Recent Advances in Tin Dioxide Materials: Some ... - ACS Publications

Review pubs.acs.org/CR

Recent Advances in Tin Dioxide Materials: Some Developments in Thin Films, Nanowires, and Nanorods Zhiwen Chen,*,†,§,∥ Dengyu Pan,‡,§ Zhen Li,†,§ Zheng Jiao,*,†,‡,§ Minghong Wu,*,†,‡,§ Chan-Hung Shek,*,∥ C. M. Lawrence Wu,∥ and Joseph K. L. Lai∥ †

Shanghai Applied Radiation Institute, ‡Institute of Nanochemistry and Nanobiology, §School of Environmental and Chemical Engineering, Shanghai University, Shanghai 200444, People’s Republic of China ∥ Department of Physics and Materials Science, City University of Hong Kong, Tat Chee Avenue, Kowloon Tong, Hong Kong Biographies Acknowledgments List of Abbreviations References

7478 7480 7480 7480

1. INTRODUCTION 1.1. Preamble

The key scientific issues in the application and development of semiconductor micro/nanodevices and optoelectronics components have driven scientists to investigate in depth the design, preparation, micro/nanostructure, and performance of semiconductor materials.1−10 Semiconductor oxides are fundamental to the development of smart and functional materials, devices, and systems.11−15 These oxide materials have two unique structural features: mixed cation valences and an adjustable oxygen deficiency, which are the bases for creating and tuning many novel material properties, from chemical to physical.16−19 An integrated device for semiconductor industry is highly desirable for versatile advanced applications. The prospect of various preparation processes to fabricate semiconductor oxide materials continues to drive research toward improving the performance of the semiconducting materials utilized in these devices.20−23 Since the properties of materials strongly depend on its micro/nanostructures, which all result from the fabrication processes, the influence of micro/ nanostructural evolution on material properties is of special interest in the field of materials science and engineering.24−32 However, challenges remain on how to further improve the material fabrication processes for various advanced applications. This optimization requires a clear understanding of the relationship between micro/nanostructures and their morphologies.33−40 Tin oxide micro/nanostructures have garnered considerable attention in recent years for their potential to facilitate both fundamental research and practical applications through their advantageous chemical and physical properties.41−51 Tin oxide is a unique material of widespread technological applications, particularly in the field of gas sensors,52−56 dye-based solar cells,57 transparent conducting electrodes,58 and catalyst supports.59 New assessment strategies for tin oxide functional materials are of fundamental importance in the development of micro/nanodevices. 60 However, the as-grown tin oxide

CONTENTS 1. Introduction 1.1. Preamble 1.2. Outline 2. Tin Dioxide Thin Films 2.1. Overview on Tin Dioxide Thin Films 2.2. Production and Fractal Assessment Strategies of Tin Dioxide Thin Films 2.3. Annealing Effects of Tin Dioxide Thin Films 2.4. Defect Evolution of Tin Dioxide Thin Films 2.5. Optical Properties of Tin Dioxide Thin Films 2.6. Electrical Properties of Tin Dioxide Thin Films 3. Tin Dioxide Nanowires 3.1. Overview on Tin Dioxide Nanowires 3.2. Preparation and Characterization of Tin Dioxide Nanowires 3.3. Microstructure and Growth of Tin Dioxide Nanowires 3.4. Defect and Properties of Tin Dioxide Nanowires 3.5. Modified Tin Dioxide Nanowires and Composites 4. Tin Dioxide Nanorods 4.1. Overview on Tin Dioxide Nanorods 4.2. Synthesis and Growth of Tin Dioxide Nanorods 4.3. Doping and Modification of Tin Dioxide Nanorods 4.4. Tin Dioxide Nanorods Planted Graphites 5. Conclusions 5.1. Summary 5.2. Concluding Remarks Author Information Corresponding Authors Notes © 2014 American Chemical Society

7442 7442 7443 7443 7443 7444 7448 7451 7456 7458 7460 7460 7460 7462 7464 7467 7470 7470 7470 7473 7475 7477 7477 7478 7478 7478 7478

Received: December 28, 2013 Published: June 19, 2014 7442

dx.doi.org/10.1021/cr4007335 | Chem. Rev. 2014, 114, 7442−7486

Chemical Reviews

Review

Figure 1. Schematic diagram illustrating the various scientific perspectives of tin dioxide thin films, nanowires, and nanorods.

materials typically possess a high density of defects,61 which would degrade their properties. Therefore, the synthesis of defect-free tin oxide materials is of great interest. In order to provide guidance for the search of better tin dioxide (SnO2) functional materials with suitable optical and electrical properties, it is necessary to investigate various effects of SnO2 thin films and nanostructured materials.62−68 For example, it was found that the influence of annealing temperature on material properties is especially remarkable.69 However, some challenges still remain in the clarification of the intricate aspects of SnO2 thin films and nanostructured materials as well as their applications.

