 192, 193, 194, 197, 198, 199, 201, 202, 203, 204, 205, 212, 213, 216, 217, 219, 222, 228 nanolithography, 9, 11, 13, 14, 15, 40, 297, 298, 299, 303 nanomaterials, viii, xi, xii, xiv, xvi, 4, 11, 42, 88, 149, 246, 293, 294, 295, 297, 298, 299, 301, 303, 309, 310, 328, 372, 374, 388, 410, 414, 479, 480, 506, 615, 619 nanomedicine, 144, 147 nanometer, vii, ix, xv, 40, 71, 72, 73, 80, 81, 82, 108, 116, 201, 265, 294, 295, 333, 338, 349, 350, 442, 518, 523 nanometer scale, 72, 80, 81, 82, 108, 295, 338, 349, 518, 523 nanometers, 11, 52, 80, 117, 171, 172, 190, 250, 303 nanoparticle synthesis, 109, 128 nanoparticulate, 147, 148 nanoribbons, 411 nanorods, viii, xiii, 4, 26, 27, 238, 245, 246, 248, 249, 250, 254, 256, 258, 261, 263, 315, 409, 410, 490, 491, 492 nanoscale structures, xiii, 117, 409 nanoscience, 238, 294 nanosheets, 42, 252, 484, 485 nanostructured materials, 23, 24, 247, 332 nanostructures, 8, 11, 50, 270, 273, 315, 384, 411, 412, 413, 415, 417, 419, 421, 423, 425, 427, 429, 431, 433 nanotechnologies, 39 nanotechnology, vii, xvi, 3, 43, 52, 108, 141, 142, 144, 146, 147, 238, 294, 303, 392, 510, 560, 596 nanotube, 26, 27, 28, 29, 30, 31, 32, 34, 35, 41, 104, 144, 315, 323, 372, 373, 381, 389, 390, 391, 397, 403, 411 National Academy of Sciences, 140, 143, 144 National University of Singapore, 331 NATO, 405, 455, 599 natural, x, 107, 126, 397, 522, 559, 560, 584, 615 Nd, xiv, 412, 479, 480, 489, 501, 562, 565, 584 necrosis, 126 NEMS, vii, 3, 21 neovascular, 119 network, xiii, 110, 224, 228, 417, 436, 585 neural tissue, 40 neuroblastoma, 146 neuroendocrine, 116 neutrophils, 119 New York, 43, 141, 147, 272, 304, 305, 307, 405, 431, 476, 556, 596, 599, 600, 602 nickel (Ni), xvi, 8, 9, 19, 20, 21, 25, 26, 37, 127, 355, 525, 527, 528, 529, 530, 531, 532, 533, 534, 536, 537, 538, 539, 540, 541, 542, 543, 544, 545, 551, 552, 553, 571 NiO, 9, 527, 528, 530, 531, 532, 533, 534, 535, 536, 539, 540 NIR, 109, 122, 123, 124, 502, 503, 504

Index nitrate, 247, 252, 615, 616 nitride, xiii, xv, 14, 84, 99, 100, 140, 239, 243, 244, 245, 246, 263, 409, 411, 412, 413, 414, 416, 417, 419, 420, 421, 422, 425, 426, 427, 429, 431, 509, 510 nitrides, 510 nitrogen, 152, 171, 262, 280, 283, 350, 376, 416, 417, 511, 512, 522, 585 nitrogen gas, 511, 512, 522 NMR, 282, 287 noise, 514, 515, 574 non-destructive, 77 nonionic, 115, 129, 482 nonlinear, xv, 187, 494, 525, 526, 527, 560, 561 non-linear, xvi, 559, 578, 581, 583 non-small cell lung cancer, 133 nontoxic, 129 nontoxicity, 618 non-uniform, 397 non-uniformity, 397 normal, ix, xv, 26, 75, 89, 90, 91, 93, 94, 95, 100, 107, 110, 111, 112, 116, 126, 127, 250, 269, 345, 348, 361, 374, 381, 383, 388, 399, 438, 468, 474, 509, 514, 516, 517, 519 novel materials, 411 n-type, 7, 17, 24, 28, 240 nuclear, 124, 145, 238 nuclear magnetic resonance, 145 nucleating agent, 565 nuclei, 21, 22, 24, 120, 246, 250, 251, 253, 269, 353, 422, 443 nucleic acid, 146, 560 nucleus, 113, 126, 402, 403, 429 numerical aperture, 73, 77 nutrients, 110, 615 nutrition, 110, 142

O observations, 53, 161, 216, 340, 343, 346, 378, 414, 429, 544, 608, 609 occlusion, 109, 128 odors, 560 oil, 35, 73, 248 oncology, 142 Oncology, 146, 149 one dimension, ix, 107, 108 optical activity, 602 optical anisotropy, 564, 575 optical fiber, 77, 78, 84, 283, 439 optical imaging, 104 optical microscopy, 102, 103, 104, 226, 584 optical phonon confinement, 257, 258 optical properties, xiv, 31, 122, 123, 124, 141, 258, 314, 321, 332, 473, 479, 480, 526, 560, 561, 562, 582, 584, 596 optical tweezers, 104 optics, 77, 93, 94, 95, 102, 103, 104, 309, 314, 506, 595

Index optimization, viii, 51, 100, 126, 159, 514 optoelectronic, 239, 258, 300, 310, 460 opto-electronic, 339 optoelectronics, xi, 239, 294, 301 ordinary differential equations, 566 organic, viii, xi, 4, 11, 13, 14, 15, 25, 29, 30, 42, 109, 117, 119, 128, 129, 133, 141, 275, 276, 277, 278, 279, 280, 281, 291, 295, 297, 302, 309, 310, 311, 312, 313, 315, 316, 317, 328, 355, 493, 584, 617 organic polymers, 309 organic solvent, 109, 128, 129, 133, 295, 297, 310, 311, 313, 315, 316, 317, 328 organic solvents, 109, 129, 133, 295, 297, 311, 315, 328 organic thin films, 277 organization, xvi, 417, 431, 559, 560, 561 orthorhombic, x, 151, 152, 159, 161, 171, 199, 201, 206, 207, 208, 218, 224, 226, 228, 229, 230, 341, 343 oscillations, 122, 312, 401, 546 oscillator, 375 osmolality, 128 osteoblasts, 149 osteosarcoma, 110 ovarian cancer, 143 oxidation, 8, 11, 31, 38, 41, 103, 121, 247, 282, 283, 316, 402 oxidative, 33, 121, 149, 282 oxidative stress, 121, 149 oxide nanoparticles, xvi, 8, 310, 525, 526, 528 oxides, viii, xvi, 29, 125, 159, 190, 265, 332, 333, 337, 340, 341, 343, 354, 363, 364, 365, 366, 368, 484, 485, 525, 617 oxygenation, x, 152, 160, 164, 167, 170, 171, 199, 224, 225, 226, 227, 228, 229, 230

