Encyclopedia of nanoscience and nanotechnology vol 10 edited by hari singh nalwa

To support this statement, we will focus on the discussion of the implications of the particularities of the intergranular coupling on the hysteretic behavior of nanocrystalline materials First, we will cover the case of the (mainly single phase) reduced magnetocrystalline anisotropy materials, in which the dipolar correlation length frequently can be much larger than the crystallite size (ca 15 nm), thus resulting in an extremely soft behavior linked to the occurrence of exchange-induced averages of the local anisotropy The second section will be dedicated to the demagnetization process of high anisotropy, single- and multiphase nanocrystalline materials characterized by exchange correlation lengths comparable with or smaller than the crystallite size In this case, the goal is either the reduction of the intergranular coupling (single-phase materials) aiming at the increase of coercivity or the achievement of large remanences linked to the occurrence of strong coupling between hard and soft grains, the latter having dimensions comparable with their exchange correlation length We will end this chapter by reviewing the state of the art of the micromagnetic modeling, a numerical technique allowing both the analysis of systems for which (as it is the largely majority case today) there are not experimental data on the local magnetic properties and the implementation of elements of device design This last section also will include a discussion of the influence on the performance of magnetic recording media of the control of the intergranular exchange SOFT NANOCRYSTALLINE MATERIALS 2.1 Introduction to the Soft Magnetic Materials Nanocrystalline materials, obtained by devitrification of the precursor amorphous alloy, displaying soft magnetic character (high magnetic permeability and low coercivity), have been the subject of increasing attention from the scientific community, not only because of their potential use in technical applications but also because they provide an excellent setting in which to study basic problems in nanostructures formation and magnetism [1–8] In fact, these materials provide a crucial point in opening up new fields of research in materials science, magnetism, and technology, such as metastable crystalline phases and structures, extended solid solubilities of solutes with associated improvements of mechanical and physical properties, nanocrystalline, nanocomposite and amorphous materials that, in some cases, have unique combinations of properties (magnetic, mechanical, corrosion, etc.) Technological development of the fabrication technique of the amorphous precursor material and studies of the structure, glass formation ability, and thermodynamics, and magnetism of amorphous alloys were intensively performed in 1960s and 1970s These aspects have been analyzed extensively in few review papers and books [9–11] Most commercial and technological interests have been paid to soft amorphous and nanocrystalline magnetic materials Initially, it was believed that ferromagnetism could not Soft and Hard Magnetic Nanomaterials exist in amorphous solids because of a lack of atomic ordering The possibility of ferromagnetism in amorphous metallic alloys was theoretically predicted by Gubanov [12] and the experimental confirmation of this improbable prediction was the main cause of the sudden acceleration of research on amorphous alloys from about 1970 onward, this onrush of activity was due both to the intrinsic scientific interest of a novel and unexpected form of ferromagnetism and also to the gradual recognition that this is the key to the industrial exploitation of amorphous ferromagnetic alloys The amorphous alloy ribbons obtained by the melt-spinning technique have been introduced widely as the soft magnetic materials in the 1970s Their excellent magnetic softness and high wear and corrosion resistance made them very attractive in the recording head and microtransformer industries In contrast with the flood of work on magnetic behavior, the study of electrical transport (i.e., magnetoimpedance effect) is very recent and is making significant progress Conventional physical metallurgy approaches to improving soft ferromagnetic properties involve tailoring the chemistry and optimizing the microstructure Significant in the optimizing of the microstructure is recognition that a measure of the magnetic hardness (the coercivity, Hc ) is roughly inversely proportional to the grain size D for a grain size exceeding ∼0 to m (where the grain size exceeds the domain wall thickness) In such cases, grain boundaries act as impediments to domain wall motion, and, thus, finegrained materials usually are magnetically harder than large grain materials Significant developments in the understanding of magnetic coercivity mechanisms have lead to the realization that for very small grain size D < ∼100 nm [13–21], Hc decreases rapidly with increasing grain size This can be understood by the fact that the domain wall, whose thickness exceeds the grain size, now samples several (or many) grains so that fluctuations in magnetic anisotropy on the grain-size length scale are irrelevant to domain wall pinning This important concept suggests that nanocrystalline and amorphous alloys have significant potential as soft magnetic materials In this section, we explore issues that are pertinent to the general understanding of the magnetic properties of amorphous and nanocrystalline materials As the state of the art for amorphous magnetic materials is well developed and much of which has been thoroughly reviewed [11, 22–24], we will concentrate on highlights and recent developments The development of nanocrystalline materials for soft magnetic applications is an emerging field for which we will try to offer a current perspective that may well evolve further with time The development of soft magnetic materials for applications requires the study of a variety of intrinsic magnetic properties as well as development of extrinsic magnetic properties through an appropriate optimization of the microstructure As intrinsic properties, we mean microstructure insensitive properties Among the fundamental intrinsic properties (which depend on alloy composition and crystal structure), the saturation magnetization, Curie temperature, magnetic anisotropy, and magnetostriction coefficient are all important In a broader sense, magnetic anisotropy and magnetostriction can be considered as extrinsic in that, Soft and Hard Magnetic Nanomaterials 2.2 Microstructural Characterization –1 1W·g An important part of the recent developments corresponding to nanostructured materials is related to those obtained by controlled crystallization, either by annealing the amorphous single phase or by decreasing the cooling rate from the liquid of metallic systems Typically, in these nanostructures, precipitate sizes range between and 50 nm embedded in an amorphous matrix with nanocrystal volume fractions of 10 to 80%, which means particle densities of 1022 to 1028 m−3 We present the most relevant aspects of the nanocrystallization process of Fe-based nanocrystalline alloys as a two-phase system, namely -Fe or -Fe(Si) grains embedded in a residual amorphous matrix, which, being ferromagnetic, results in a material with extremely good soft magnetic properties The microstructural analysis of the primary crystallization of Fe-rich amorphous alloys usually has been done by using conventional techniques such as differential scanning calorimetry (DSC), X-ray diffraction (XRD), transmission Mössbauer spectroscopy (TMS), and transmission electron microscopy (TEM) In this way, through the combined structural analysis of these techniques, useful information, such as the dependence of the microstructure upon time and temperature of treatment, can be obtained The kinetics of metastability loss of the disordered system above glass transition, i.e., under less than equilibrium conditions, is a key subject, because it provides new opportunities for structure control by innovative alloy design and processing strategies Several examples include soft and hard magnets and high-strength materials [30–32] Most studies focus on the crystallization onset as a measure of kinetic stability under heat treatment and recognise the product phase selection involved in nucleation and the role of competitive growth kinetics in the evolution of different microstructural morphologies [33, 34] Differential scanning calorimetry has become quite effective as a means of studying the nature of nanoscale structures and their stability It has been established that the initial annealing response allows one distinguish between a sharp onset for a nucleation and growth or a continuous grain growth of pre-existing grains [35] Kinetic data on the transformation often are obtained from this technique Figure shows the DSC curves of the as-prepared amorphous alloy (Fe73 Cu1 Nb3 Si17 B5 , trademark Finemet), as well as of that of alloy previously annealed for hour at 703 K and 763 K, respectively [36] For the as-prepared alloy, the calorimetric signal shows some relaxation before the exothermic nanocrystallization process, as well as with the qmc Curie temperature of the amorphous phase (TC ≈ 595 K) When the sample has been annealed at 703 K, relaxation is no longer apparent in the calorimetric signal, the Curie temperature is shifted about 15 K, to higher values, and the nanocrystallization process is slightly advanced in temperature (2 K) On further increasing the annealing temperature, the calorimetric signal shows no clear changes in the Curie temperature of the annealed sample with respect to annealing at 703 K However, a clear shift of the nanocrystallization onset toward higher temperatures (40 K) can be observed, as well as a significant decrease of its area [36] Consequently, DSC measurement permits evaluation of the maximum temperature to prevent partial crystallization during previous heating and the maximum heating rate to control the temperature of the sample by analyzing the transformation peak connected to the primary crystallization Calorimetric measurements, however, give information about microstructural development Microstructure determination by the use of several techniques (XRD, TMS and TEM) allows us to complete the understanding of the mechanisms of the primary crystallization Transmission Mössbauer spectroscopy has the main advantage of giving local information of an active element (Fe nuclei in these alloys) Because Fe is present both in the nanocrystalline precipitates and in the disordered matrix, TMS provides information on local ordering in both phases, which can be correlated with the changes in the short-range order of the amorphous phase and with the composition of nanocrystalline phase X-ray diffraction and TEM provide a Heat flow / W·g–1 for a two-phase material (in aggregate), they depend on the microstructure A vast literature exists on the variation of intrinsic magnetic properties with alloy composition Although new discoveries continue to be made in this area, it can be safely stated that a more wide open area in the development of magnetic materials for applications is the fundamental understanding and exploitation of the influence of the microstructure on the extrinsic magnetic properties Important microstructural features include grain size, shape, and orientation; defect concentrations; compositional inhomogeneities; magnetic