microstructure, growth, defects, properties, structural modification, and composites of SnO2 nanowires will be discussed in detail. Finally, we turn our attention to SnO2 nanorods, emphasizing the synthesis, growth, doping, and structural modification of SnO2 nanorods and SnO2 nanorods planted graphites. This information may inspire a novel approach to improve their performance and promote better design of micro/nanodevices. Once their unique properties have been understood and mastered, SnO2 functional materials with a variety of fascinating micro/nanostructures will offer vast and unforeseen opportunities in the semiconductor industry as well as in other fields of science and technology.

1.2. Outline

2. TIN DIOXIDE THIN FILMS

This review summarizes the results of recent research and development on the micro/nanostructures and properties of SnO2 functional materials, including thin films, nanowires, and nanorods. It mainly focuses on the synthesis of SnO2 thin films, nanowires, and nanorods by using various preparation techniques, including pulsed laser deposition (PLD), microemulsion, etc., and their characterization and application. It is an interdisciplinary work that integrates the areas of chemistry, physics, and materials science. These results may enable SnO2 functional materials with appropriate micro/nanostructures to be tailor-made for a large number of applications and provide new opportunities for future studies of SnO2 architectures with the goal of optimizing functional material properties for specific applications. In order to facilitate reading, we provide below a sketch (Figure 1) of the outline of this review encompassing various effects, microstructural evolution, and related properties of SnO2 thin films, nanowires and nanorods. This schematic diagram illustrates the various scientific perspectives of SnO2 thin films, nanowires, and nanorods, including preparation, heat treatment, growth, doping, modification of micro/nanostructures, and related properties. First, we present the preparation and formation processes of SnO2 thin films analyzed by fractal assessment strategies. Annealing effects, defect evolution, and optical and electrical properties will be described. Second, we focus on SnO2 nanowires. The preparation, characterization,

2.1. Overview on Tin Dioxide Thin Films

In recent years, semiconductor oxides become one of the most advanced materials for a wide range of industrial applications.70−73 Key components for these applications are wide band gap semiconductors with oxides of different origins serving as passive as well as active components, similar to that in conventional semiconductors. Special attention has been focused on tin oxide during the past few decades and this material is now established as one of the most promising materials for producing numerous novel functional properties due to its excellent chemical and physical properties.74−77 In view of the increasing public concern on air pollution and the need to monitor concentration levels of gases such as CO, CO2, NOx, O3, SO2, etc., research and development of many kinds of sensors and control systems has begun in earnest in recent years. Many combustible and toxic gases such as CO, NH3, NO2, H2S, and CH4 can be detected by sensors constructed using SnO2 as a gas sensing material.78 Sintered SnO2 powders are often used in commercial sensors, but SnO2 thin films are now gaining increasing popularity.79,80 Reduction in device size and a concomitant increase in the speed of response can be achieved with the advent of advanced thin film technology. More cost-effective and reproducible devices can now be made by using SnO2 thin films. 7443


Chemical Reviews

Review

prepared by PLD techniques will be described in the next section. Results on the experimental preparation of SnO2 thin films at different substrate temperatures with interesting fractal features will be reviewed. The review also includes the investigation of the micro/nanostructure evolution of SnO2 thin films using X-ray diffraction (XRD) and scanning electron microscopy (SEM), evaluation of its structure by fractal methodology, and the correlation between fractal dimensions and substrate temperature in SnO2 thin films. The experimental evidence indicating the formation of fractal clusters with various sizes, densities, and fractal dimensions in SnO2 thin films prepared under different substrate temperatures will be reviewed. The review also highlights the unusual formation of significant fractal features and their sensitive dependence on substrate temperature, a key parameter affecting gas sensing behavior. The accumulated knowledge on this material today may enable application-specific novel SnO2 functional materials with appropriate fractal structures to be tailor-made. Apart from the immediate use in the monitoring of harmful gases in the environment, the information presented may stimulate new ideas in future studies of fractal structure SnO2 architectures, with the ultimate goal of optimizing material properties for specific applications.