P p53, 143 PAA, viii, ix, 38, 51, 53, 54, 55, 56, 58, 59, 60, 64, 65, 66, 68 paclitaxel, 142 pain, 619 paints, 311 palladium, 9, 36, 79 PAN, 318, 325 pancreatic, 133 PANI, 23, 24, 32, 33, 34, 37, 38 parabolic, 180, 381, 382, 393 paramagnetic, 261, 288, 289, 544 parameter, xvi, 56, 59, 73, 96, 153, 176, 177, 179, 183, 199, 302, 344, 346, 403, 437, 462, 465, 467, 473, 504, 525, 538, 552, 565, 575, 578, 580, 582, 588, 589 particle mass, 569 passivation, 17 passive, 7, 113, 116, 119, 125, 127 pathways, 144, 277, 278, 280, 281 patients, 128, 140, 143, 145, 620

635

pattern recognition, 300 patterning, 9, 11, 40, 297, 429, 431 Pb, xii, 120, 153, 331, 332, 333, 337, 339, 342, 343, 344, 345, 353, 356, 362, 363, 364, 365, 366, 367, 369 PCA, 335 PDMS, 606, 607, 608, 609, 611, 612 peptide, 147, 148, 149, 606 peptides, 120 percolation, 568, 569 performance, ix, xi, 20, 24, 71, 80, 171, 239, 270, 277, 291, 293, 294, 300, 342, 345, 347, 355, 361, 362, 460, 480, 488, 493, 505, 516, 618 pericytes, 110 periodic, 25, 27, 411, 416, 426, 526 periodic table, 526 periodicity, 9 permeability, 111, 140, 150 permeation, 111 permit, 9, 199 permittivity, 356 perovskite, 340, 341, 342, 343, 344, 345, 346, 347, 348, 350, 354, 365, 366, 367, 369 perovskites, 365 peroxide, 323 perturbation, 11, 102, 278 PET, 28 pH, 6, 7, 27, 31, 38, 114, 115, 128, 135, 143, 146, 147, 245, 247, 249, 252, 314, 315, 560, 561, 563, 566, 570, 575, 576, 577, 579, 580, 585, 591, 592, 594, 595 pH values, 570 phagocyte, 147 pharmaceutical, 148 phase diagram, 154, 155, 157, 341, 362 phase transformation, 167, 335, 342 phase transitions, 569 phenol, 316 phenomenology, 596 phonon, xi, 74, 75, 79, 80, 81, 82, 83, 84, 86, 87, 90, 94, 95, 96, 98, 100, 101, 102, 237, 255, 256, 257, 258, 269 phonons, ix, 71, 74, 75, 90, 102, 255, 258, 498 phospholipids, 112, 147 phosphor, 6, 483, 484, 498, 506 phosphors, 326 photobleaching, 109, 119 photocatalysis, 560 photocatalysts, 618 photochemical, xi, 30, 275, 278, 281, 283, 286, 287, 291 photodetectors, 294, 300, 301, 302, 303 photoelectron spectroscopy, 412, 527 Photoelectronic, 307 photoexcitation, 376 photographs, 497 photoionization, 560 photolithography, 297, 299

636 photoluminescence, xi, xiv, xv, 237, 241, 269, 459, 460, 464, 473, 525, 526, 527, 550 Photoluminescence, xiv, 259, 461, 463, 464, 465, 467, 469, 471, 473, 475, 477, 479, 551 photoluminescence spectra, xi, 269, 464 photolysis, 29, 121 photon, 117, 258, 460, 465, 466, 467, 475, 493 photonic, 35, 36, 311, 460, 468, 469, 475 photons, xiv, 117, 459, 465, 467, 469 photosynthesis, xvi, 559 photovoltaic, 6, 7 photovoltaics, 6 physical properties, x, xi, 237, 239, 240, 241, 258, 269, 272, 275, 287, 411, 566 physicists, 410 physicochemical, 33, 121, 143, 149 physicochemical properties, 121 physico-chemical properties, 561 physics, 145, 233, 378, 404, 506, 599 physiological, 128, 133 physiology, 118 piezoelectric, xii, 27, 239, 331, 332, 335, 339, 340, 347, 349, 350, 351, 352, 354, 357, 359, 360, 362, 363, 366, 368, 369, 370 piezoelectric properties, 239, 347, 349, 350, 352, 362, 363 PL spectrum, 259, 260, 550 planar, 239, 269, 278, 352, 386, 429, 578 planetary, 333, 334, 335, 337, 340, 343, 351, 352, 353, 355 plasma, xv, 28, 126, 133, 336, 437, 438, 439, 443, 447, 454, 455, 456, 523, 525, 527, 539, 546 plasmons, 102, 123, 582 plastic, x, 14, 53, 151, 152, 161, 165, 171, 173, 174, 176, 177, 178, 181, 183, 184, 187, 188, 189, 190, 191, 192, 193, 197, 216, 217, 218, 276, 311, 373, 397, 402, 518 plastic deformation, x, 151, 161, 165, 171, 173, 177, 178, 181, 187, 188, 189, 190, 191, 216, 217, 218, 373, 397, 402, 518 plasticity, 188, 190, 210, 277 platforms, 13, 299 platinum, 14, 15, 17, 36, 369 plausibility, 373 play, 24, 158, 224, 238, 245, 278, 279, 505, 526, 578 PLC, 512 PLD, 241, 244, 461, 462, 467, 471 PLGA, 129, 130, 131, 132, 140, 145, 147, 148, 594, 595 PMMA, 311, 313, 317, 325 PN, 582 point defects, 354, 460, 540, 541 point-to-point, 172 Poisson, 179, 224 Poisson ratio, 224 polarity, 240 polarizability, 90, 91, 99, 563, 564, 574, 582, 583, 585