domains; and domain walls The interaction of magnetic domain walls with microstructural impediments to their motion is of particular importance to the understanding of soft magnetic behavior Extrinsic magnetic properties important in soft magnetic materials include magnetic permeability and coercivity, which typically have an inverse relationship Thorough discussions of soft magnetic materials are available [25–28, 29] as prepared after h annealing at 703 K after h annealing at 763 K 600 650 700 750 800 850 900 T (K) Figure DSC signals of Fe73 Si13 B9 Cu1 Nb3 alloy obtained after heating at 40 K/min For the as-quenched amorphous alloy (solid line); after h isothermal annealing at 703 K (dashed line) and at 763 K (dotted line) Reprinted with permission from [37], M T Clavaguera-Mora et al., Progress in Materials Science 47, 559 (2002) © 2002, Elsevier Science Soft and Hard Magnetic Nanomaterials close look at the developed microstructure and permit the characterization of the precipitates, showing their morphology and grain-size distribution Figure 2(a–c) shows a TEM image of a Finemet alloy after heating at 763 K at three annealing times (3, 10, and 30 minutes), while Figure 2(d) corresponds to the case (b) with increasing contrast, which allows the shape and size of the precipitates to be determined [37] As can be seen, the microstructure is characterized by a homogeneous, ultrafine grain structure of -Fe(Si), with grain sizes around 10 nm and a random texture, embedded in a still amorphous matrix The formation of this particular structure is ascribed to the combined effects of elements as Cu (which promotes the nucleation of grains) and Nb, Ta, Zr, Mo,—(which hinders their growth) and their low solubility in -Fe(Si) [38–40] Nevertheless, the size and morphology of the nanocrystals in these alloys, as well as their distribution, could be analyzed by the application of local probe techniques These techniques, such as scanning tunneling microscopy (STM) and atomic force microscopy (ATM) provide threedimensional (3D) topographic images at the nanometer level [41, 42] and represent powerful tools to study the surface properties and structures of metals and alloys 2.3 On the Ferromagnetism in Amorphous and Nanocrystalline Materials There are, when concerning the magnetic order in materials having structural disorder (such as is the case of amorphous and nanocrystalline alloys), some fundamental questions related to the existence of such a well-defined magnetic order Ferromagnetic interactions of the magnetic materials can be immediately considered as ferromagnetic structures In this naive idea, the magnetic anisotropy effects have been neglected Magnetic moments tend to arrange their orientations parallel to each other via exchange interactions; they (a) (b) (c) (d) Figure TEM images of Finemet samples after (a) min; (b) 10 min; (c) annealing at 763 K; and (d) is the case (b) after contrast increasing Reprinted with permission from [37], M T Clavaguera-Mora et al., Progress in Materials Science 47, 559 (2002) © 2002, Elsevier Science this when lying along a magnetic easy axis that is in the same direction at every point in the material However, if the easy axis orientation fluctuates from site to site, a conflict between ferromagnetic coupling and anisotropy arises As long as we imagine lattice periodicity, a ferromagnetic structure is a consequence of ferromagnetic exchange interactions, the strength of the anisotropy being irrelevant In this situation, we are assuming a major simplification, namely, the direction of the easy axis is uniform throughout the sample With this simple picture, we present crucial questions related to the influence of an amorphous structure on magnetic order Regarding the magnetic order in amorphous and nanocrystalline materials, we know that it stems from two contributions: exchange and local anisotropy The exchange arises from the electron–electron correlations The mechanism of the electrostatic interactions between electrons has no relation to structural order and is sensitive only to overlapping of the electron wave functions With respect to magnetic anisotropy, it also originated by the interaction of the local electrical field with spin orientation, through the spin-orbit coupling Therefore, magnetic anisotropy also is a local concept Nevertheless, the structural configuration of magnetic solids exerts an important influence on the macroscopic manifestation of the local anisotropy As a consequence, when the local axes fluctuate in orientation owing to the structural fluctuation (amorphous and nanocrystalline materials as examples), calculations of the resultant macroscopic anisotropy become quite difficult In the case of amorphous ferromagnetic alloys, the usual approach to the atomic structure of a magnetic order connected to a lattice periodicity is not applicable These materials can be defined as solids in which the orientation of local symmetry axes fluctuate with a typical correlation length l = 10 A The local structure can be characterized by a few local configurations with icosahedral, octahedral, and trigonal symmetry These structural units have randomly distributed orientation The local magnetic anisotropy would be larger in the units with lower symmetry In general, these units are characterized by fluctuations of the orientation local axis It is remarkable that with these types of structures, the correlation length, l, of such a fluctuation is typically the correlation length of the structure and ranges from 10 A (amorphous) to 10 nm (nanocrystals) and mm (polycrystals) Fluctuations in the interatomic distances associated with the amorphous structure also should contribute to some degree of randomness in the magnetic interactions of the magnetic moments Nevertheless, such randomness is expected not to affect the magnetic behavior qualitatively [11, 43] Moreover, random distribution of the orientation of the easy axis drastically affects the magnetic properties The random anisotropy model developed by Alben et al [44] provides a successful explanation of how the correlation length, l, exerts a relevant influence on magnetic structure The important question is What is the range of orientational correlation of the spins? Let L be the correlation length of the magnetic structure If we assume L > l, the number of oriented easy axes in a volume L3 should be N = L/l The effective anisotropy can be written as: Keff = K/N 1/2 (1) Soft and Hard Magnetic Nanomaterials where K is the local anisotropy where strength is assumed to be uniform everywhere By minimizing the total energy with respect to L, the following expression can be deduced: L = 16A2 / 9K l3 (2) where A is the exchange stiffness parameter If we consider A = 10−11 J/m and l = 10−9 m, which are typical values of ferromagnetic metallic glasses [45], L in equation (2) becomes 105 /K For 3D-based alloys, we can take the value of K corresponding to crystalline samples (∼104 J/m3 ) leading to L around 10−9 m, which is equal to the structural correlation length of an amorphous material In addition, the random anisotropy model provides the following expression for the average macroscopic anisotropy: K = K l6 /A3 (3) Equation (3) points out that the macroscopic structural anisotropy is negligible in 3D amorphous alloys K ∼ 10−9 K); this is a consequence of the averaging of several local easy axes, which produces the reduction in magnitude Special attention has been paid, in the last decade, to the study of nanocrystalline phases obtained by suitable annealing of amorphous metallic ribbons owing to their attractive properties as soft magnetic materials [1, 15, 19–21, 23, 46–48] Such soft magnetic character is thought to have originated because the magnetocrystalline anisotropy vanishes and there is a very small magnetostriction value when the grain size approaches 10 nm [1, 12, 46] As was theoretically estimated by Herzer [12, 46], average anisotropy for randomly oriented -Fe(Si) grains is negligibly small when grain diameter does not exceed about 10 nm Thus, the resulting magnetic behavior can be well described with the random anisotropy model [12, 19, 23, 46–48] According to this model, the very low values of coercivity in the nanocrystalline state are ascribed to small effective magnetic anisotropy (Keff around 10 J/m3 However, previous results [19, 21, 49] as well as recently published results by Varga et al [50] on the reduction of the magnetic anisotropy from the values in the amorphous precursors (∼1000 J/m3 ) down to that obtained in the nanocrystalline alloys (around 300–500 J/m3 ), is not sufficient to account for the reduction of the coercive field accompanying the nanocrystallization The enhancement of the soft magnetic properties should, therefore, be linked to the decrease of the microstructure– magnetization interactions These interactions, originating in large units of coupled magnetic moments, suggest a relevant role of the magnetostatic interactions, as well a role in the formation of these coupled units [19, 49] In addition to the suppressed magnetocrystalline anisotropy, low magnetostriction values provide the basis for the superior soft magnetic properties observed in particular compositions Low values of the saturation magnetostriction are essential to avoid magnetoelastic anisotropies arising from internal or external mechanical stresses The increase of initial permeability with the formation of the nanocrystalline state is closely related to a simultaneous decrease of the saturation magnetostriction Partial crystallization of amorphous alloys leads to an increase of the frequency range, where the permeability presents high values [51] These high values in the highest possible frequency range are desirable in many technological applications involving the use of ac fields It is remarkable that a number of workers have investigated the effects on the magnetic properties of the substitution of additional alloying elements for Fe in the Fe73 Cu1 Nb3 Si13 B9 alloy composition, Finemet, to further improve the properties, e.g., Co [52–56], Al [20, 57, 58], varying the degree of success Moreover, it was shown in [20] that substitution of Fe by Al in the classical Finemet composition led to a significant decrease in the minimum of coercivity, Hcmin ≈ A/m, achieved after partial devitrification, although the effective magnetic anisotropy field was quite large (around Oe) [59] The coercivity behavior was correlated with the mean grain size, and a theoretical low effective magnetic anisotropy field of the nanocrystalline samples was assumed in contradiction with those experimentally found in [49, 50, 58] Although amorphous Fe-, Co-, and Ni-based ribbons are slightly more expensive compared with conventional soft magnetic materials, such as sendust, ferrites, and supermalloys (mostly due to the significant content of Co and Ni), they found considerable applications in transformers (400 Hz), ac powder distributors (50 Hz), magnetic recording as a magnetic heads, and magnetic sensors The main reason for using amorphous alloys such as soft magnetic materials is a saving of the electric energy wasted by magnetic cores Besides, the combination of high magnetic permeability and good mechanical properties of amorphous alloys may be used successfully in magnetic shielding and in magnetic heads [51] Production of about millions heads per year in Japan in the mid-1980s has been reported [51] The internal stresses, as the main source of magnetic