The number of electrons in the conduction band of SnO2 is affected by the adsorption of gaseous species on its surface.81 It is this key property which enables this material to be used as a gas sensor. Gas molecules, such as CO, are oxidized by oxygen species (O2− and O−) on the semiconductor surface.82 As a consequence, the height of the Schottky barrier of the material is lowered and the conductance is increased.83,84 SnO2 gas sensors have high sensitivity to humidity and inflammable gases and are therefore extensively used in many situations. Change in the material’s electrical impedance due to the adsorption of gas molecules on the SnO2 surface is utilized as a means to measure the concentration of various gases in this type of sensors. In a normal air environment, oxygen is adsorbed by the SnO2. This captures its electrons and raises its electrical resistivity. When a reducing gas is present, the SnO2 resistivity decreases because the former competes for the adsorbed oxygen. With less adsorbed oxygen, fewer electrons are captured and the resistivity decreases. However, material fabrication parameters can also have a strong influence on the electrical properties of SnO2. For high sensitivity, a small SnO2 grain size is desirable. This is because gas sensing is based on adsorption mechanisms on the SnO2 grain surface. A small grain size would lead to a high adsorption area per unit volume, thereby increasing the sensor’s sensitivity.85,86 The scientific principle underpinning the operation of resistive gas sensors is based on the correlation between the concentration of certain ambient gases and the surface conductance of semiconducting oxides. The nature of the sensing mechanism is related to the electrical response of the gas sensor to reactive gases and the change in the sensor’s surface resistance provides an indication of the gas concentration.87 Operation of these sensors involves interaction between gaseous molecules and microstructural defects on the surface and at grain boundaries. Although the detailed mechanism of interaction is complicated, this does not restrict the versatility of these sensors and they can be used to detect oxygen, flammable gases and common toxic gases. The surface conductance can increase or decrease, depending on whether the solid has n-type or p-type conductivity.88,89 Factors that can significantly affect the gas sensing properties of semiconducting SnO2 include microstructural characteristics such as grain size, grain geometry, and specific surface area. It is necessary to understand the detailed mechanism of micro/nanostructure evolution of SnO2 thin films in order to control these microstructural characteristics. One potentially powerful technique to characterize these micro/nanostructures is fractal method. This technique has been applied to analyze microstructural features in SnO2 thin films.60 In the following sections, we will show some examples of geometric microstructures of SnO2 thin films. We will then discuss in detail the applicability and relevance of fractal theory in the analysis of the micro/nanostructure and gas sensing behavior of SnO2 thin films. In order to improve versatility in advanced applications, an integrated sensing device for different gas species is highly desirable. Although SnO2 has high sensitivity to many gas species, it suffers from certain drawbacks. For example, it is often susceptible to electrical drift which entails long stabilization periods. Moreover, after extended periods of operation the material can be permanently poisoned. Fractal assessment strategies can shine a new light on the future development of micro/nanodevices made from this material. A new insight on fractal assessment strategies of SnO2 thin films

2.2. Production and Fractal Assessment Strategies of Tin Dioxide Thin Films

One method to produce SnO2 thin films is to use PLD. The starting point is the production of a sintered SnO2 target. This involves the synthesis of pure nanocrystalline SnO2 powder which can be done by the sol−gel method.90 The fabrication method is described in detail in the following. A cold ethanol solution of SnCl4 (27%) is treated with an aqueous ammonia solution (28%) until a suitable pH value is reached. This process produces meta-stannic acid sol (parent sol) by precipitation. The parent sol is washed repeatedly by deionized water. After drying, powder with average grain size of about 4 nm could be obtained. The powder is then compacted under uniaxial pressure of about 0.4 GPa, and sintered at 1150 °C for 2 h to produce SnO2 discs. This process should produce high purity (about 99.8%) SnO2 with a cassiterite structure. The sintered SnO2 disc can then be used to make SnO2 thin films by the pulse laser deposition technique. It is important to clean the target with methanol in an ultrasonic bath before installation to minimize contamination. A KrF excimer laser (e.g., Lambda Physik, LEXtra 200, Germany) producing pulse energies of 350 mJ at a wavelength of 248 nm and a frequency of 10 Hz could be used. Successful thin films could be made with duration of 34 ns for each excimer laser pulse. A typical setup involves a high vacuum chamber with an ultraviolet (UV)-grade fused silica window made of UV-grade fused silica lens through which laser energy is transmitted onto the target. It is necessary to rotate the target during the thin film production process in order to avoid excessive wear of material at a fixed location (i.e., drilling). A rotation rate of about 15 rpm is recommended. A typical setting involves laser energy intensity of 5 J/cm2, corresponding to a total of approximately 1.5 × 105 laser pulses. This should produce SnO2 thin films with a growth rate of about 0.3 nm/s (or about 1 μm/h). The substrate for the deposition of laser ablated material could be Si (100) mounted on a holder a few cm away from the target. The temperature of the substrate is an important parameter and is typically in the range 300−450 °C. It is necessary to maintain high vacuum in the deposition chamber with a cryogenic pump 7444