Index polarization, ix, 10, 41, 42, 72, 76, 77, 88, 89, 92, 93, 95, 96, 97, 98, 100, 101, 104, 239, 255, 287, 340, 342, 357, 366, 373, 394, 593, 594 polarized light, 77, 89, 164, 226, 590, 593 polarized light microscopy, 164 pollutants, 618 poly(lactic-co-glycolic acid), 129 poly(methyl methacrylate), 111, 326 poly(vinylpyrrolidone), 144 polyamine, 580, 584, 585, 588 polyaniline, 13, 22, 23, 24, 31, 32, 33, 37 polyaniline (PANI), 22, 31, 37 polycarbonate, 40 polycondensation, 40, 319 polycrystalline, 4, 5, 7, 15, 161, 194, 195, 247, 248, 289, 335, 362, 490, 496, 497, 548, 549 polydispersity, 310, 313 polyester, 616 polyethylene, 28 polymer chains, 310, 315, 317, 320, 324 polymer film, 9, 33 polymer materials, 171, 310 polymer matrix, 33, 122, 132 polymer solutions, 53, 57 polymer structure, 22, 324 polymer-based, x, 107 polymerization, 23, 31, 32, 38, 41, 129, 149, 309, 310, 311, 315, 316, 323, 324, 325, 327, 328 polymerization time, 31, 32, 38, 41 polymerizations, xii, 309, 310, 326 polymers, viii, 4, 10, 13, 21, 22, 23, 31, 33, 34, 35, 40, 42, 52, 68, 119, 123, 128, 129, 130, 132, 142, 172, 176, 309, 310, 311, 312, 313, 319, 328, 569 polypeptide, 594, 595 polypeptides, 590 polysaccharides, 129 polystyrene, viii, 4, 24, 28, 35, 36, 37, 318, 327, 460 polystyrene latex, 36 population, 110, 161, 171, 207, 263 pore, 6, 16, 17, 25, 27, 29, 30, 52, 63, 111, 112, 311, 482 pores, 16, 17, 24, 25, 27, 31, 40, 58, 63, 228, 311, 481 porosity, 199, 201, 205, 207, 208, 210, 213, 216, 217, 226, 227, 228, 345, 575 porous, 17, 22, 25, 31, 33, 36, 55, 66, 353, 419, 490, 569 porous materials, 569 porphyrins, 560, 561, 563, 564, 565, 566, 570, 572, 573, 574, 575, 578, 579, 585, 586, 590, 591, 594, 595 potassium, 31 potential energy, 400, 401 p-polarized, 89, 97, 98 precipitation, 6, 7, 27, 129, 343, 549 prediction, 263, 280, 291, 336, 411, 573 pre-existing, 110 pressure, 87, 88, 93, 100, 104, 129, 153, 154, 155, 158, 174, 182, 221, 223, 224, 276, 277, 353, 375,

637

Index 392, 396, 411, 437, 438, 440, 443, 451, 454, 461, 511, 512, 514, 518, 522, 523 prevention, 142, 144, 147 printing, 358, 431 pristine, 282, 283, 287, 320, 414, 418, 431 private, 294 probability, 375, 397, 437, 566, 567, 568, 569, 570, 571, 579, 580, 588 probe, vii, xi, 3, 9, 10, 11, 12, 42, 77, 79, 104, 177, 181, 237, 376, 377, 498, 595 process control, 335 process gas, 455 production, 17, 23, 27, 36, 159, 247, 276, 352, 354, 358, 362, 372, 411, 615 productivity, 269, 333, 337 progesterone, 313 program, 7, 506 proliferation, 108, 137, 148 promote, 108, 301, 356, 411 promoter, 607, 612 propagation, 14, 89, 160, 161, 170, 184, 187, 201, 216, 217, 228, 255, 439, 451, 469 property, xv, 179, 197, 199, 268, 280, 286, 287, 290, 436, 438, 461, 509, 510, 512, 515, 527, 542, 543, 544, 561, 616, 617, 619 proportionality, 568 propylene, 111 prostate, 108, 119, 133, 134, 142, 143 prostate cancer, 119, 133, 134, 142, 143 prostatectomy, 143 proteases, 110 protection, 119, 616, 617 protective clothing, 619 protective coating, 414, 510 protein, 9, 118, 143, 149, 615 proteins, 115, 119, 121, 125, 126, 129, 137, 146, 149 protocol, 296, 578 protons, 125 prototype, 42, 286 prototyping, 296, 299 pseudo, 155 PSS, 28, 29, 38, 320 PST, 343, 345, 346, 349, 405 PT, 335, 336, 337, 338, 339, 340, 343, 344, 347, 348, 350 p-type, 7, 12, 14, 17, 29, 240, 241, 412 public, 294, 304 publishers, 413, 425, 432 pulse, 18, 19, 20, 21, 42, 125, 207, 375, 376, 377, 378, 388, 395, 404, 412, 418, 425, 426, 461, 469, 493, 511, 512 pulsed laser, xiii, xiv, 241, 242, 243, 410, 431, 459, 461, 526 pulsed laser deposition, xiv, 241, 242, 243, 459, 526 pulses, 10, 11, 16, 18, 19, 20, 21, 300, 303, 376 pumps, 375 purification, 312, 410, 412, 618 PVA, viii, 51, 53, 66, 67, 68, 130, 131, 159 PVP, 315

pyramidal, 189 pyrolysis, xv, 325, 332, 390, 392, 490, 605, 606, 607, 608, 609, 611, 612 pyrolytic graphite, 8, 9, 13, 41 pyrrole, 13, 21, 32, 35, 41

Q QDs, 117, 118, 119, 120, 121, 141 quadrupole, 146 quality control, 171 quantitative estimation, 283 quantum, viii, x, 4, 7, 8, 17, 107, 108, 109, 117, 118, 140, 141, 142, 143, 144, 145, 148, 149, 265, 267, 321, 372, 396, 465, 527, 532, 542, 563, 592 quantum confinement, 117, 532 quantum dots, viii, x, 4, 7, 8, 107, 108, 109, 140, 141, 142, 143, 144, 145, 148, 149, 321 quantum yields, 109, 118 quartz, 53, 283, 412, 439, 551, 606