anisotropy in amorphous and nanocrystalline materials, are due to the magnetoelastic coupling between magnetization and internal stresses through magnetostriction Consequently, these materials are interesting for field- and stresssensing elements because the Fe-rich amorphous alloys exhibit high magnetostriction values ( s ≈ 10−5 ) and, therefore, many of magnetic parameters (i.e., magnetic susceptibility, coercive field, etc.) are extremely sensitive to the applied stresses The discovery of Fe-rich nanocrystalline alloys carried out by Yoshizawa et al [1] was important owing to the outstanding soft magnetic character of such materials Typical compositions of the precursor amorphous alloys, which, after partial devitrification, reach the nanostructure character with optimal properties, are FeSi and FeZr, with small amounts of B to allow the amorphization process, and smaller amounts of Cu, which act as nucleation centers for crystallites, and Nb, which prevents grain growth This effect is provided by Zr in FeZr alloys After the first step of crystallization, FeSi or Fe crystallites are finely dispersed in the residual amorphous matrix In a wide range of crystallized volume fraction, the exchange correlation length of the matrix is larger than the average intergranular distance, d, and the exchange correlation length of the grains is larger than the grain size, D Magnetic softness of Fe-rich nanocrystalline alloys is due to a second complementary reason: the opposite sign of the magnetostriction constant of crystallites and residual amorphous matrix, Soft and Hard Magnetic Nanomaterials which allows reduction and compensation of the average magnetostriction Figure shows the thermal variation of the coercive field (Hc ) in a Finemet-type (Ta-containing) amorphous alloy This behavior is quite similar to that shown in the case of Nbcontaining ones and particularly, evidences the occurrence of a maximum in the coercivity linked to the onset of the nanocrystallization process [60, 61] Considering the grain size, D, to be smaller than the exchange length, Lex , and the nanocrystals are fully coupled between them, the random anisotropy model implies a dependence of the effective magnetic anisotropy K , with the sixth power of average grain size, D The coercivity is understood as a coherent rotation of the magnetic moments of each grain toward the effective axis leading to the same dependence of the coercivity with the grain size [16]: Hc = pc K K D6 = pc Js Js A K = K14 D A3 (4) where K1 = kJ/m3 is the magnetocrystalline anisotropy of the grains, A = 10−11 J/m is the exchange ferromagnetic constant, Js = T, is the saturation magnetic polarization and pc is a dimensionless prefactor close to unity The predicted D6 dependence of the coercive field has been widely accepted to be followed in a D range below Lex (around 30 to 40 nm) for nanocrystalline Fe–Si–B–M–Cu (M = Iva to Via metal) alloys [21, 46, 62–65] A clear deviation from the predicted D law in the range below Hc = A/m was reported by Hernando et al [19] Such deviation was ascribed to effects of induced anisotropy (e.g., magnetoelastic and field induced anisotropies) on the coercivity could be significant with respect those of the random magnetocrystalline anisotropy As a consequence, the data of Hc D were fitted by assuming a contribution from (i) the spatial fluctuations of induced anisotropies and (ii) Ku to a2 + bD was found with Hc (i.e., a dependence Hc = a = A/m representing the contribution originating from the induced anisotropies) To investigate the effect of the grain size on coercivity, this dependence of Hc D in alloys treated by Joule heating was obtained Experimental results on this dependence are shown in Figure [21] The fitting of this dependence appears to follow, surprising, a rough dependence of the Hc ∝ D3−4 type (the best regression was found fitting the D law) It must be noted that our data of Hc D correspond to a grain size variation between about 10 to 150 nm As it is well known, an analysis of this Hc D data in terms of the random anisotropy model is only justified if the grain size is smaller than Lex and, hence, could not be applicable (in the framework of the random anisotropy model) to the range grain size above Lex , which results in being only two points of our data in Figure [21] These points should correspond to a magnetic hardening due to the precipitation of the iron borides In this case, the random anisotropy model should be applied by taking into account the volume fraction and the different anisotropy of the iron borides This indicates that K1 should vary as D3 , contrary to the theoretical D6 law This indicates that Hc is mainly governed by K1 , which varies as D3 , contrary to the theoretical D6 law This contradiction of the Hc D law between the theory and the experimental has recently been explained by Suzuki et al [62] considering the presence of long-range uniaxial anisotropy, Ku , which influences the exchange correlation value and length, and yields an anisotropy average given by: K = Ku + · Ku · K1 · D3 A3/2 (5) The second part of (5) corresponds to K1 and if Ku is coherent in space or if its spatial fluctuations are negligible to K1 , this second part ultimately determines the grain size influence on the coercivity Such influence changes from the D6 law to a D3−4 one when the coherent uniaxial anisotropies dominate over the random magnetocrystalline anisotropy An additional point in order to justify the Eq (5) 100 10 Hc (A/m) Hc (A/m) 103 10 102 10 Fe73.5Si13.5B9Cu1Ta3 1 10 20 30 40 50 60 Jann (A/mm2) Figure Evolution of the coercive field, measured at room temperature, as a function of the current density after: (o) and ( ) 10 of annealing time Reprinted with permission from [21], N Murillo and J González, J Magn Magn Mater 218, 53 (2000) â 2000, Elsevier Science 0.1 101 Hc - 2.9ì10 D (nm) –4 3.35 D 102 Figure Dependence of the coercive force, Hc , with the average grain diameter, D, for the two studied compositions (Fe73 Si13 B9 Cu1 Nb3 and Fe73 Si13 B9 Cu1 Ta3 Reprinted with permission from [21], N Murillo and J González, J Magn Magn Mater 218, 53 (2000) © 2000, Elsevier Science Encyclopedia of Nanoscience and Nanotechnology www.aspbs.com/enn Zinc Oxide Nanostructures Chun-Hua Yan, Jun Zhang, Ling-Dong Sun Peking University, Beijing, People’s Republic of China CONTENTS Introduction Fabrication and Morphology of ZnO Nanostructures Optical Properties of ZnO Nanostructures Conclusion Glossary References INTRODUCTION Nanostructured semiconductors, in particular their versatile fabrications and unique chemical or physical properties for potentially technological applications, have been stimulating considerable research interests in the past decade [1] Tremendous progress has been made to understand the quantum-size behaviors and to investigate the size- and morphology-dependent properties With this respect, great efforts and contributions from the fruitful groups, such as Yang’s [2] and Wang’s groups [3] for ZnO nanostructures, have been made persistently to enrich the diversiform morphological world of nanostructures and show their possibilities for versatile utilizations in room temperature laser and other highly technological fields It is demonstrated so much that the properties of materials on nanoscale, especially those with one dimension less than 10 nm, are virtually dependent on their sizes and morphologies [4] We therefore reasonably need to summarize the achievements obtained so far in the field of fabrication and morphology control of the most focused nanostructures of ZnO ZnO, known as a very important semiconductor with wide bandgap (3.37 eV) and large exciton binding energy of 60 meV [5] at room temperature, has been investigated extensively due to its wide technological applications from catalytic, electrical, optoelectronic, and photochemistry fields [6] to the room temperature blue-ultraviolet (UV) laser region [7] For instance, ZnO as a gas sensing materials is sensitive to many sorts of gases with satisfactory stability [8], while as a promising material for dye-sensitized solar cells, ZnO demonstrates improved performance [9] In the form of thin film, ZnO is a potential candidate for flat ISBN: 1-58883-066-7/$35.00 Copyright © 2004 by American Scientific Publishers All rights of reproduction in any form reserved display screen usage [10], and a recent report reveals that ZnO would be an alternative to the TiO2 nanostructured electrode in Grützel-type photocells [11] Moreover, ZnO is also regarded as a hopeful material for realizing efficient light-emitting diodes, UV light-emitting diodes [12], and room temperature lasers, which has been demonstrated by several groups [2g, 2h, 13] In addition, ZnO can be made as transparent and highly conductive, or piezoelectric components as well [14] Therefore, studies with respect to synthesis and properties of ZnO nanostructures have been presented to meet the demands of potential applications on morphology of materials So far, synthetic strategies including physical process-related methods such as chemical vapor deposition [15], thermal evaporation [16], vapor deposition [17], thermal decomposition [18], arc plasma [19], sputtering and laser ablation [13a, 20], molecular-beam epitaxy [21], hydrolysis [22], electrochemical deposition [23], and chemical route such as hydrothermal and solvothermal [24], solid state reaction [25], sol–gel [26], precipitation [27], templated-based and solution-based processes (discussed in detail later) have been employed to manufacture ZnO nanostructures Well-defined nanostructures of ZnO with abundant morphological forms such as nanorods [28], nanobridges and nanonails [29], nanobelts [3b, 3h], hierarchical nanostructures [30], nanocables and nanotubes [3e], nanoneedle [31], and dendritic nanowires [32] have been achieved and other unusual morphological forms controlled as flower-, snowflake-, prism-, and prickly sphere-like shape have been presented as well [33] The morphogenesis shows the possibility of controlling the morphology of ZnO nanostructrues by employing suitable synthetic process To give an insight into the capability of controlling the size and morphology of ZnO nanostructures, it is desirable to integrate the knowledge of all aspects of the morphology controlling preparation We hereby present this chapter and try to highlight fabrication and control strategies of ZnO nanostructures based on our efforts on morphologically controllable synthesis by solution-based routes [33a, 34] For the purposes of this chapter, the synthetic methods can be tentatively classified into several categories according to a previous publication [3b]: vapor-phase growth including thermal evaporation, chemical vapor-phase deposition, metal-organic chemical vapor-phase deposition, arcdischarge, laser ablation, templated-based, sol–gel, and solution-based growth methods, etc Among these synthetic Encyclopedia of Nanoscience and Nanotechnology Edited by H S Nalwa Volume 10: Pages (767–780) 768 processes, the solution-based approach may be the simplest, without involving high temperature and expensive equipment, but powerful enough to lead to morphology control by adjusting the reaction conditions This chapter will mainly review recent progress on synthesis and morphology control of ZnO by the solution-based route FABRICATION AND MORPHOLOGY OF ZnO NANOSTRUCTURES 2.1 Early Synthetic Route ZnO being a useful semiconductor has been studied for a long time Long before, attention