Chemical Reviews

Review

with the base pressure maintained at about 1 × 10−6 mbar prior to laser ablation. The oxygen partial pressure during laser ablation should be set at about 3 × 10 −2 Pa. Characterization of the deposited thin films usually includes XRD and SEM. The Fractal Images Process Software (FIPS) may be used to analyze SEM images. A typical method involves dividing these digitized SEM images into boxes of 360 × 360 size followed by analysis using the fractal method.91 A suitable number of intact fractal patterns could be selected from these digitized images for evaluation. The box counting method could be used to determine the average values of fractal dimensions (D), the fractal density and the average size of the fractal clusters from these digitized fractal patterns.92−96 After installing the SnO2 thin film into the sensor, the following technique may be used to assess its carbon monoxide (CO) gas sensing property. Test CO gas is introduced into a chamber. The gas concentration is controlled by an injector with variable volume. A typical range is 25−500 ppm. Electrical measurement may then be performed after the sensor is stabilized. The variation of electrical properties (e.g., resistivity) with carbon monoxide concentration determined will form the basis of assessment of the sensitivity of the gas sensor. The mineral form of SnO2 is known as cassiterite. It has a tetragonal rutile crystalline structure with point group D14 4h and space group P42/mnm. There are two metal atoms and four oxygen atoms in the unit cell. The lattice parameters are a = 4.7382(4) Å, and c = 3.1871(1) Å. These crystallographic data form the basis of characterization of the SnO2 thin film by XRD. Figures 2A, B, C, and D show XRD patterns of SnO2 thin

peaks relative to the background. The crystallinity of the thin films is enhanced by increasing substrate temperature. This is manifested by the intensity and sharpness of the XRD peaks of the SnO2 thin films. This dependence on substrate temperature dependence can be explained by the mobility of the atoms in the thin films. The surface mobility of the vapor species is low at low substrate temperatures. As a consequence, they are scattered at different positions on the surface. This low mobility of the species will inhibit full crystallization of the thin films. As the substrate temperature increases, the vapor species will have sufficient mobility to arrange themselves at the lower energy positions in the crystalline cell.97−99 The SnO2 average grain size can be determined using the Scherrer formula: D = Kλ/β cos θ, where D is the diameter of the nanoparticles. In the data presented in Figure 2, K = 0.9, λ(Cu Kα) = 1.5406 Å, and β is the full-width-at-half-maximum of the diffraction lines. Calculation using the above method shows that the average grain size of the SnO2 nanoparticles at different substrate temperatures falls within the range of 25.3−27.8 nm. Nanoparticle grain size increases from 25.3 nm at 300 °C to 26.2 nm at 350 °C. Further increase in substrate temperature to 400 and 450 °C increases the nanoparticle grain size to 27.0 and 27.8 nm, respectively. The data suggest that SnO2 nanostructures can work as sensitive and selective chemical sensors. SnO2 nanostructure sensor elements can be configured as resistors or as barrier junction devices. In the former, the conductance can be modulated by charge transfer across the surface. In the latter, the properties can be controlled by applying an electrical potential across the junction. With a better understanding of the influence of significant micro/ nanostructural features, further improvement of sensing ability can be achieved through such functionalization of the surface. Figure 3 presents SEM images of SnO2 thin films prepared on Si (100) substrate at temperatures of (A) 300, (B) 350, (C) 400, and (D) 450 °C. These figures show self-similar fractal patterns in all these thin films, even though they were produced under different substrate temperatures.60 It can be seen from Figure 3 that the fractal patterns are open and loose structures, becoming coarser with increasing substrate temperature. The average sizes of the fractal patterns (or clusters) estimated by measurements undertaken on the fractal regions are about 0.307 (Figure 3A), 0.906 (Figure 3B), 1.202 (Figure 3C), and 1.608 μm (Figure 3D). A typical measuring procedure involves choosing a number of fractal patterns at random for each SEM image for average size determination.100−104 It was found that increasing substrate temperature results in an increase in the average size of the fractal clusters. Fractal structure is characterized by its box size, L, and the number of boxes occupied by the fractal clusters, N. A plot of ln(N) versus ln(1/L) would indicate scale invariance if the graph is linear. Figure 4 shows such plots of ln(N) versus ln(1/ L) of the fractal cluster regions in Figure 3. It can be seen that all plots are reasonably linear, indicating that the morphologies of SnO2 clusters have scale invariance within these ranges. Thus, the SnO2 clusters have fractal properties. The fractal dimension (D) can be obtained by fitting a linear relationship to the function ln(N) versus ln(1/L). When the data presented in Figure 4 are analyzed in this way, the fractal dimensions obtained are 1.896 at 300 °C, 1.884 at 350 °C, 1.865 at 400 °C, and 1.818 at 450 °C. It is apparent that the fractal dimension (D) decreases with increasing substrate temperature. Smaller fractal dimension implies finer branches in the open and loose fractal structure of the SnO2 thin films. Figure 5A−C show the