R R&D, 506 radial distribution, 417 radiation, ix, xiv, 18, 107, 111, 121, 124, 127, 397, 412, 436, 451, 452, 453, 492, 512, 517, 564, 574, 615, 616, 617, 618 radiation therapy, ix, 107, 111 radical, xi, xii, 143, 275, 278, 279, 282, 309, 310, 315, 316, 324, 441, 451 radical polymerization, xii, 309, 310, 315, 316, 328 radio, 125, 252 radiofrequency, 127 radiotherapy, 123 radium, 223 radius, xiv, 117, 122, 123, 176, 181, 182, 184, 190, 266, 269, 374, 375, 376, 379, 380, 381, 383, 384, 385, 386, 388, 389, 390, 393, 479, 480, 494, 565, 566, 567, 568, 569, 570, 571, 572, 575, 580, 583 Raman scattering, ix, 71, 72, 73, 74, 75, 77, 82, 85, 86, 90, 102, 103, 241, 255, 257, 258, 473, 583, 584, 586 Raman spectra, ix, xiii, 71, 73, 75, 80, 82, 86, 87, 90, 94, 95, 99, 100, 104, 284, 287, 435, 441, 445 Raman spectroscopy, ix, 71, 72, 75, 101, 102, 103, 104, 412, 439 Raman-scattering, 102, 103 random, 16, 17, 349, 379, 390, 416, 548, 549 Rayleigh, 52, 561, 583, 584, 586, 587, 588 reactant, 5, 7, 241, 242, 437, 456, 578, 579, 580, 611 reaction rate, 18, 567 reaction temperature, 340 reaction time, 412 reactivity, 11, 324, 333, 340, 345, 353, 362, 374, 522 reagent, 6, 578, 579, 580, 585 reagents, 145, 238, 578, 579, 587 real time, 380, 588 reality, 404

638 receptors, 112, 113, 115, 116, 119, 147 recognition, 119, 152, 300, 560 recombination, 260, 263, 264, 541, 546, 551, 579 reconstruction, 397 recovery, 125, 126, 176, 187, 191, 192, 210, 212, 302, 354 recrystallization, 368, 549 recrystallized, 534, 549, 550 red shift, 260 red wine, 618 redox, 41, 133, 282, 291 redshift, 257, 259, 260, 261, 263, 269 reduction, ix, 8, 9, 15, 25, 27, 33, 36, 41, 52, 71, 72, 75, 86, 99, 101, 187, 207, 215, 229, 247, 334, 343, 344, 345, 350, 354, 355, 361, 520, 523, 606 reference frame, 90, 91 reflectance spectra, 286 reflection, 76, 77, 79, 85, 87, 91, 92, 102, 103, 286, 346, 470, 474, 513 refractive index, 73, 84, 92, 93, 468, 470, 473, 474, 547, 569 regression, 144, 563 regression analysis, 563 regular, xiv, 6, 37, 170, 380, 423, 429, 448, 479, 489, 501, 502, 615 regulation, 104 relationship, xv, xvi, 11, 146, 188, 191, 192, 215, 250, 265, 313, 345, 416, 447, 465, 509, 510, 513, 525, 537, 549, 552 relationships, 213 relative size, 121 relaxation, 125, 288, 290, 346, 403, 489, 500, 572, 575, 577 relaxation rate, 572, 575, 577 relevance, 294, 590 Reliability, 305remission, 127 renal, 116, 119, 141 renal cell carcinoma, 141 repair, 297 replication, 27, 462 research, vii, x, xii, 24, 88, 107, 108, 139, 144, 146, 239, 241, 276, 294, 299, 303, 312, 321, 324, 333, 337, 371, 374, 376, 397, 401, 404, 411, 430, 448, 475, 506, 510, 560, 606, 616, 617 researchers, 89, 118, 119, 121, 127, 140, 152, 160, 276, 295, 302, 303, 310, 351, 356, 410, 460 reservoir, 10, 153 reservoirs, 41, 42 residuals, 560 residues, 493 resin, 392 resistance, 19, 119, 169, 171, 184, 190, 191, 192, 193, 212, 300, 302, 303, 373, 403, 510, 615 resistive, 24, 373 resistivity, 17, 249, 252, 277, 278, 279, 282, 283, 284, 285, 286, 287, 288, 291, 302 respiratory, 133 responsiveness, 319 retention, 111

Index returns, 125 Reynolds, 273, 477 RF, 125, 461 rhombohedral, 340, 341, 343 RIE, 461, 462, 471 rings, 373, 380, 391, 410, 496, 497, 520 rods, 246, 394, 417, 427, 428, 429, 537, 578, 591 room temperature, 32, 42, 53, 77, 126, 159, 161, 169, 175, 197, 201, 218, 239, 265, 266, 283, 314, 333, 347, 348, 412, 460, 461, 463, 464, 470, 471, 479, 523, 526, 527, 538, 543, 545, 546, 550, 551, 555, 606 room-temperature, 239, 240, 248, 260 rotational matrix, 91 roughness, 20, 175, 254, 286, 448, 467 Rouleau, 271 Royal Society, 413, 418 RP, 619 rubber, 171, 311, 606 rubber products, 311 Russia, 456 rutile, 367, 480, 488, 489

S SAD, 535, 536, 537, 549, 552 saline, 124 salt, 8, 129, 246, 332, 343, 344, 350, 575, 576, 591 salts, xi, 129, 245, 275, 277, 278, 280, 282, 332 samarium, 482 sampling, 512 sapphire, 101, 173, 242, 243, 259, 493, 551 sarcomas, 132 satellite, 528, 530, 537 saturation, 156, 242, 246, 265, 267, 269, 287, 378, 422, 544, 554 scaffold, 461, 471 scalable, 23 scaling, 567, 570, 581, 583, 584, 585, 587 scaling law, 570, 583, 584, 585, 587 scandium, 343, 345, 365 Scanning electron, 20, 21, 135, 286, 413, 616, 617 scanning electron microscopy, xiv, xv, 32, 215, 242, 412, 459, 605, 606 scanning electronic microscope, 439 scanning tunneling microscopy, 9, 374 scatter, 109, 123, 208, 592 scattered light, 97, 98, 562, 592 schema, 225 Schmid, 104, 234 Schottky, 291 scientific community, 295 scientists, vii, ix, xii, 3, 107, 120, 309, 310, 315, 316, 317, 325, 410 SCs, 152 security, xii, 331 sediments, 412