was mostly paid to the synthesis of ZnO nanocrystallites and nanocrystalline epitaxial films, in which ultraviolet emission at room temperature was observed [2h, 13d, 13g, 13i] Physical methods such as thermal decomposition, vapor evaporation, pyrolysis, laser ablation, arc plasma, etc and chemical techniques such as hydroor solvothermal, microemulsion, precipitation, sol–gel, etc are popular ways used to achieve high quality ZnO crystals For example, Meulenkamp et al [35a] employed sol– gel method via the addition of LiOH into ethanolic solution of zinc acetate to prepare morphological well-defined ZnO nanoparticles By finely controlling the water content and acetate anions, the nanoparticles size can be well controlled Later, more combined methods, for instance a hydrothermal coupled with microemulsion method, were developed to prepare nanocrystallites with controllable morphology, which show novel optical and electrical properties attributed to the size reduction and the accordant quantum confinement effects By high temperature physical methods, nano- or microwiskers are easy to obtain but mainly on the micrometer scale and with a wide size distribution [36] In contrast, by wet chemical routes, more morphological forms rather than nanoparticles were achieved However, less effort has been made on morphology controllable synthesis purposively Chittofrati et al [33b] has achieved ZnO whiskers and intertwined morphologies by hydrothermally treating zinc nitrate solution in presence of alkali such as NaOH, KOH, NH4 OH, triethanolamine and ethylenediamine at elevated temperature Wang et al [37] fully investigated the growth mechanism of bipyramidal dumbbell-like twinning morphology of ZnO generated by the hydrothermal method, indicating that growth conditions may have crucial effects on final morphologies of ZnO, which preliminarily provides a possibility for morphology control of ZnO A chemical precipitation method [38] was adopted to produce rodlike and other unusual morphologies of ZnO Summarily, it is noticeable that early synthetic methods may focus on the preparation of spherical and rodlike related morphology of ZnO, which might be attributed to the inadequate understanding of the diversity of materials morphology and of its importance for finely tuning properties 2.2 Vapor-Phase Growth Vapor-phase growth is one kind of approach involving a high temperature process for synthesizing nanomaterials Several methods can be assigned to this category, such as thermal evaporation [16], chemical vapor deposition [15], Zinc Oxide Nanostructures and metal organic chemical vapor deposition [17i] However, in consideration of morphology control, the innovative vapor deposition process and thermal evaporation method recently developed by Yang et al and Wang et al have been proved to be the two most effective approaches so far to obtain uniform ZnO nanostructures Generally, vapor deposition is carried out by heating the mixture of ZnO and graphite powder to high temperature, normally around 1000 C under a constant flow of argon, to make ZnO nanowires grow on a Au-coated silicon substrates This method is employed by Yang et al to prepare uniform ZnO nanowire arrays via a vapor–liquid–solid growth mechanism, in which the natural laser cavities are formed to give room temperature lasing action With the catalyst of Au film on substrate, the diameter of ZnO nanowires with growing orientation along the c-axis can be controlled as 20–150 nm and length up to 40 m This process provides the possibilities to get shortwavelength nanolasers for potential applications in optical computing, information storage, and nanoanalysis [2] The thermal evaporation process is in principle one technique to make source materials evaporate and then condense to finally obtain the materials on nanoscale The experiments are conducted with rigid reaction parameters of temperature, pressure, carrier gas, substrate, and evaporation time, etc., whose settings will be ultimately crucial to the morphology and phase structure of as-prepared nanostructures The process is employed successfully by Wang et al to manufacture novel one-dimensional nanostructures called nanobelts The unique geometrical configuration of nanobelts with a rectangle-like cross-section of 30 to 300 nanometers width and to 10 width-to-thickness ratios, and lengths up to several millimeters, may potentially open up new application fields of nanostructures The achievement of ZnO nanobelts thus further pushes the research developments forward in the synthesis of novel one-dimensional (1D) nanostructured candidates for fully understanding dimensional confinement effects and constructing nanodevices Up to now, ZnO nanobelts tentatively used as nanocantilever, field-effect transistors have been demonstrated [3] 2.3 Template-Based Methods Roughly the templates used for synthesis of nanostructures of ZnO can be divided into hard and soft features Most common hard templates are based on anodic alumina membranes (AAM), which have hexagonally ordered porous structure (channel diameter below 10 to 200 nm, lengths from to over 100 m, and channel density in the range of 1010–1012 cm−2 ) [39] Since the nanochannel of AAM can be adjusted by finely varying preparation parameters and AAM itself possesses good chemical and thermal stabilities, it can be used as an ideal template to deposit desired materials on nanoscale based on the porous shape of AAM Coupled with other techniques such as electrophoretic deposition [40] and electrochemical deposition [41], this method has been employed widely to build nanowires or nanoarrays For this approach, the method of source material injection and porous structure of AAM are of vital importance for material size and morphology modulation Li et al [41c] reported the successful preparation of ZnO nanowire arrays 769 Zinc Oxide Nanostructures embedded in AAM by oxidizing Zn nanowire arrays electrodeposited in an AAM nanochannel The nanowire arrays possessing polycrystalline feature with diameters from 15 to 90 nm exhibit blue emission By using an anodic aluminum oxide (AAO) templated method coupled with the sol–gel process, uniform polycrystalline ZnO nanotubules were assembled in the membrane pores of AAO Soft templates generally refer to those chemically formed structures without rigidity such as mesoporous silica [42], polymeric filter membranes [43], poly(styrene) beads [44], and copolymer [45] These structures are normally on nanoscale with well-controlled shapes, which can be used as templates to morphologically controllable synthesis Zhang et al have made the attempt to prepare ZnO nanoclusters in the channels of MCM-41 by high temperature calcinations of MCM41 with zinc cations attachment Zheng et al employed an AAO membrane as the template for preparation of an array of zinc oxide (ZnO) and polyaniline coaxial nanowires, which shows enhanced photoluminescence Neves and coworkers obtained the hollow structure of ZnO by the calcination of hydrozincite coated poly(styrene) beads All soft templates afford versatile ways for achieving nanostructures with diverse morphologies 2.4 Solution-Based Approaches 2.4.1 Overview A solution-based approach has been attractive for many years for the achievement of ZnO nanostructures In comparison with other methods, this approach possesses the advantages of relatively low cost, mild reaction conditions, inexpensive equipment, and easy control by simply adjusting the chemical process of reactions Up to now, a solutionbased approach might be one of the lowest cost ways but popular and powerful in the synthesis field of nanostructures Recently much research on the synthesis of ZnO with novel and unusual morphologies via this synthetic method has been demonstrated Boyle et al [46] have grown perpendicularly oriented ZnO rods along the c-axis on thin ZnO template from aqueous solution of zinc acetate and hexamethylenetetraamine The ZnO nanorods deposited on substrate exhibit more homogenous and ordered features than the previous reports in which ZnO was obtained by aqueous thermal decomposition of Zn2+ amino complex [47] Vayssieres et al showed a successful preparation of three-dimensional arrays of highly oriented crystalline ZnO microtubes by employing aqueous solution of Zn2+ ions and metheneamine as starting materials [48] Recently, there have been several continual reports [28] on the preparation of nanodots, regular nanorods, and columns of ZnO by solution-based methods with assistance of additives or microemulsions, indicating the versatile utilizations of the solution-based approach for morphology control Our preparation is based on the facts that previous reports on preparation of ZnO by a wet chemical route such as hydrothermal, sol–gel, and precipitation processes are mainly dependent on the decomposition of hydroxide or products of hydrolyzed zinc salt in pure water or a simple solvent system, which cannot led to diversiform morphologies with uniform size distribution Therefore, we like the ideas of using a single precursor solution, a one-step process, and a low temperature condition with microemulsion modulation to realize morphology control of nanostructured and ZnO We employ simple Zn2+ complexes, Zn(OH)2− 2+ Zn(NH3 )2+ in base , obtained at different pH values of Zn solution dependent on the fact that Zn2+ is liable to form a soluble complex when the base is excessive, as precursors, since previous investigations on nanostructured ZnO formation suggest that Zn(OH)2− is proposed to be the growth unit and can be directly incorporated into ZnO crystallites under given conditions [49] The processes for precursor decomposition into ZnO nanostructures are simple in terms of chemical reaction However, the advantages of this approach may be seen in several aspects: microemulsions and solvents as a driving force to modulate crystal nucleation and growth process, lower hydrothermal temperature to ensure sufficient crystallinity, and single source precursors to integrate into nanostructures with the desired building method By adjusting parameters of solvent, precursor, solution basicity, reaction temperature, as well as time, the preparative process becomes controllable to get desired high quality ZnO with controllable morphologies In our preparation, Zn(OH)2− precursor solution is prepared by mixing 0.5 mol/L ZnAc2 and mol/L NaOH solutions (volume ratio, v/v = 1:1, pH ≈ 14), while Zn(NH3 )2+ precursor is achieved by mixing 0.5 mol/L ZnAc2 and or fresh ammonia (pH ≈ 10 or 12) Typically, Zn(OH)2− Zn(NH3 )2+ precursor solution, surfactant cetyltrimethylam4 monium bromide (CTAB) (absence in some cases), cosurfactant n-hexanol (absence in some cases), and solvent were mixed with various ratio (see Table 1) to form a solution or microemulsion in a vessel under constant stirring The mixture is then transferred into a 25 mL Teflon-lined autoclave for hydrothermal treatment to a given temperature for a certain time (see Table 1) After the reaction finished, the autoclave containing the samples is cooled to room temperature naturally The white precipitate deposited on the bottom of the autoclave is collected and washed with absolute ethanol and distilled water several times Finally, ZnO samples are obtained by centrifugation and dehydration of the precipitate in vacuum at 60–70 C All reaction parameters and morphologies are summarized in Table 2.4.2 One-Dimensional Nanostructures Nanorods By thermally treating Zn(OH)2− precursor solution in a CTAB