Figure 2. XRD patterns of SnO2 thin films prepared on Si (100) substrate at temperatures of (A) 300, (B) 350, (C) 400, and (D) 450 °C. Reprinted with permission from ref 60. Copyright 2010, American Chemical Society.

films prepared on Si (100) substrate at 300, 350, 400, and 450 °C respectively.60 Using the standard values for bulk SnO2 (International Center for Diffraction Data (ICDD), PDF file no. 77-0447), the major diffraction peaks of some lattice planes can be indexed to the tetragonal unit cell structure of SnO2 with lattice constants a = 4.738 Å and c = 3.187 Å. The (110), (101), (200), (211), (220), and (002) peaks are indicated in the figure. If the SnO2 was contaminated with crystalline impurities, these may be manifested as additional peaks in the XRD pattern. Presence of noncrystalline impurities can also be revealed by enhanced noise background in these figures. In Figure 2, the thin films are predominately pure crystalline SnO2, indicated by the high intensity of the relevant crystalline (hkl) 7445


Chemical Reviews

Review

Figure 3. SEM images of SnO2 thin films prepared on Si (100) substrate at temperatures of (A) 300, (B) 350, (C) 400, and (D) 450 °C. Reprinted with permission from ref 60. Copyright 2010, American Chemical Society.

Figure 4. Plots of ln(N) versus ln(1/L) of the fractal cluster regions in Figure 3, where L is the box size and N is the number of boxes occupied by the SnO2 crystalline structure for substrate temperatures at (A) 300, (B) 350, (C) 400, and (D) 450 °C. Reprinted with permission from ref 60. Copyright 2010, American Chemical Society.

Figure 5. (A) Fractal average size; (B) fractal dimension; (C) fractal density versus substrate temperature. Reprinted with permission from ref 60. Copyright 2010, American Chemical Society.

relationship between the fractal average size, fractal dimension and fractal density with substrate temperature. The fractal average size increases with increasing substrate temperature, while the opposite is true for the fractal dimension and fractal density. In general, the initial nucleation rate of the core crystal is a key factor in determining the fractal density. In Figure 5C, the fractal densities are 18, 6, 3, and 2 mm−2 at 300, 350, 400, and 450 °C, respectively. These results show that fractal density

gradually decreases with increasing substrate temperature. The nucleation rate could be affected by strain relaxation effects in the material’s temperature field during annealing. At 300 °C, the fractal density and fractal occupation area are high due to the relatively short-range temperature field. At higher substrate temperature, the relatively long-range temperature field may promote new nuclei and subsequent growth, leading to fractal 7446


Chemical Reviews

Review

Figure 6. Formation process of SnO2 nanocrystals and fractal clusters. (A) laser; (B) target; (C) plasma; (D) plume; (E) nucleation; (F) grain rotation; (G) coalescence; (H) growth; and (I) fractal. Reprinted with permission from ref 60. Copyright 2010, American Chemical Society. (i) Laser ablation of the cassiterite SnO2 target. (ii) Interaction between the pulsed laser and SnO2 target to produce high-temperature and highpressure SnO2 plasma at the solid−liquid interface. (iii) Cooling of the SnO2 as a result of subsequent expansion of the high-temperature and highpressure SnO2 plasma.105−108 Since the interval between two successive laser pulses is much longer than the life of the plasma, the next laser pulse does not interact with the plasma produced by the earlier pulse. (iv) The SnO2 plume is deposited on the Si (100) substrate after the disappearance of the plasma. Nucleation of SnO2 nanocrystals begins. (v) Grain rotation occurs to minimize the area of high energy interfaces, culminating in a lowenergy configuration.109,110(vi) Coalescence takes place, with the consequence of replacing high energy grain boundaries by lower energy coherent grain boundaries due to grain rotation, and the formation of larger SnO2 nanocrystals. (vii) Growth of SnO2 nanocrystals along preferred crystallographic directions. These directions could be predicted by an analysis of the surface energy in several crystallographic orientations. (viii) As SnO2 crystallizes and nucleates at high energy interfaces such as grains boundaries, fractal structure is formed.