Index seed, xiii, 158, 159, 161, 169, 208, 224, 238, 252, 254, 409, 412, 419, 420, 421, 422, 425, 427, 428, 429, 430, 431 seeding, 158, 159, 345 seeds, 28, 246, 251, 362, 423, 426, 427, 430, 578, 586, 588 selected area electron diffraction, 483, 521, 527, 535, 537, 549 selectivity, 11, 300, 303, 560 selenium, x, 107, 109, 133, 134, 135, 136, 137, 138, 139, 142, 144, 145, 146, 147, 148, 149 self-assembling, 299, 410 self-assembly, vii, viii, xi, xiii, xiv, xvi, 3, 4, 24, 40, 237, 238, 239, 241, 251, 254, 269, 276, 299, 319, 409, 410, 411, 424, 429, 430, 431, 460, 461, 471, 473, 485, 490, 559, 560, 561, 578, 590, 596 Self-cleaning, 620 self-organization, 411, 425, 430, 431 self-similarity, 567 SEM micrographs, 23 semiconductor, vii, x, xiv, 3, 5, 6, 7, 8, 11, 14, 22, 27, 29, 72, 121, 141, 142, 143, 149, 161, 237, 238, 239, 241, 264, 265, 266, 269, 299, 301, 414, 479, 480, 506, 526 semiconductors, xii, 4, 5, 7, 8, 10, 27, 37, 42, 72, 117, 247, 263, 276, 277, 278, 299, 301, 302, 309, 316, 354, 411 sensing, xi, 27, 34, 42, 65, 66, 123, 141, 149, 261, 294, 299, 300, 302, 314, 590 sensitivity, viii, 31, 40, 51, 66, 68, 77, 82, 84, 87, 88, 95, 100, 103, 255, 258, 278, 303, 332, 394, 595 sensitization, xiv, 479, 480, 501 sensors, viii, xii, 9, 23, 24, 27, 34, 51, 65, 66, 240, 276, 294, 300, 302, 303, 314, 328, 331, 332, 339, 342, 347 separation, viii, 19, 20, 51, 54, 57, 67, 68, 118, 301, 309, 312, 313, 380 series, 10, 177, 288, 291, 334, 380, 392, 465, 519, 599 serum, 130, 139, 149 serum albumin, 130 Shanghai, 51, 317, 605, 612 shaping, 17, 392, 393, 394 sharing, xiii, 409 shear, x, 151, 172, 174, 177, 191, 447, 450 Shell, 460, 471 short-range, 255, 258, 398, 568 shoulder, 75, 80, 86, 539 Si3N4, 84, 85, 87, 88, 98, 99, 100, 606 side effects, ix, 107, 112, 126, 140 Siemens, 16 sign, 193, 333, 561, 591, 593, 595 signals, ix, 71, 72, 73, 82, 85, 89, 94, 95, 98, 99, 103, 120, 125, 261, 375, 464, 493, 514, 592 signal-to-noise ratio, ix, 72, 99, 100 signs, 127, 593 silane, 327 silanol groups, 311 silica, 23, 40, 121, 122, 184, 311, 312, 326, 532, 551

639

silica glass, 532 silicon dioxide, 84, 310 siloxane, xv, 605, 606 silver, 102, 103, 104, 244, 278, 280, 283, 321, 615, 616, 617, 619, 620 simulated emission, 248 simulation, xiv, 276, 374, 376, 379, 399, 401, 404, 459, 460, 469, 470, 475, 585 simulations, xii, 10, 187, 371, 373, 380, 397, 398, 400, 401, 402, 404, 568, 582, 584 Singapore, 331 single crystals, xi, xvi, 153, 158, 159, 165, 211, 214, 275, 276, 277, 278, 281, 282, 286, 457, 525, 530, 531, 532, 539, 540, 541, 546, 547, 548 single walled carbon nanotubes, 34 single-crystalline, 16, 17, 28, 238, 246, 410 sintering, xii, 158, 331, 332, 335, 336, 338, 344, 346, 347, 349, 351, 353, 355, 358, 361, 362, 364, 367, 368 SiO2, viii, xv, 11, 51, 53, 67, 68, 84, 85, 312, 325, 326, 460, 474, 475, 546, 547, 551, 605, 606, 607, 608, 609, 610, 611, 612 sites, xiv, 22, 39, 112, 113, 116, 119, 127, 140, 156, 157, 161, 207, 241, 242, 399, 403, 423, 479, 493, 497, 498, 499, 500, 506, 611, 617 skeleton, 17, 228, 517, 520 skin, 148, 617 skin cancer, 148 Sm, xiv, 479, 480, 501 smooth muscle, 148 smoothing, 17, 377 SNR, ix, 71, 72, 99, 100 sodium, 135, 138, 318, 320 soft matter, viii, 4, 10, 31, 33 software, 197, 218, 300 SOI, 460, 465 solar, 27, 29, 31, 52, 240, 301, 302, 618 solar cell, 27, 29, 31, 52, 240, 301, 302 solar cells, 27, 29, 31, 52, 240, 301 solar energy, 301 sol-gel, 31, 40, 245, 332, 336, 337, 339, 342, 343, 350, 352, 357, 358, 359, 361, 369, 489, 490, 492, 498, 506 solid phase, 246 solid solutions, 339 solid state, xi, 275, 277, 278, 282, 287, 288, 344, 345, 367, 369, 480 solid tumors, 142, 145 solidification, 60, 156, 354 solid-state, xii, 165, 331, 332, 333, 335, 338, 342, 343, 344, 345, 346, 347, 349, 350, 352, 353, 354, 362, 367, 369 sols, 369 solubility, 129, 277, 310, 316, 317, 328, 332, 480 solution phase method, 241 solutions, 5, 24, 25, 27, 33, 38, 52, 53, 57, 66, 67, 128, 135, 136, 145, 239, 245, 246, 303, 316, 339, 347, 352, 358, 361, 563, 566, 567, 569, 570, 572, 575, 576, 577, 579, 580, 581, 588, 590, 596