quaternary microemulsion-based system at various reaction conditions, ZnO nanorods with variable diameter and length can be produced [50] Figure shows a typical transmission electron microscopy (TEM) image of ZnO nanorods obtained by hydrothermally treating a microemulsion system consisting of g CTAB, 1.2 mL Zn(OH)2− solution, 5.0 mL n-hexanol, and 8.2 mL n-heptane at 180 C for 13 h From the low magnification TEM image shown in Figure 1a, it can be seen clearly that ZnO nanorods possessing a good crystallinity are formed and most of them are relatively uniform, with an average diameter less than 100 nm and length up to tens of micrometers The high resolution TEM (HRTEM) shown in Figure 1b proves the wurtzite phased structure of ZnO with a deduced lattice spacing of 2.59 Å corresponding 770 Zinc Oxide Nanostructures Table Summarized morphologies and reaction conditions Reaction media pHa Temperature ( C) Time (h) g CTAB/1.2 mL Zn(OH)2− solution/ 5.0 mL n-hexanol/8.2 mL n-heptane g CTAB/1.2 mL Zn(OH)2− solution/ 2.0 mL n-hexanol/11.2 mL n-heptane g CTAB/1.2 mL Zn(OH)2− solution/ 2.0 mL n-butanol/11.2 mL n-heptane g CTAB/1.2 mL Zn(OH)2− solution/ 3.0 mL n-hexanol/10.2 mL n-heptane 1.2 mL Zn(NH3 )2+ solution/ 13.2 mL ethanol g CTAB/1.2 mL Zn(NH3 )2+ solution/ 3.0 mL n-hexanol/10.2 mL n-heptane 1.2 mL Zn(OH)2− solution/ 13.2 mL H2 O 1.2 mL Zn(OH)2− solution/ 13.2 mL n-heptane 1.2 mL Zn(NH3 )2+ solution/ 13.2 mL ethanol 1.2 mL Zn(OH)2− solution/ 13.2 mL ethanol 1.2 mL Zn(OH)2− solution/ 13.2 mL ethanol 1.2 mL Zn(OH)2− solution/ 13.2 mL H2 O 1.2 mL Zn(OH)2− solution/ 13.2 mL H2 O 14 180 13 14 180 13 14 180 13 14 140 13 12 180 13 10 100 13 14 180 13 14 180 13 10 100 13 14 100 13 14 180 13 14 180 10 14 180 Morphology Rodlike (Fig 1) Rodlike (Fig 3a) Rodlike (Fig 3b) Wireslike Tubelike Tubelike Flowerlike Snowflakelike Prismlike Prickly sphere-like Regular rodlike Columnlike Multiarm architecture 2+ The pH of Zn(OH)2− or Zn(NH3 )4 precursor solutions 500 nm Figure TEM (a) and HRTEM (b) images (the inset is ED pattern) of ZnO nanorods fabricated via a solution-based route (for detailed preparative information, see Table 1) The lattice spacing of 2.59 Å corresponds to (002) plane, implying the preferred growth orientation of ZnO nanorods along (002) plane 20 30 40 60 70 (202) (103) (110) 50 (200) (112) (201) a (102) b b (100) a ratio of n-heptane and hexanol from 8.2/5.0 to 11.2/2.0, but maintaining other compositions of microemulsion and reaction conditions, the diameters of ZnO nanorods achieved are gradually increased Figure 3a shows the ZnO structures obtained at the ratio of 11.2/2.0, which exhibits rodlike morphology as well, but more uniform and with an average diameter (around 350 nm) larger than that of the (002) (101) to the (002) plane, implying the preferred growth orientation of ZnO nanorods along the (002) plane The electron diffraction (ED) pattern in the inset of Figure 1b indicates the single crystalline nature of ZnO nanorods The X-ray diffraction (XRD) pattern in Figure 2a also reveals the wurtzite structure of ZnO (hexagonal phase, space group P 63mc), which is consistent with the ED characterization shown in the inset of Figure 1b All diffraction peaks can be well assigned to hexagonal-phase ZnO (JCPDS card, No 36-1451) Compared with the standard diffraction pattern, as shown in Figure (the bottom), the diffraction intensity differences along [101] and [002] directions may imply the preferred growth orientation of ZnO By increasing the Relative Intensity / a.u a 80 2Theta / ° Figure XRD patterns of ZnO nanorods fabricated via a solutionbased route (a) Sample shown in Figure and (b) sample shown in Figure The bottom is a standard XRD pattern of ZnO (JCPDS No.36-1451) 771 Zinc Oxide Nanostructures a b µm a µm Figure TEM images of ZnO nanorods fabricated via a solutionbased route in different microemulsion systems of (a) g CTAB, 1.2 mL Zn(OH)2− solution, 2.0 mL n-hexanol, and 11.2 mL n-heptane, and (b) the same microemulsion compositions as (a) except replacing n-hexanol by n-butanol previous sample shown in Figure 1, indicating microemulsion compositions are of significant importance in modulating the morphology during the formation process of ZnO nanostructures The ED pattern in the inset of Figure 3a displays the single crystalline nature of ZnO with the wurtzite phased structure consistent with the XRD result shown in Figure 2b When replacing n-hexanol by butanol and maintaining other reaction conditions, the diameters of obtained ZnO nanorods (1 m) become much larger than that of comparative samples, as shown in Figure 3b, confirming the modulation function of microemulsion compositions on ZnO morphology formation To fully understand the formation process of ZnO nanorods, the time-dependent experiments are carried out to monitor the evolving process by recording the shapes of the sample achieved at stepwise increased heating time from 10 to h The series scanning electron microscopy (SEM) images in Figure show the morphology evolution with prolonged reaction time Figure 4a shows that when operating the reaction for 10 min, the aggregates of spherical morphology ZnO with an average diameter as smaller as 30 nm, other than nanorods, are observed This implies that at this stage nanostructured ZnO forms in a fashion of spherical shape in microemulsion droplets and the interactions such as collision, conglutination, and amalgamation among the microemulsion droplets would lead to the aggregates of nanoparticles Figure 4b indicates the morphology exhibition of ZnO obtained after 30 heating treatment At this stage, besides the aggregated nanoparticles, ZnO nanorods are observed, which indicates that parts of nanoparticles observed in Figure 4a have evolved into nanorods Although the length and diameter of the nanorods are not uniform, it can be deduced from the image that the diameters of most nanorods not exceed 100 nm The reason for this observation is that with the extension of heating time, the nucleation of ZnO may gradually be completed and the directed aggregation growth mediated by microemulsion droplets might occur By aggregation, the microemulsion droplets might connect into a linearlike shape, and then by experiencing an Ostwald ripen process, the linearlike aggregates recrystallize into perfect single-crystalline 1D structure along the preferred orientation For a synthetic strategy by employing microemulsion as reaction media, it is proved that the shape of a microemulsion droplet has important effects on morphology formation of nanostructures [51] With the microemulsion composition b 500 nm c 500 nm d µm µm e f µm µm Figure Series of SEM images of shape evolution of ZnO nanorods shown in Figure 3a with stepwise prolonged reaction time (a) 10 min, (b) 30 min, (c) h, (d) h, (e) h, and (f) h, exhibiting the evolution of ZnO nanoparticles to nanorods used in our preparation, the microemulsion droplets are liable to form spherical shape at room temperature and their sizes can be adjusted by changing the ratio of solvent to cosurfactant (n-hexanol to n-heptane in present case) or altering cosurfactant species [51] With this regard, the diameter of ZnO nanorods can be adjusted by changing the size of microemulsion droplet The proposed formation process is further proved by the morphological changes during continually extending reaction time to 1, 2, 4, and h The representative SEM images corresponding to each stage are shown in Figure 4c–f Figure 4c shows uniform rod shape with an average diameter of about 200 nm after driving the reaction for h Figure 4d–f displays the Ostwald ripen processes of rodlike ZnO, which gradually makes the crystallinity improved and the diameter grow With the Ostwald ripening, the ends of nanorods show the evolved tendency from pyramid- to prismlike shapes, which can be seen clearly from the comparison of images shown in Figures 4e and 3a This accords with the crystal growth habit of ZnO with different growth velocities along different crystal planes [49, 52] In addition, from images shown in Figure 4b–f, it can be seen that some nanorods are liable to form a homocentric bundle, which is more evident in Figure 4e This phenomenon has also been observed in other shapes of ZnO, implying the detailed growth fashion [33a] Nanowires In consideration of the strategies we used to adjust the diameters of nanorods, it is reasonably believed that by varying the reaction conditions, the nanowires can 772 (101) (202) (112) 70 80 (pH ∼ 12) and ethanol (v/v = 1/11) at 180 C for 13 h Following the typical preparation procedures, the tubular ZnO is obtained XRD structural characterization (Fig 7) indicates the wurtzite phase (hexagonal, space group P 63mc) with good crystallinity Figure displays the SEM and TEM images showing different aspects of the tubular structure of ZnO The image in Figure 8a shows a few tubular ZnO coexisted with some small nanoparticles, while Figure 8b shows the magnified hollow structure of the individual tube marked by the arrow From these two images it can be deduced that the outer and inner diameters of the hollow tube are about ∼450 and ∼250 nm, respectively The average length of the tube is about ∼4 m And the tube wall with thickness of ∼100 nm is not smooth as it is built up by small polycrystalline nanoparticles Figure 8c gives the TEM image of an individual tube with the destroyed ends, showing clearly the a 20 30 40 50 (103) 60 70 (202) (102) 300 nm (200) (112) (201) b c 500 nm (201) (103) 60 Figure X-ray diffraction pattern of ZnO nanowires shown in Figure (002) (101) 10 µm 50 (200) (110) (102) (002) 40 2θ (degree) Intensity b 30 (100) a 20 (110) Nanotubes Apart from our works [34a], recently there have been several other reports concerning tubular ZnO nanostructures [16f, 22d, 39d, 43, 54] The methods for preparing tubular ZnO are mostly based on the processes of pyrolysis, thermal reduction routes, or AAO based template methods A low cost, solution-based approach for growing tubular ZnO still remains a big challenge Our preparation is conducted with the mixture of Zn(NH3 )2+ aqueous solution Intensity (Counts) be achieved Based on the nanorod morphology, to get the nanowires, the Ostwald ripening process should be slowed down by decreasing the reaction temperature and the nuclei should be smaller as well Then by aggregating smaller nanoparticles into 1D nanostructures, the nanowires would possibly be achieved The nanowire preparation is carried out by thermally treating a microemulsion consisting of g solution, mL n-hexanol, and CTAB, 1.2 mL Zn(OH)2− 10.2 mL n-heptane at 180 C for 13 h [34b] The representative image shown in Figure 5a reveals that the obtained sample exhibits nanowire structure with a diameter ranging from 30 to 150 nm The aspect ratio of nanowires is estimated to be larger than 50 nm, as a typical image in Figure 5c shows a single nanowire with aspect ratio more than 200 (diameter of 50 and length up to approximately 14 m) The TEM image (Fig 5b) shows another shorter single ZnO nanowire with a diameter of ∼30 nm and a length up to m The ED pattern (inset of Fig 5b) indicates that the ZnO nanowires have a single crystal nature with preferred growth orientation along the (110) crystal plane based on the calculation of the diffraction dots The blurry diffraction dots in the inset image might hint at the existence of branch crystallites around the nanowires The XRD pattern of ZnO nanowires (Fig 6) shows the same fashion as that of ZnO nanorods (Fig 2) However, the