growth of fine branches and a concomitantly lower fractal density. The above experimental observations have led to the development of a novel model to describe the formation process of SnO2 nanocrystals and fractal clusters.60 The model consists of eight steps, as illustrated in detail in Figure 6. When SnO2 crystals are formed, latent heat of crystallization is released. This leads to a rise in temperature in the vicinity of the crystal. As this temperature field propagates, fresh nucleation occurs in the surrounding region.111,112 The newly formed nuclei will in turn release latent heat and stimulate further nucleation. This iterative process continues until SnO2 fractal patterns are formed. Figure 6 illustrates the formation process of SnO2 nanocrystals and fractal structures based on the above proposed mechanism. Laser ablation is an appropriate method in the synthesis of environmentally functional materials with controlled composition, morphology and nanocrystal size. These are parameters which the gas detection sensitivity of SnO2 thin films depend. Carbon monoxide is one of the common toxic gases which gas sensors are designed to detect. Thus, the sensitivity of SnO2 thin films to carbon monoxide concentration is a topic of major

importance. Figure 7 shows the CO gas sensing behavior of SnO2 thin films prepared on Si (100) substrate at (A) 300, (B) 350, (C) 400, and (D) 450 °C. The sensitivity was determined at room temperature with CO concentrations of 25, 50, 75,

Figure 7. CO gas sensing behavior of SnO2 thin films prepared on Si (100) substrate at temperatures of (A) 300, (B) 350, (C) 400, and (D) 450 °C. Reprinted with permission from ref 60. Copyright 2010, American Chemical Society. 7447


Chemical Reviews

Review

used to deposit SnO2 thin films, such as RF-sputtering,128 DCmagnetron sputtering,129 thermal evaporation,130,131 ion beam deposition,132 chemical vapor deposition,133−137 spray pyrolysis,138 successive ionic layer deposition (SILD),139 and other chemical methods.123,140 Sberveglieri has reviewed the techniques used for tin oxide thin films deposition.141 High annealing temperature is usually required in order to fabricate good quality polycrystalline films. However, the surface of the thin films can be damaged by elevated temperature. Moreover, the interface thickness also increases with increasing annealing temperature. This may have deleterious effects on optical properties, especially waveguide efficiency. Pulsed laser deposition is a thin film growth technique in which photonic energy is coupled to the bulk starting material via electronic processes.142,143 The technique involves focusing an intense laser pulse onto a target made from the starting material. The laser pulse passes through an optical window of a vacuum chamber. Significant material removal occurs in the form of an ejected luminous plume above a certain power density. Depending on the target material, its morphology, and the laser pulse wavelength and duration, the threshold power density needed to produce a plume might be of the order of 10−500 MW/cm2 for ablation using ultraviolet excimer laser pulses of 10 ns duration. Thin film growth takes place after the material from the plume is allowed to recondense on a substrate. The ablation plume species in the gas phase and the surface reaction could be affected by the presence of a passive or reactive gas, or an ion source. These may be introduced to supplement the process.144 Pulse laser deposition is a flexible and versatile process for thin film growth and the types of thin films that could be produced by this method are very diverse. The advantages of pulse laser deposition over other deposition techniques are its pulsed nature, and the ability to assemble molecules at the substrate surface to form a material of composition which is far from that at thermal equilibrium. Under favorable conditions, it has the ability to reproduce in thin films the same elemental ratios of highly chemically complex bulk ablation targets.145 The underlying physical processes in pulse laser deposition are highly complex and interrelated. Important factors include laser pulse parameters and the properties of the target material.146 There are many advantages of thin film growth using laser ablation: (i) there is a greater degree of flexibility in the use of materials and geometrical arrangements because, unlike vacuum-installed devices, the energy source (laser) is outside the vacuum chamber; (ii) ablation can be applied to almost any condensed matter material; (iii) thin film growth rates may be controlled to any desired amount due to the pulsed nature of pulse laser deposition; (iv) only the material localized at the area defined by the laser focus will be vaporized; (v) even for chemically complex systems, the ratios of the elemental components of the bulk and thin films can be controlled to tight limits under optimal conditions; (vi) the kinetic energy of the ablated species can be controlled to lie within a range that promotes surface mobility while avoiding bulk displacements; (vii) it has the potential to produce novel or metastable materials that would be unattainable under ordinary thermal conditions due to its ability to produce species with electronic states far from the chemical equilibrium configuration. Pulse laser deposition techniques have been successfully applied to the growth of quality tin oxide thin films produced by ablation of either tin target in oxidizing atmosphere124 or SnO2 target,72,127,147 The use of laser light, control of the