640 solvent, viii, 4, 7, 32, 52, 54, 55, 56, 57, 58, 59, 61, 65, 129, 130, 141, 147, 310, 422, 490, 560, 574, 589 solvents, viii, 30, 31, 51, 54, 55, 56, 68, 129 space-time, 579 Spain, 151, 170, 293 spatial, ix, 18, 71, 72, 77, 172, 186, 283, 294, 376, 403, 569, 578, 579, 588, 595 spatial confinement, 18 species, xvi, 147, 239, 241, 242, 245, 248, 277, 278, 282, 288, 291, 300, 302, 303, 324, 345, 372, 375, 376, 399, 411, 437, 438, 439, 454, 505, 527, 559, 560, 561, 565, 566, 567, 568, 579, 590, 611 specific heat, 143 specific surface, 68 specificity, ix, 107, 120, 143 spectral dimension, 583, 586 spectroscopic methods, 288 spectroscopy, ix, xiv, 9, 27, 71, 72, 93, 95, 102, 103, 104, 135, 139, 464, 473, 479, 481, 493, 497, 500, 527, 552, 588, 590, 606 speed, viii, ix, xv, 4, 11, 13, 14, 15, 20, 51, 55, 56, 60, 64, 66, 67, 71, 110, 333, 334, 335, 341, 351, 356, 402, 412, 414, 425, 431, 509, 512 spermine, 580, 581, 584, 585, 586, 587 spheres, viii, 4, 24, 36, 37, 140, 173, 280, 311, 383, 414, 418, 461, 462, 473, 491 spin, viii, 4, 125, 239, 296 spleen, 111, 133 SPR, xv, 525, 526, 527, 540, 541, 547 sputtering, 252, 266, 526 SQUID, 527, 542 S-shaped, 260, 263, 269 stability, xi, 10, 39, 102, 109, 116, 121, 123, 131, 134, 175, 251, 293, 300, 310, 312, 313, 317, 344, 348, 364, 366, 369, 373, 376, 377, 378, 397, 404, 411, 437, 500, 560, 618 stabilization, 66, 210, 368, 369 stabilize, 124, 278, 437 stabilizers, 130 stages, ix, 10, 20, 107, 108, 111, 140, 158, 172, 176, 210, 288, 290, 352, 385, 398, 422, 423, 424, 430, 496 stainless steel, 20, 108, 134, 135, 137, 138, 334, 343, 352, 412 stainless steels, 20 standard deviation, 63 Staphylococcus, 616, 617 Staphylococcus aureus, 616, 617 stars, 372 STD, 520 steel, xv, 108, 134, 135, 137, 138, 173, 333, 334, 343, 351, 352, 412, 456, 509, 510, 511, 512, 516, 520, 523 STEM, 323, 527, 533, 536, 538 stem cell therapy, 145 steric, 578 sterilization, 618

Index stiffness, 172, 175, 176, 177, 179, 181, 183, 198, 218, 219 stimulus, 115 STM, 9, 10, 11, 14 stock, 566, 578, 579, 582 stoichiometry, 352 Stokes shift, 583, 586 storage, xiii, xv, 109, 116, 152, 158, 172, 173, 190, 312, 342, 351, 409, 410, 525, 526 strain, ix, 71, 72, 73, 74, 75, 76, 77, 80, 81, 82, 84, 87, 88, 90, 92, 93, 100, 102, 103, 104, 165, 166, 175, 181, 182, 188, 190, 191, 218, 223, 240, 342, 400, 403, 494 strains, ix, 71, 72, 80, 81, 173, 187, 335, 342, 403 strategies, viii, 4, 16, 24, 25, 52, 109, 114, 245, 277, 297, 300, 431 stress level, 177 stress-strain curves, x, 151, 174, 181, 218 stretching, 372, 401, 402, 403, 416, 417 strong interaction, 38, 314, 562 strontium, 352, 361, 368, 370 structural changes, 226, 574 structural dimension, 422 structural relaxation, 403 structure formation, 38 students, 291 styrene, 38, 310, 324, 326 substances, 276 substitutes, 500 substitution, 498 substrates, ix, xiv, 5, 7, 8, 9, 13, 14, 15, 18, 23, 25, 26, 27, 28, 39, 71, 72, 77, 84, 105, 134, 135, 136, 137, 138, 139, 168, 241, 243, 246, 247, 249, 250, 252, 266, 269, 276, 310, 320, 358, 436, 447, 459, 488, 526, 551 sulfur, 611, 612 sulfuric acid, 12 sulphate, 140 sulphur, 38 superconducting, 152, 153, 154, 158, 167, 170, 224, 527, 542 superconducting materials, 153, 170 superconductivity, 27, 153, 236, 278 superconductor, x, 151, 152, 163, 170 superconductors, 152, 153, 158, 159, 161, 170, 224, 291 supercritical, 291 superhard, xv, 411, 509, 510, 513, 514, 515, 516, 517, 518, 519, 520, 521, 522, 523 superimpose, 258 superlattice, 346 superlattices, 5 superposition, 386, 581 supply, 21, 53, 110, 156, 374, 394, 511 suppression, ix, 72, 73, 89, 97, 99, 101, 396 supramolecular, xvi, 559, 560, 561, 578, 590, 591, 593, 594, 595 surface area, 13, 23, 30, 31, 38, 64, 311, 549, 617 surface chemistry, 145

Index surface diffusion, 242, 437 surface diffusivity, 438 surface energy, 144, 240, 245, 251, 269, 315, 450 surface layer, 9, 90, 261, 263, 376, 379 surface modification, 13, 117 surface properties, 109, 114, 132, 250, 312 surface region, 401, 414, 526 surface roughness, 20, 175 surface structure, 4, 470 surface tension, 52, 56, 437 surface-initiated atom transfer radical polymerization, xii, 309, 310, 328 surfactant, 121, 128, 129, 130, 482 surfactants, 40, 410 surgery, ix, 107 Surgery, 107 surgical, 108 surging, 52 susceptibility, xv, 125, 287, 288, 289, 525, 526 suspensions, 134, 357, 431 switching, xv, 318, 351, 366, 525, 526 SWNTs, 34, 315, 323 symbols, 577 symmetry, 25, 184, 238, 255, 414, 470, 498, 499, 500, 561, 574, 590 synchronous, 141 synchrotron, xiv, 18, 436, 451 synchrotron radiation, xiv, 436, 451 Synchrotron radiation, 452 synthetic polymers, 310