obtained result of ZnO nanowire growth along (110) crystal plane deduced from the diffraction calculation is different from that of nanorods achieved from HRTEM before, which adopts oriented growth along (002) crystal plane, indicating the possibility to obtain 1D ZnO nanostructures with different orientated growth direction by the same approach The same results of SnO2 nanorods achieved via a solution-based route may be another good example to prove this [53] It is believed that the nanowire formation mechanism is analogous to that of ZnO nanorods (100) Zinc Oxide Nanostructures 80 2Theta (deg) Figure SEM (a) and TEM (b and c) images of ZnO nanowires synthesized via a solution-based route The inset of in (b) is an ED pattern of a single nanowire Figure X-ray diffraction pattern of polycrystalline tubular ZnO fabricated via a solution-based route 773 Zinc Oxide Nanostructures a a b µm µm 500 nm Figure SEM images (a and b) of monocrystalline tubular ZnO fabricated via a solution-based route (for detailed preparative information, see Table 1) (b) The magnified image of the white pane part in (a), showing the hollow structure b 100 nm c Figure 9, it can be deduced that the individual microtube of ZnO with a wall thickness of ∼200 nm may be approximately 1–2 m wide and 8–10 m long It is proposed that during the formation process, the decomposition of Zn(NH3 )2+ precursor may result in the generation of NH3 which may be responsible for the formation of hollow structures 2.4.3 Other Unusual Morphologies 500 nm d –500 Figure TEM (a, b and c) and SEM (d) images of polycrystalline tubular ZnO fabricated via a solution-based route hollow structure, and the SEM image in Figure 8d further affords a spatial view of the destroyed end of the tube The tubular sample is obtained by hydrothermal treatprecursor solution in ethanol to lead ment of Zn(NH3 )2+ to the polycrystalline tubular ZnO Interestingly, when changing the reaction media into microemulsion with solution, compositions of g CTAB, 1.2 mL Zn(NH3 )2+ 3.0 mL n-hexanol, and 10.2 mL n-heptane, and lowering the pH value of the precursor solution to around 10 and temperature to 100 C, hexagonal structure ZnO microtubes with single crystalline nature are obtained The SEM image in Figure 9a shows finely crystallized hexagonal structures, and some of them grow separately and others have the growing feature of embedded into each other From the destroyed parts of the hexagonal structures, the hollow structure can be seen clearly The embedded feature of hollow structures may be attributed to the homocentrically growing tendency of ZnO crystals under our reaction conditions, which is observed for the homocentrically growing branch of nanorods Figure 9b shows the magnified image of hollow structures of one destroyed end of a ZnO microtube marked by a white pane in Figure 9a, displaying the evident hollow structure of ZnO single crystal microtubes From Flower- and Snowflake-Like Nanostructures Based on the successful preparation of 1D ZnO nanostructures via a solution-based route with the assistance of microemulsion modulation, it is reasonably believed that reaction media play an important role in modulating the morphology during the reaction process Therefore, some pure solvents other than microemulsion were also employed to comparatively investigate the possibilities of morphologically controllable synthesis by varying reaction media By employing 2+ the mixture of Zn(OH)2− or Zn(NH3 )4 precursor solution and solvent according to the volume ratio of 1:11, some unusual nano- or microstructures are achieved by following the same reaction procedures The summarized morphologies and reaction conditions are listed in Table By decomposing Zn(OH)2− precursor in H2 O solvent at 180 C for 13 h, flowerlike morphology of ZnO is achieved, as the representative SEM image shows in Figure 10a The flowerlike ZnO with average size about ∼1.5 m displays a feature of homocentric growth due to multiple nuclei twinning at the onset of growth, which finally becomes an individual crystalline nucleus to make all branch growth at different directions The flowerlike structure of ZnO is further confirmed by the TEM image shown in Figure 10b The ED pattern inset in Figure 10b indicates the single crystalline nature and wurtzite structure of ZnO, which is consistent with the results of XRD characterization shown in Figure 11b From the ED pattern, it is also deduced that the flowerlike shape of ZnO is not the simple aggregation of small crystallites but is formed by monocrystallines of nanorods growing homocentrically The similar morphology of ZnO has been previously observed by Chittofrati et al [33b] and appeared in a very recent publication [55] Figure 10c and d gives the magnified SEM images of two individual flowerlike ones marked by a white pane in Figure 10a, showing the detailed features of two different growing patterns of ZnO flowers When changing the polar solvent of H2 O by nonpolar solvent of n-heptane, snowflakelike structures of ZnO are observed, as 774 Zinc Oxide Nanostructures a b f f Intensity e d c 20 30 40 50 70 (202) (103) 60 (200) (112) (201) (110) a (102) d (100) c b 500 nm (002) (101) µm 80 2Theta (deg) Figure 11 X-ray diffraction pattern of ZnO nanostructures with various morphologies (a) Standard ZnO (JCPDS #36-1451), (b) flower-, (c) snowflake-, (d) prism-, (e) prickly sphere-, and (f) regular rodlike ZnO e f a µm b 500 nm Figure 10 SEM (a, c, d, and e) and TEM (b and f) images of ZnO fabricated via a solution-based route (a) and (b) General morphology of flowerlike ZnO The inset of (b) is the ED pattern (c) and (d) Magnified images of two individual ZnO flowers marked with the white pane in (a) (e) and (f) Snowflakelike ZnO (for detailed preparative information, see Table 1) µm c shown in Figure 10e and f The size of snowflakelike morphology is in the range of 0.5–1.0 mm, having the same growth fashion but composed of needle ZnO nanocrystallites XRD in Figure 11c shows the wurtzite phase structure of ZnO Prism-, Prickly Sphere-, and Regular Rodlike Nanostructures When thermally treating the reaction media containing Zn(NH3 )2+ precursor at 100 C for 13 h, prismlike morphology of ZnO is obtained, with average diameter ∼500 nm and length ∼1 m, as can be seen clearly from the SEM image in Figure 12a for the aspect ratio, and the TEM image in Figure 12b for the top view of the regular hexagonal shape of the prism When operating the reaction in ethanol, at 100 C for 13 h by decomposing Zn(OH)2− prickly sphere-like instead of prismlike structures of ZnO are achieved From the SEM image of Figure 12c, it can be deduced that the average diameter of prickly sphere-like ZnO is approximately 1.5 m, while from the TEM image in Figure 12d, it is concluded that the surface of the prickly sphere-like ZnO is built up by needlelike rods several tens of nanometers in width With identical reaction conditions 500 nm d µm e 500 nm f 100 nm 100 nm Figure 12 SEM (a, c and e) and TEM (b, d and f) images of ZnO fabricated via a solution-based route (a) and (b) Prismlike morphology (c) and (d) Prickly sphere-like morphology (e) and (f) Regular rodlike morphology of ZnO (for detailed preparative information, see Table 1) The inset in (f) is an ED image 775 Zinc Oxide Nanostructures for preparing a prickly sphere-like ZnO sample but increasing temperature to 180 C, regular rodlike nanostructures of ZnO are formed, as SEM amd TEM images show in Figure 12e and f The regular nanorods of ZnO have unique aspect ratio near 5, with an average diameter of 100 nm and length of 500 nm The ED pattern (the inset of Fig 11f) shows the orientated growth of ZnO nanorods along the c-axis of hexagonal phase structured ZnO XRD patterns in Figure 11d, e, and f show the wurtzite phase structures of prism-, prickly sphere-, and regular rodlike ZnO Columnlike Morphology and Multiarmed Architectures While keeping all the preparative conditions for synthesizing flowerlike structures of ZnO, we conduct the experiments for 10 at a higher basicity by adding one time more volume of NaOH solutions (5 M) to ZnAc2 solution when precursor solution ZnO obtained at preparing Zn(OH)2− this condition exhibits the aggregated column feature composed of flakelike branches The size of the column is up to several micrometers, as the SEM image shows in Figure 13a When prolonging the reaction to h at the same conditions, the columnlike ZnO totally evolves into uniform multiarmed architectures formed by branch rods sharing one growth center The uniform ZnO multiarmed architecture seems to be formed by nanorods growing homocentrically, which is similar to the formation of flowerlike structure of ZnO (Fig 10a), but the diameter of the rod (∼300 nm) deduced from the SEM image in Figure 13b is smaller and a µm b µm Figure 13 SEM images of (a) columnlike morphology and (b) multiarmed architecture of ZnO the aspect ratio is bigger It is believed that the growth process of the multiarmed architecture is analogous to that of flowerlike shape but having a different crystal growth rate along a certain crystal face due to the effects of various solution basicities on crystal growth habit 2.4.4 Formation Mechanism In fact, for the solution-based route employed in our preparations, it can be expediently classified into two types of microemulsion- and solvent-based ways in terms of the reaction media in order to present the formation mechanism To understand the formation mechanism of 1D nanostructures of ZnO achieved in microemulsion systems, the roles of microemulsion should be taken into consideration In principle, the crystallization of ZnO shows a tendency toward 1D nanostructures under hydrothermal conditions [56], but without the assistance of additives Most times irregular ellipsoidal shapes resulted and no regular nanorods can be obtained To avoid these disadvantages, we employ the reaction media of microemulsions in consideration of modulating crystal formation It is proposed that the formation of homogeneous ZnO nanorods and nanowires might be induced and achieved via a directed aggregation growth process mediated by the microemulsion droplets, as suggested in previous works [57] Following this mechanism, the microemulsion may function both on the nucleation and growth stages of the formation of 1D ZnO nanostructures On the nucleation stage, on the one hand, microemulsion droplets can act as microreactors and play a role in controlling nucleation rate and nuclei size, which would eventually affect the final morphologies of the samples On the other hand, the characteristic kinetics processes of collision, conglutination, and amalgamation among microemulsion droplets may, as a result, result in the linear aggregations of ZnO nanoparticles While in the growth stage, the linear aggregations will grow into the well-crystallized 1D nanostructures by experiencing a recrystallization of the Oswald ripen process confined by miroemulsion droplets Therefore, changing the microemulsion compositions to make droplet smaller, which means the nuclei would be smaller as well, may induce nanowires rather than nanorods The formation mechanism for 1D nanostructures along the preferred growth orientation is analogous to another investigation [58] on the transformation process of nanodots to nanorods by experiencing an “oriented attachment” mechanism, but with a confinement of microemulsion droplets Besides, it is also suggested that surfactant CTAB is beneficial to the transport and orderly stacking of