100, 200, 300, 400, and 500 ppm. The sensitivity increases with increasing CO concentration and substrate temperature. Other data produced by Cooper and Cicera using different procedures also indicate that SnO2 thin film sensors have higher sensitivity to CO.113,114 A better understanding of the role of the fractal structure in gas sensing is necessary for further development of this type of SnO2 thin film gas sensors. Experimental results obtained so far show that the CO gas sensing behavior clearly depends on the fractal dimension, fractal density, and average size of the fractal clusters (Figures 5 and 7). A random tunnelling junction network (RTJN) mechanism has been proposed to explain this gas sensing behavior.60 After the fractal structure is formed, typical fractal clusters in SnO2 thin films comprise of SnO2 grains with the morphology of fine dendrite-like nanocrystals incorporating many tunnelling junctions of varying sizes. The whole SnO2 thin film is considered to be made up of a series of tunnelling junctions from the electron transport perspective. Different substrate temperatures produce SnO2 thin films with fractal branches of different dimensions. This leads to differences in the height of the Schottky barrier of the tunnelling junctions, resulting in different breakdown voltages. Gas detection relies on the fact that reducing gas molecules such as CO will react with the oxygen species (O2− and O−) ionized on the surface of the SnO2 particles. This lowers the height of the Schottky barrier, and increases the conductance.82−84 For example, let us consider a junction i in a SnO2 thin film deposited at a relatively low substrate temperature of 300 °C. In view of the relatively large fractal dimension produced at this substrate temperature, the junction i will possess high resistance due to thick fractal branches. Thus, it is more difficult for the external voltage Vi to lower the Schottky barrier Si and the junction i cannot be easily broken. This results in low gas sensitivity (Figure 7A). On the other hand, SnO2 thin film deposited at high substrate temperatures (e.g., 450 °C) will have small fractal dimension. Fine fractal branches will lead to low junction resistance. In this case, the external voltage can readily lower the Schottky barrier and the junction can be easily broken, resulting in high gas sensitivity (Figure 7D). The fractal dimension and the size of the fractal branches are closely related. The number of the fine branches increases with decreasing fractal dimension. Therefore, the smaller the fractal dimension, the larger will be the number of junctions with the smaller Schottky barrier Si in the lower resistance state. This is an important discovery. It shows how fractal studies can lead to better SnO2 thin film gas sensors to be designed. 2.3. Annealing Effects of Tin Dioxide Thin Films

Tin dioxide (SnO2) has high electrical conductivity (8 × 10−4 Ωcm), high transparency in the visible wavelength range, high chemical stability, and a direct band gap.115−117 It has been widely used in many different situations as an n-type semiconductor (Eg = 3.6 eV at 300 K). Examples of applications of this material include solid-state gas sensors, liquid crystal displays, photovoltaic cells, and transparent conducting electrodes.72,118−120 SnO2 is also a material of relatively low cost. This makes it particularly attractive for commercial devices where cost competitiveness is a major factor in material selection.121−124 Annealing temperature can have a significant influence on the microstructure and defects density in SnO2 thin films. In recent years, SnO2 thin films have attracted much attention due to their high surface area.125−127 Many different techniques can be 7448


Chemical Reviews

Review

composition of deposited structure and in situ doping have led to reduced contamination in the thin film production process. By varying the experimental deposition conditions, it can produce nanoparticles of desired size and composition.142 The effects of annealing on microstructural change, and morphological properties, and root-mean-square (RMS) surface roughness of SnO2 thin films have been investigated.148 SnO2 thin films deposited on glass substrates by pulse laser deposition techniques have been investigated before and after annealing in the temperature range from 50 to 550 °C at 50 °C intervals. XRD, SEM, transmission electron microscopy (TEM), and selected area electron diffraction (SAED) are applied to study the influence of annealing temperature on the microstructural and morphological properties of SnO2 thin films. With increasing annealing temperatures, transformation of the amorphous microstructure into polycrystalline SnO2 phase with preferred orientations related to the (110), (101), and (211) crystal planes is observed. The best crystalline properties, viz optimum growth conditions, are found in the thin film annealed at 200 °C. The minimum average rootmean-square roughness of 20.6 nm with average grain size of 26.6 nm is observed in the thin film annealed at 100 °C. A sintered SnO2 disc is used as the target for pulse laser deposition. Pure SnO2 powder for the circular target consisting of high-purity cassiterite SnO2 (99.8%) is produced by direct oxidation reaction at 1050 °C, viz.: Sn + O2 → SnO2, carried out in a horizontal quartz tube. The target (about ϕ15 mm × 4 mm) is cleaned with methanol in an ultrasonic cleaner before installation to minimize contamination. The laser used for pulse laser deposition is a KrF excimer laser, producing pulse energy of about 350 mJ at a wavelength of 248 nm and a frequency of 10 Hz. The duration of each excimer laser pulse is 34 ns. The laser energy is transmitted through an UV-grade fused silica window with an UV-grade fused silica lens onto the target in a high-vacuum chamber. Rotation of the target at a rate of about 15 rpm is necessary to minimize excessive material loss at the region where the laser beam was focused. The pulse fluence is 5 J/cm2, corresponding to a total of approximately 1.5 × 105 laser pulses. Under these experimental conditions, the growth rate is about 3 × 10−1 nm/s (or about 1 μm/h). To collect the ablated material, a clean glass substrate mounted on a holder 2.5 cm away from the target is used. A cryogenic pump is used to maintain the high vacuum in the deposition chamber. In a typical experiment, the base pressure prior to laser ablation is about 1 × 10−6 mbar, and the working pressure during laser ablation is about 2 × 10−6 mbar. The as-prepared thin films are annealed at a pressure of about 2 × 10−3 Pa at various temperatures ranging from 50 to 550 °C at 50 °C intervals for a fixed time of 30 min at each temperature. Structure of these films is examined by XRD, SEM, TEM, and SAED. Atomic force microscopy (AFM) at atmospheric pressure and room temperature is used to determine the mean RMS surface roughness values for various films taken at different sites of each film. In a typical experiment, AFM measurements are carried out by a digital instrument nanoscope scanning probe microscope in tapping mode using a window size of 6 × 6 μm2 with a resolution of 256 × 256 pixels. A number of sites are selected at random for each AFM image to obtain an average value. The systematic error is affected by many factors, e.g. scan size, pixel resolution, tip−sample force, and tip radius etc., and is of the order of ±0.2 nm.148