T tactics, 247 targets, 111, 113, 118, 148, 149, 511, 512 Taxol, 142 technology, vii, xiii, xvi, 3, 17, 18, 72, 108, 276, 277, 291, 295, 297, 300, 303, 310, 319, 360, 363, 372, 374, 392, 404, 409, 410, 453, 456, 526, 615 tellurium, 5 temperature annealing, 505, 526 temperature dependence, 77, 283, 342 temperature gradient, 158, 159, 162, 438 tensile, 90, 169, 184, 400, 403, 404, 442 tensile strength, 403, 404 tensile stress, 90, 184, 403, 442 tension, 228, 398, 402 terraces, 448 Tesla, 125, 127 testicular cancer, 133 textile, 615, 616, 619, 620 textiles, xvi, 52, 615, 617, 618, 619 thallium, 11 theory, 179, 190, 191, 223, 255, 256, 265, 268, 302, 394, 397, 404, 547, 563, 564, 565, 566, 584, 588 therapeutic agents, 115, 120, 127, 140 therapeutics, 109, 123, 128, 142, 148, 149 therapy, x, 107, 120, 123, 124, 141, 142, 143, 144, 145, 146, 149, 560

641

thermal activation, 352 thermal analysis, 282, 339, 492 thermal decomposition, 489 thermal energy, 123, 239, 389, 506 thermal evaporation, 79 thermal expansion, 165, 167, 175, 228, 335 thermal mechanical analysis, 340 thermal plasma, 456 thermal properties, 480 thermal stability, 344, 348, 364, 500 thermal treatment, 158, 159, 162, 345, 537, 538 thermodynamic, 152, 251, 278, 397, 411, 561, 578 thermodynamic method, 278 thermodynamic parameters, 561 thermodynamic stability, 251, 411 thermodynamics, 247, 411, 493 thermogravimetry, 282 thermolysis, 323 thin film, vii, 3, 4, 5, 7, 24, 29, 40, 56, 57, 95, 159, 171, 181, 184, 207, 276, 277, 280, 300, 455, 456, 481, 488, 489, 523, 581 Thomson, 3 three-dimensional, viii, xiii, 17, 20, 51, 52, 410 threshold, 11, 186, 187, 248, 373, 384, 526, 569 thresholds, 186 thymocytes, 147 thyroid, 133 time consuming, 332 time resolution, 376 TiO2, viii, xiv, xvi, 4, 27, 29, 30, 31, 336, 337, 338, 339, 349, 350, 351, 353, 363, 414, 479, 480, 481, 482, 483, 484, 488, 489, 490, 491, 492, 493, 494, 495, 496, 497, 498, 500, 501, 502, 503, 504, 505, 506, 525, 527, 530, 538, 544, 545, 606 tissue engineering, 52, 149, 619 titania, 24, 137, 480, 481, 482, 484, 489, 490, 491, 618, 619 titanium, xv, 29, 30, 31, 108, 134, 135, 136, 137, 138, 139, 243, 310, 480, 505, 509, 522, 618 Titanium, xv, 29, 134, 135, 509, 511, 513, 515, 517, 519, 521, 523 titanium dioxide, 29, 618 tolerance, 128, 175 toluene, 36, 37, 326 top-down, vii, viii, xi, 3, 4, 23, 42, 293, 300 topographic, 73, 74, 87, 88, 373, 385 topological, 354, 372, 403 topology, 372, 376 toughness, x, 151, 161, 170, 172, 173, 174, 184, 185, 186, 194, 195, 196, 197, 199, 204, 205, 215, 216, 217, 224, 228, 455, 514 toxic, 109, 120, 121, 128, 133, 134, 138 toxic metals, 121 toxicity, 121, 127, 135, 148 TPA, 442 trajectory, 383 transducer, 370 transfer, vii, xi, xvi, 3, 11, 17, 27, 35, 38, 39, 115, 141, 275, 277, 278, 279, 282, 300, 302, 303, 309,

642

Index

310, 316, 324, 367, 447, 480, 484, 489, 491, 499, 505, 559, 560, 561, 562 transferrin, 149 transformation, xii, 91, 156, 159, 167, 228, 229, 318, 335, 339, 342, 371, 429, 453 transformation matrix, 91 transformations, 373 transforming growth factor, 110 transistors, 24, 29, 72, 101, 276, 299, 526 transition metal, 282, 315 transition metal ions, 315 transition temperature, 152, 335, 368, 461 transitions, xiv, 263, 264, 467, 479, 482, 486, 489, 490, 498, 500, 501, 502 translational, 562, 572, 574, 575, 578, 583 transmission, xi, xiv, xv, 32, 75, 77, 84, 91, 92, 114, 167, 237, 321, 373, 375, 412, 459, 479, 509, 512, 526, 527, 533, 535, 592, 605, 606, 617, 618 transmission electron microscopy, xi, xiv, xv, 114, 237, 321, 412, 459, 526, 527, 535, 605, 606 transparent, 28, 29, 30, 239, 265, 342, 436, 461, 489, 618 transport, x, 7, 14, 23, 30, 41, 111, 114, 144, 152, 170, 238, 294, 372, 404, 410, 414, 429, 454 transpose, 90 travel, 278, 377, 400 travel time, 377, 400 treatment methods, ix, 107, 111 trend, 56, 137, 210, 276, 398, 514, 519 TRF, 326 trifluoroacetic acid, 32 Tsunami, 493 tubular, 42, 578 tumor, ix, 107, 108, 110, 111, 112, 113, 114, 115, 116, 118, 119, 120, 124, 125, 126, 127, 128, 141, 143, 144, 145, 146, 147, 150 tumor cells, ix, 107, 108, 110, 112, 113, 114, 120, 126, 146 tumor growth, 116 tumor progression, 143 tumors, 108, 109, 110, 111, 112, 115, 116, 119, 122, 124, 125, 127, 128, 142, 143, 145 tungsten, 20, 173, 333, 334, 335, 338, 339, 343, 353, 361, 392, 456, 538 tungsten carbide, 173, 333, 334, 335, 338, 339, 343, 361 tunneling, 11, 41, 265, 374 twins, 165, 166, 224, 225, 416, 456, 555 two-dimensional (2D), viii, xiii, 9, 51, 156, 239, 240, 245, 409, 410, 560