the crystal growth units [16k] The growth processes of ZnO samples under solvothermal reaction conditions are described as follows At the initial heating stage, the temperature is comparatively low, and many individual nuclei are formed in the reaction media During this stage, the formed nuclei are individually dispersed in the reaction media Within the ramping range of the temperature until the solvent is boiling, the individual nuclei gradually grow and begin to amalgamate in an airtight autoclave Since the boiled solvent exhibits the dropletlike dispersion in the autoclave under gas–liquid equilibrium, there are many nuclei in one droplet Thus, all nuclei in one droplet have the tendency to amalgamate into a big 776 Zinc Oxide Nanostructures one As the amalgamation of the nuclei occurred, the crystal growth also took place on various crystal faces with different growth rates As a result, the dendrite of ZnO grew into rodlike form from one big center nucleus by the amalgamating of a large amount of nuclei in one droplet Therefore, the morphological features of ZnO samples are liable to be dendriticlike, such as flowers, snow flakes, and prickly spheres Of course, due to influences of variation of the precursors, solvents, and reaction condition, the growth habit of ZnO changes accordingly, which causes the variation of final morphologies of ZnO The influences of solvents, precursors, reaction time, and temperature as well as basicity may have influence on the ZnO morphologies as discussed in detail in our previous publication [33a] For the solvent, it is found that the polarity and saturated vapor pressure of solvents may affect the resulting morphology by adjusting the homogenization of the reactants in reaction medium, the amount of individual nucleus formation, the amalgamation, and the direction preference of growing nucleus For the precursors, we employ two kinds of single source precur2+ sor of Zn(OH)2− and Zn(NH3 )4 solutions The precursors may follow different reaction equations to form ZnO nanostructures shown as follows: Zn NH3 2+ a µm b µm c + 2OH− −→ ZnO + 4NH3 + H2 O Zn OH 2− −→ ZnO + H2 O + 2OH− The first reaction may produce a by-product of NH3 , which may play a role in preventing the amalgamation of nucleus in the supersaturated solvents during reaction process Therefore, it is easy to lead to dispersed morphology of prismlike or hollow tubular structures by using Zn(NH3 )2+ as precursors Time and temperature are the key factors responsible for morphology evolutions Series images in Figure 14 record the process of morphology evolution of flowerlike ZnO from the morphology of sheetlike branch aggregates at 10 (Fig 14a), to rodlike branch aggregates at and h (Fig 14b, c), and finally to a well-developed flowerlike shape (Fig 4d) [33a] For other morphologies of ZnO, it has the same evolution feature under solvothermal conditions Figure 15a–d shows the initial shape of snowflake-, prickly sphere-, rod-, and prismlike ZnO, respectively, indicating the discrepancies of initial and final morphologies due to the crystal growth habits under different reaction conditions OPTICAL PROPERTIES OF ZnO NANOSTRUCTURES As it is mentioned that ZnO is a wide bandgap semiconductor with a large exciton binding energy at room temperature, so it is a suitable candidate in forms of disordered particle, thin film, and nanowire arrays for the realization of ultraviolet lasing action However, the photoluminescence of ZnO in the visible region of green instead of desired emission at UV range is often observed due to the high density of crystal defects, which may quench effective exciton emission by irradiative transition and cause a deeplevel or trap-state emission [13i] This deficiency hinders the progress of the applications of ZnO in optoelectronic and lasing devices In this regard, the development of suitable µm d µm Figure 14 Series SEM images of morphology evolution of flowerlike ZnO with the stepwise prolonged reaction time (a) 10 min, (b) h, (c) h, and (d) h synthetic approaches toward high quality ZnO is the basis of realizing UV emission and lasing Our morphologically controllable synthesis may afford a possible way to solve this problem The room temperature photoluminescence spectra of nanostructured ZnO with various morphologies are measured by using a Jobin Yvon–Labram spectrometer with a He–Cd laser focused to ca m as the excitation source ( ex = 325 nm) The photoluminescence of ZnO nanowires shown in Figure 16 exhibits a strong UV emission at 385 nm arising from the recombination of excitonic center of the nanowires [2i], and a green emission at 485 nm attributed to 777 Zinc Oxide Nanostructures a 1250 Intensity (a.u.) 1000 750 500 µm b 250 300 350 400 450 500 550 600 Wavelength (nm) Figure 16 Room temperature photoluminescence of ZnO nanowires µm 4000 c Intensity (a.u.) 3000 µm 2000 1000 d 325 350 375 400 425 450 λ (nm) Figure 17 Room temperature photoluminescence of polycrystalline tubular ZnO µm Flower Figure 15 SEM images of morphology exhibitions of (a) snowflake-, (b) prickly sphere-, (c) regular rod-, and (d) prismlike ZnO when driving reaction for 10 Regular rod Intensity / a.u the radiative recombination of a photogenerated hole with an electron occupying the oxygen vacancy [59] While the photoluminescence of polycrystalline tubular ZnO shown in Figure 17 displays a strong UV emission at 385 nm and no defect emission at the visible region is observed The discrepancies of photoluminescence between 1D structure of ZnO nanowires and tubes may be ascribed to the crystal quality in association with size and morphology, since the tubular ZnO is built up by sufficiently crystallized nanoparticles, which may have less possibility to generate singly ionized oxygen vacancy to lead to green emission [2d] The photoluminescence of other unusual morphologies is comparatively given in Figure 18 The photoluminescence has Snow flake Prickly sphere Prism 350 400 450 500 550 600 λ / nm Figure 18 Room temperature photoluminescence of ZnO with unusual morphologies 778 different features in UV and green emissions The prismlike ZnO shows a strong UV emission at ca 385 nm, due to its prismatic morphology with sufficient crystallization, which only results in the recombination of exciton The regular rodlike ZnO exhibits a strong UV emission with a very weak green emission (ca 510 nm) due to the recombination of electrons in singly occupied oxygen vacancies with photoexcited holes, as observed in ZnO nanowires The relative density of oxygen vacancies can be estimated by comparing the green emission intensity ZnO samples with three other kinds of morphologies (i.e., flower-, snowflake-, and prickly sphere-like) show both the UV and green emissions, but the relative intensity of UV emission gradually decreased with the morphology changing from flower-, snowflake-, to prickly sphere-like, which is a strong evidence indicating the increase of oxygen vacancies CONCLUSION This chapter reviews the solution-based route by making use of the advantages of microemulsion and hydrothermal preparative techniques to reach the morphology control aim of ZnO micro- and nanostructures in our laboratory The novel structures of ZnO including nanorods, nanowires, microtubes, and other unusual morphologies such as dendritic flower- and snowflake-, prickly sphere-, prism-, columnlike shape and multiarm architecture are successfully realized by adjusting the parameters of microemulsion compositions, solvent species, precursors, and other conditions such as time and temperature of solution synthesis The growth mechanism for each kind of morphology formation under microemulsion-based conditions is basically understood by a microemulsion droplet mediated aggregation process and other morphology achieved in pure solvent can be considered as results in terms of competition of nuclei diffusion and aggregation process The optical properties of various typical morphologies show a morphology-dependent features induced by the crystal quality due to morphology variations Such an abundant morphology world of microand nanostructured ZnO obtained via the current approach may afford, in consideration of fundamental research, better understanding of crystal nucleation and growth mechanism from various aspects and would provide, in consideration of technological applications, possible candidates for nanodevice constructions for various utilizations We believe this contribution would give some hints for fabrication and morphology control of other kinds of semiconductors GLOSSARY Hydrothermal and solvothermal The term hydrothermal usually refers to any heterogeneous reaction in the presence of aqueous solvents of mineralizers under high pressure and temperature conditions to dissolve and recrystallize (recover) materials that are relatively insoluble under ordinary conditions When the solvent is replaced by other matter, it is known as “solvothermal.” Microemulsion A microemulsion is a thermodynamically stable dispersion of one liquid phase into another, stabilized by an interfacial film of surfactant This dispersion may Zinc Oxide Nanostructures be either oil-in-water or water-in-oil Microemulsions are typically clear solutions, as the droplet diameter is approximately 100 nm or less The interfacial tension between the two phases is extremely low Ostwald ripening processes Ostwald ripening is the process by which larger particles (or, for emulsions, droplets) grow at the expense of smaller ones because of the higher solubility of the smaller particles (Gibbs–Thomson or Kelvin effect) and molecular diffusion through the continuous phase Template-based synthesis In this approach, the matter functions as the template simply serves as a scaffold within (or around) which a different material is generated in-situ and shaped into a nanostructure with its morphology complementary to that of the template Vapor-phase growth Vapor-phase growth is one kind of approach toward crystal growth in which the crystal is grown by depositing material directly from the vapor or gaseous state (i.e., the immediate precursor is in a vapor or gaseous state) Several methods can be assigned to this category, such as thermal evaporation, chemical vapor deposition, and metal organic chemical vapor deposition ACKNOWLEDGMENTS This work was supported by NSFC (20001002, 20023005, and 20221101), MOST (G19980613), MOE (the Foundation for University Key Teacher), and the Founder Foundation of PKU REFERENCES C M Lieber, Solid State Commun 107, 607 (1998); (b) J Hu, T W Odom, and C M Liber, Acc Chem Res 32, 435 (1999) (a) J Goldberger, R R He, Y F Zhang, S W Lee, H Q Yan, H J Choi, and P D Yang, Nature 422, 599 (2003) (b) H Q Yan, R R He, J Pham, and P D Yang, Adv Mater 15, 402 (2003) (c) Y Y Wu, H Q Yan, and P D Yang, Top Catal 19, 197 (2002) (d) P D Yang, H Q Yan, S Mao, R Russo, J Johnson, R Saykally, N Morris, J Pham, R R He, and H J Choi, Adv Funct Mater 12, 323 (2002) (e) J C Johnson, H Q Yan, R D Schaller, P B Petersen, P D Yang, and R J Saykally, Nano Lett 2, 279 (2002) (f) Y Y Wu, H Q Yan, M Huang, B Messer, J H Song, and P D Yang, Chem Eur J 8, 1261 (2002) (g) H Kind, H Q Yan, B Messer, M Law, and P D Yang, Adv Mater 14, 158 (2002) (h) M H Huang, S Mao, H Feick, H Q Yan, Y Y Wu, H Kind, E Weber, R Russo, and P D Yang, Science 292, 5523 (2001) (i) M H Huang, Y Y Wu, H Feick, N Tran, E Weber, and P D Yang, Adv Mater 13, 113 (2001) (a) W L Hughes and Z L Wang, Appl Phys Lett 82, 2886 (2003) (b) Z R Dai, Z W Pan, and Z