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.
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(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. φand 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.
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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)
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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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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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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
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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.
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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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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.
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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.
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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.
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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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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 45. Load/depth-of-penetration results for YBCO monodomain ab-planes, with a 25 μm spherical indenter, for Bridgman technique at 100 nm of penetration depth.
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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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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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
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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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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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