U Ukraine, 371 ultra-fine, 52, 343 ultrasound, 115, 127 ultraviolet, 102, 121, 239, 248, 269, 282, 482, 546, 615 Ultraviolet, 82, 618

uncertainty, 175, 179, 257, 389 underreported, 108 uniform, ix, xiii, 9, 17, 22, 23, 25, 32, 36, 37, 58, 63, 71, 73, 87, 131, 238, 252, 253, 299, 311, 345, 358, 359, 374, 394, 409, 410, 420, 431, 450, 517, 520, 521, 606, 612 unions, 301 urethane, 144 urokinase, 110 UV, ix, 40, 72, 73, 121, 239, 240, 248, 258, 259, 260, 261, 264, 282, 283, 284, 287, 290, 294, 295, 297, 300, 303, 484, 532, 538, 539, 546, 551, 615, 616, 617 UV absorption, 532, 539 UV exposure, 121 UV irradiation, 283, 284, 290 UV light, ix, 72, 258, 282, 283, 284, 300, 303 UV radiation, 121, 616, 617

V vacancies, 161, 258, 302, 403, 427, 503, 540, 541, 549, 551 vacuum, xv, 53, 65, 73, 93, 266, 276, 350, 374, 380, 396, 399, 410, 509, 511, 522, 527, 546 valence, 263, 265, 282, 301, 349 validation, 148 validity, 376, 397 van der Waals, 410, 560, 570 van der Waals forces, 570 vapor, 239, 241, 242, 243, 245, 247, 250, 266, 269, 277, 376, 392, 454, 455, 456, 457, 527, 546, 606, 611 vapor phase deposition, 247 vapor-liquid-solid, 238, 241, 611 variable, 4, 6, 341, 392, 399 variables, 335 variation, ix, 27, 28, 52, 71, 72, 80, 95, 97, 100, 175, 191, 351, 353, 395, 473, 494, 514, 544, 554 vascular endothelial growth factor, 110, 111 vascular endothelial growth factor (VEGF), 111 vascularization, 110, 111 vasculature, 112, 142 vector, 90, 120, 226, 468, 564, 572 VEGF, 111 velocity, 168, 253, 265, 351, 437, 438 versatility, 15, 109, 128 vessels, 110, 111, 118, 144, 147 vibration, 74, 75, 80, 84, 87, 175, 255, 334 vibrational modes, 76, 84, 416 Vickers hardness, 510 vinblastine, 133 viscosity, 52, 53, 57, 66, 164, 168 visible, 29, 41, 73, 103, 117, 121, 250, 258, 261, 265, 282, 420, 426, 489, 504, 515, 523, 538, 539, 540, 546, 562, 588, 596 vision, 149 visualization, 184 VLS, 238, 241, 242, 243, 611

643

Index voiding, 5 voids, 36, 37, 414, 417, 421, 422, 534, 549 volatility, 332 vortex, 561, 590 vortices, 158

X-ray analysis, 412 X-ray diffraction, xi, xv, 194, 237, 412, 413, 439, 447, 479, 526, 527, 531, 532 X-ray diffraction (XRD), xiv, 249, 250, 251, 252, 253, 254, 287, 288, 290, 336, 337, 338, 339, 341, 343, 344, 346, 350, 412, 413, 439, 442, 446, 450, W 479, 485, 487, 488, 492, 493, 494, 495, 496, 505, 512, 516, 518, 519, 520, 521, 523, 526, 527, 531, waste products, 110 532, 533, 534, 539, 606, 607, 609, 611 water, 10, 11, 14, 17, 41, 53, 117, 125, 128, 129, 246, X-ray photoelectron spectroscopy (XPS), 526 249, 252, 282, 317, 318, 319, 328, 414, 420, 421, Y 460, 461, 485, 486, 492, 493, 560, 564, 565, 566, 571, 575, 579, 591, 615, 616 water evaporation, 414 Y-axis, 587 water-soluble, 114 YBCO, x, 151, 152, 153, 154, 156, 158, 159, 160, wave vector, 257, 468, 572 162, 164, 165, 167, 168, 170, 171, 184, 194, 195, wavelengths, 82, 83, 117, 118, 282, 300, 486, 592 196, 197, 198, 199, 200, 201, 202, 206, 207, 208, wealth, 10 209, 210, 211, 212, 213, 214, 215, 216, 217, 218, wear, 186, 510, 512, 515, 516 219, 221, 222, 223, 224, 225, 226, 228, 229, 230 Weibull, 213 yeast, 133 weight loss, 493 yield, x, xv, 8, 9, 31, 151, 158, 172, 174, 175, 178, weight ratio, 335, 341, 351 182, 188, 191, 192, 211, 218, 221, 222, 224, 243, wettability, 319, 429 245, 355, 374, 410, 480, 605, 606, 607, 612 WG, 144 YSZ, xvi, 525, 527, 529, 530, 534, 535, 536, 539, wide band gap, 539, 546 540, 544, 551, 552, 553 windows, 453 yttria-stabilized zirconia, 527 wires, xii, 9, 10, 17, 24, 40, 371, 372, 381, 388, 394, yttrium, 156, 157 396, 397, 401, 402, 403, 578

Z

X xenograft, 116, 146, 150 xenon, 493, 497, 498 XPS, 135, 136, 261, 262, 263, 282, 283, 412, 526, 527, 528, 529, 530, 531, 534, 539, 540, 541 X-ray absorption, 282

zinc (Zn), viii, 27, 51, 53, 67, 242, 244, 245, 246, 247, 248, 250, 252, 255, 258, 269, 294, 310, 343, 362, 366, 546 zirconia, 364, 527 ZnO nanorods, 245, 246, 248, 249, 256, 261, 263 ZnO nanostructures, 240, 245, 253, 255