L Wang, Adv Funct Mater 13, (2003) (c) M S Arnold, P Avouris, Z W Pan, and Z L Wang, J Phys Chem B 107, 659 (2003) (d) Y Berta, C Ma, and Z L Wang, Micron 33, 687 (2002) (e) X D Wang, P X Gao, J Li, C J Summers, and Z L Wang, Adv Mater 14, 1732 (2002) (f) P X Gao and Z L Wang J Phys Chem B 106, 12653 (2002) (g) K Mcguire, Z W Pan, Z L Wang, D Milkie, J Menendez, and A M Rao, J Nanosci Nanotech 2, 499 (2002) (h) Z W Pan, Z R Dai, and Z L Wang, Science 291, 1947 (2001) H Gleiter, Acta Mater 48, (2000) D C Reynolds, D C Look, B Jogai, J E Hoelscher, R E Sherriff, M T Harris, and M J Callahan, J Appl Phy 88, 2152 (2000) Zinc Oxide Nanostructures C R Lee, H W Lee, J S Song, W W Kim, and S Park , J Mater Synth Process 9, 281 (2001); L Vayssieres, k Keis A Hagfeldt, and S E Lindquist, Chem Mater 13, 4386 (2001); D S King and R M Nix, J Catal 160, 76 (1996); T Minami, Mater Res Soc Bull 25, 38 (2000); D M Bagnall, Y F Chen, Z Zhu, T Yao, S Koyama, M Y Shen, and T Goto, Appl Phys Lett 70, 2230 (1997); J Zhong, A H Kitai, P Mascher, and W Puff, J Eletrochem Soc 140, 3644 (1993) Y F Chen, D M Bagnall, H Koh, K Park, K Hiraga, Z Zhu, and T Yao, J Appl Phys 84, 3912 (1998) T Seiyama and A Kato, Anal Chem 34, 1502 (1962); L F Dong, Z L Cui, and Z K Zhang, Nanostruct Mater 8, 815 (1997); R J Shen, D Z Jia, K Liang, X Q Xin, and J Y Wang, Chin J Inorg Chem 16, 906 (2000) K Hara, T Horiguchi, T Kinoshita, K Sayama, H Sugihara, and H Arakawa, Sol Energy Mater Sol Cells 64, 115 (2000) 10 D S Ginley and C Bright, Mater Res Bull 25, 15 (2000); B G Lewis and D C Paine, Mater Res Bull 25, 22 (2000); A J Freeman, K R Poeppelmeier, T O Mason, R P H Chang, and T J Marks, Mater Res Bull 25, 45 (2000) 11 K Keis, L Vayssieres, S E Lindquist, and A Hagfeldt, Nanostruct Mater 12, 487 (1999) 12 R F Service, Science 276, 895 (1997) 13 (a) P Zu, Z K Tang, G K L Wong, M Kawasaki, A Ohtomo, H Koinuma, and Y Segawa, Solid State Commun 103, 459 (1997) (b) M Kawasaki, A Ohtomo, H Koinuma, Y Sakurai, Y Yoshida, Z K Tang, P Yu, G K L Wang, and Y Segawa, “Silicon Carbide III—Nitrides and Related Materials,” Pts and 2, Vol 264-2, p 1459 (1998) (c) A Ohtomo, M Kawasaki, Y Sakurai, Y Yoshida, H Koinuma, P Yu, Z K Tang, G K L Wong, and Y Segawa, Mater Sci Eng B 54, 24 (1998) (d) Z K Tang, G K L Wong, P Yu, M Kawasaki, A Ohtomo, H Koinuma, and Y Segawa, Appl Phys Lett 72, 3270 (1998) (e) M Kawasaki, A Ohtomo, I Ohkubo, H Koinuma, Z K Tang, P Yu, G K L Wong, B P Zhang, and Y Segawa, Mater Sci Eng B 56, 239 (1998) (f) V A Dijken, E A Meulenkamp, D Vanmaekelbergh, and A Meijerink, J Lumin 87, 454 (2000) (g) Y Sun, J B Ketterson, and G K L Wong, Appl Phys Lett 77, 2322 (2000) (h) A Ohtomo and M Kawasaki, IEICE Trans Electron E83C, 1614 (2000) (i) Y C Kong, D P Yu, B Zhang, W Fang, and S Q Feng, Appl Phys Lett 78, 407 (2001) (j) A Mitra and R K Thareja, J Appl Phys 89, 2025 (2001) (k) A Mitra, R K Thareja, V Ganesan, A Gupta, P K Sahoo, and V N Kulkarni, Appl Surf Sci 174, 232 (2001) 14 M Hiramatsu, K Imaeda, N Horio, and M Nawata, J Vac Sci Technol A 16, 669 (1998) 15 S C Liu and J J Wu J Mater Chem 12, 3125 (2002); J J Wu, and S C Liu, J Phys Chem B 106, 9546 (2002); J Zhang, W Y Yu, and L D Zhang, Phys Lett 299, 276 (2002); J J Wu and S C Liu, Adv Mater 14, 215 (2002) 16 (a) H T Ng, B Chen, J Li, J E Han, M Meyyappan, J Wu, S X Li, and E E Haller, Appl Phys Lett 82, 2023 (2003) (b) H Kim and W Sigmund, Appl Phys Lett 81, 2085 (2002) (c) C K Xu, G D Xu, Y K Liu, and G H Wang, Solid State Commun 122, 175 (2002) (d) Y Dai, Y Zhang, Q K Li, and C W Nan, Chem Phys Lett 358, 83 (2002) (e) R Wu, C S Xie, H Xia, J H Hu, and A H Wang, J Cryst Growth 217, 274 (2000) (f) J Q Hu, Q Li, X M Meng, C S Lee, and S T Lee, Chem Mater 15, 305 (2003) (g) J S Lee, M I Kang, S Kim, M S Lee, and Y K Lee, J Cryst Growth 249, 201 (2003) (h) R Wu, J Wu, C S Xie, J Zhang, and A H Wang, Mater Sci Eng A 328, 196 (2002) (i) B D Yao, Y F Chan, and N Wang, Appl Phys Lett 81, 757 (2002) (j) H J Yuan, S S Xie, D F Liu, X Q Yan, Z P Zhou, L J Ci, J X Wang, Y Gao, L Song, L F Liu, W Y Zhou, and G Wang, Chem Phys Lett 371, 337 (2003) (k) X M Sun, Z X Deng, and Y D Li, Mater Chem Phys 80, 366 (2003) (l) Y K Tseng, I N Lin, K S Liu, T S Lin, and I C Chen, J Mater Res 18, 714 (2003) 779 17 (a) W I Park, D H Kim, S W Jung, and G C Yi, Appl Phys Lett 80, 4232 (2002) (b) Y J Zhang, N L Wang, S P Gao, R R He, S Miao, J Liu, J Zhu, and X Zhang, Chem Mater 14, 3564 (2002) (c) S C Lyu, Y Zhang, H Ruh, H J Lee, H W Shim, E K Suh, and C J Lee, Chem Phys Lett 363, 134 (2002) (d) W I Park, Y H Jun, S W Jung, and G C Yi, Appl Phys Lett 82, 964 (2003) (e) K Ogata, K Maejima, S Fujita, and S Fujita, J Cryst Growth 248, 25 (2003) (f) S Y Li, C Y Lee, and T Y Tseng, J Cryst Growth 247, 357 (2003) (g) C J Lee, T J Lee, S C Lyu, Y Zhang, H Ruh, and H J Lee, Appl Phys Lett 81, 3648 (2002) (h) W Il Park, G C Yi, M Y Kim, and S J Pennycook, Adv Mater 14, 1841 (2002) (i) K S Kim and H W Kim, Physica B 328, 368 (2003) 18 N Audebrand, J P Auffredic, and D Louer, Chem Mater 10, 2450 (1998); P D Cozzoli, M L Curri, A Agostiano, G Leo, and M Lomascolo, J Phys Chem B 107, 4756 (2003) 19 L F Dong, Z L Cui, and Z K Zhang, Nanostruct Mater 8, 815 (1997) 20 P Ayyub, R Chandra, P Taneja, A K Sharma, and R Pinto, Appl Phys A 73, 67 (2001) 21 Y W Heo, V Varadarajan, M Kaufman, K Kim, D P Norton, F Ren, and P H Fleming, Appl Phys Lett 81, 3046 (2002) 22 (a) O Milosevic, V Gagic, J Vodnik, A Mitrovic, L Karanovic, B Stojanovic, and L Zivkovic, Thin Solid Films 296, 44 (1997) (b) X Y Zhao, S C Zhang, C Z Li, B C Zheng, and H C Gu, J Mater Sci Procs 5, 227 (1997) (c) S Sakohara, M Ishida, and M A Anderson, J Phys Chem B 102, 10169 (1998) (d) J J Wu, S C Liu, C T Wu, K H Chen, and L C Chen, Appl Phys Lett 81, 1312 (2002) (e) T Tani, L Madler, and S E Pratsinis, J Nanopart Res 4, 337 (2002) (f) Y H Lin, Z L Tang, Z T Zhang, F L Yuan, Y B Ling, J L Lee, and S L Huang, J Am Ceram Soc 83, 2869 (2000) 23 R Liu, A A Vertegel, E W Bohannan, T A Sorenson, and J A Switzer, Chem Mater 13, 508 (2001); L Dloczik, R Engelhardt, K Ernst, S Fiechter, I Sieber, and R Konenkamp, Appl Phys Lett 78, 3687 (2001) 24 S H Yu, J Yang, Y T Qian, and M Yoshimura, Chem Phys Lett 361, 362 (2002); W J Li, E W Shi, Y Q Zheng, and Z W Yin, J Mater Sci Lett 20, 1381 (2001); W J Li, E W Shi, M Y Tian, and W Z Zhong, Sci Chin Ser E 41, 449 (1998); W J Li, E W Shi, M Y Tian, W Z Zhong, and Z W Yin, J Mater Res 14, 1532 (1999); X M Sun, X Chen, Z X Deng, and Y D Li, Mater Chem Phys 78, 99 (2003) 25 R J Shen, D Z Jia, Y M Qiao, and J Y Wang, J Inorg Mater 16, 625 (2001); Y L Du, M S Zhang, Y Deng, Q Chen, Z Yin, and W C Chen, Chin Phys Lett 19, 372 (2002); R J Shen, D Z Jia, K Liang, X Q Xin, and J Y Wang, Chin J Inorg Chem 16, 906 (2000) 26 D Mondelaers, G Vanhoyland, H V Rul, J D’Haen, M K Van Bael, J Mullens, and L C V Poucke, Mater Res Bull 37, 901 (2002) 27 L N Wang and M Muhammed, J Mater Chem 9, 2871 (1999); Q S Yi, X M Wu, and Z C Tan, J Inorg Mater 16, 620 (2001); J E Rodriguez-Paez, A C Caballero, M Villegas, C Moure, P Duran, and J F Fernandez, J Eur Ceram Soc 21, 925 (2001) 28 L Guo, Y L Ji, H B Xu, P Simon, and Z Y Wu, J Am Chem Soc 124, 14864 (2002); B Liu and H C Zeng, J Am Chem Soc 125, 4430 (2003); P V Radovanovic, N S Norberg, K E McNally, and D R Gamelin, J Am Chem Soc 124, 15192 (2002); Z R R Tian, J A Voigt, J Liu, B Mckenzie, and M J Mcdermott, J Am Chem Soc 124, 12954 (2002) 29 J Y Lao, J Y Huang, D Z Wang, and Z F Ren, Nano Lett 3, 235 (2003) 780 30 J Y Lao, J G Wen, and Z F Ren, Nano Lett 2, 1287 (2002) 31 W I Park, G.-C Yi, M Y Kim, and S J Pennycook, Adv Mater 14, 1841 (2002) 32 H Q Yan, R R He, J Johnson, M Law, R J Saykally, and P D Yang, J Am Chem Soc 125, 4728 (2003) 33 (a) J Zhang, L D Sun, J L.Yin, H L Su, C S Liao, and C H Yan, Chem Mater 14, 4172 (2002) (b) A Chittofrati and E Matijevic, Colliod Surf 48, 65 (1990) 34 (a) J Zhang, L D Sun, C S Liao, and C H Yan, Chem Commun 3, 262 (2002) (b) J Zhang, L D Sun, H Y Pan, C S Liao, and C H Yan, New J Chem 26, 33 (2002) 35 (a) E A Meulenkamp, J Phys Chem B 102, 5566 (1998) (b) M Hilgendorff, L Spanhel, C Rothenhäusler, and G Müller, J Electrochem Soc 145, 3632 (1998) 36 T Matsushita, K Kodaira, and J Satio, J Cryst Growth 26, 147 (1974); S D Sharma and S C Kashyap, J Appl Phys 42, 5302 (1971); N Sakagami, J Cryst Growth 99, 905 (1990); H Nishizawa, T Tani, and K Matsuka, in “Committee of the American Ceramic Society,” 1984, p C-99 37 B G Wang, E W Shi, and W Z Zhong, Cryst Res Tech 33, 937 (1998) 38 L N Wang and M Muhammed, J Mater Chem 9, 2871 (1999) 39 (a) Y Li, G S Cheng, and L D Zhang, J Mater Res 15, 2305 (2000) (b) G Shi, C M Mo, W L Cai, and L D Zhang, Solid State Commun 115, 253 (2000) (c) S H Yu and M Yoshimura, Adv Mater 14, 296 (2002) (d) Z Wang and H L Li, Appl Phys A 74, 201 (2002) 40 Y C Wang, I C Leu, and M H Hon, J Mater Chem 12, 2439 (2002); Y C Wang, I C Leu, and M H Hon, Electrochem Solid St 5, C53 (2002); Y C Wang, I C Leu, and M H Hon, J Cryst Growth 237, 564 (2002) 41 (a) M J Zheng, L D Zhang, G H Li, and W Z Shen, Chem Phys Lett 363, 123 (2002) (b) Z X Zheng, Y Y Xi, P Dong, H G Huang, J Z Zhou, L L Wu, and Z H Lin, Phys Chem Comm 9, 63 (2002) (c) Y Li, G W Meng, L D Zhang, and F Phillipp, Appl Phys Lett 76, 2011 (2000) (d) W Choi, T W Kang, and L A Nosova, J Korean Phys Soc 42, S216 (2003) Zinc Oxide Nanostructures 42 W H Zhang, J L Shi, L Z Wang, and D S Yan, Chem Mater 12, 1408 (2000) 43 B B Lakshmi, P K Dorhout, and C R Martin, Chem Mater 9, 857 (1997); B B Lakshmi, C J Patrissi, and C R Martin, Chem Mater 9, 2544 (1997) 44 M C Neves, T Trindade, A M B Timmons, and J D P de Jesus, Mater Res Bull 36, 1099 (2001) 45 R F Mulligan, A A Iliadis, and P Kofinas, Abstr Pap Am Chem S 220, 278-PMSE (2000) 46 D S Boyle, K Govender, and P O’Brien, Chem Commun 80 (2002) 47 K Keis, S E Lindquist, and A Hagfeldt, J Phys Chem B 105, 3350 (2001) 48 L Vayssieres, K Keis, A Hagfeldt, and S E Lindquist, Chem Mater 12, 73 (2001) 49 W J Li, E W Shi, W Z Zhong, and Z W Yin, J Cryst Growth 203, 186 (1999) 50 J Zhang, L D Sun, X C Jiang, C S Liao, and C H Yan, submitted for publication 51 J Zhang, L D Sun, C S Liao, and C H Yan, Solid State Commun 124, 45 (2002) 52 W Z Zhong and J Tang, J Cryst Growth 166, 91 (1996) 53 D F Zhang, L D Sun, J L Yin, and C H Yan, Adv Mater., in press 54 J Q Hu and Y Bando, Appl Phys Lett 82, 1401 (2003); L Vayssieres, K Keis, A Hagfeldt, and S E Lindquist, Chem Mater 13, 4395 (2001) 55 R A Bride, J Melley, and D E Cormack, J Mater Chem 13, 1196 (2003) 56 C H Lu and C H Yeh, Ceram Int 26, 351 (2000) 57 D B Zhang, L M Qi, J M Ma, and H M Cheng, Chem Mater 13, 2753 (2001); L M Qi, J Li, and J M Ma, Adv Mater 14, 300 (2002) 58 C Pacholski, A Kornouski, and H Weller, Angew Chem Int Ed 41, 1188 (2002) 59 K Vanheusden, W L Warren, C H Seager, D R Tallant, J A Voigt, and B E Gnade, J Appl Phys 79, 7983 (1996) ... nucleated at particle ends 5.1 10 –1 s s - 60 1.50 1.25 1.00 0.75 0.50 0.25 0.00 -10 10 10 10 10 10 10 5 10 Pulse Width (s) Figure 21 Coercive field as a function of an applied field duration for... (2) Encyclopedia of Nanoscience and Nanotechnology Edited by H S Nalwa Volume 10: Pages (27–42) 28 Sol–Gel Derived Semiconductor Oxide Gas Sensors The result is the formation of the M O M bond... and a phenomenological damping term Since the period of the precessional motion of a magnetic moment is of the order of 10? ??11 s, the total elapsed time of these calculations is of the order of