Tetragonal phase SnO2 rutile structure is characterized by lattice parameters of a = 4.7382(4) Å and c = 3.1871(1) Å.149 Physically stable thin films with good substrate adhesion could be deposited on clean glass substrates. Figure 8 shows the XRD

Figure 8. XRD patterns of as-prepared and annealed SnO2 thin films at different annealing temperatures for 30 min. (a) as-prepared thin film; (b) 100, (c) 200, (d) 300, (e) 400, and (f) 500 °C. Reprinted with permission from ref 69. Copyright 2009, Elsevier.

patterns of SnO2 thin films annealed at different temperatures for 30 min.69 The thin film deposited at room temperature is predominantly amorphous as manifested by the relatively broad diffraction peaks (Figure 8a). When the annealing temperature is increased to 100 °C, the diffraction peaks become sharper. The thin film becomes more polycrystalline with increasing annealing temperature (Figures 8b-f). This temperature dependence can be explained by the mobility of the atoms in thin films at different temperatures.97−99,150 SnO2 average grain size could be estimated from the XRD data by using the Scherrer formula: D = Kλ/β cos θ. Direct observation using TEM could be used as a comparison. Previous work shows that that the average grain size obtained by TEM is in agreement with those estimated from XRD data.69 Table 1 shows that the Table 1. Average Particle Size versus Annealing Temperature for As-Prepared (RT) and Annealed SnO2 Thin Filmsa annealing temperature (°C) average particle size (nm) annealing temperature (°C) average particle size (nm) a

RT

50

100

150

200

250

27.6

27.4

26.6

24.6

23.7

25.4

300

350

400

450

500

550

27.1

26.6

25.7

26.15

27.7

28.9

Reprinted with permission from ref 69. Copyright 2009, Elsevier.

average grain size of the as-prepared and annealed SnO2 nanoparticles is in the range of 23.7−28.9 nm. SnO 2 nanoparticle size first decreases from 27.6 nm (at room temperature) to less than 23.7 nm at 200 °C, and then sharply increases to larger value of 27.1 nm at 300 °C, followed by gradual decrease to 25.7 nm at 400 °C. Finally, the average grain size increases to 28.9 nm after further annealing at higher temperatures. Typical SEM images of as-prepared and annealed SnO2 thin films are shown in Figure 9. Figure 9a shows the formation of inhomogeneous SnO2 particles at room temperature (RT). Closer examination of this figure reveals the presence of 7449


Chemical Reviews

Review

Figure 9. Typical SEM images of the as-prepared and annealed SnO2 thin films. (a) as-prepared thin film; (b) 100, (c) 200, and (d) 300 °C. Reprinted with permission from ref 69. Copyright 2009, Elsevier.

nanometer sized SnO2 particles. There are agglomerations of smaller particles to form clusters of the order of several micrometers in size. At room temperature, the surface of the deposited SnO2 thin film is rough (Figure 9a). The surface roughness of the thin films changes with annealing temperature. At 100 °C, the surface becomes relatively smooth and dense (Figure 9b). At 200 °C, the surface roughness increases, but it decreases again at 300 °C (Figure 9c,d). Transmission electron micrographs of the as-prepared and annealed SnO2 thin films are shown in Figure 10. In the transmission electron bright-field image of the as-prepared thin film (Figure 10a), the bigger particles dispersed randomly in the amorphous matrix are formed by the agglomeration of smaller particles. In Figure 10a, the particles are within the range of 12−36 nm. SAED patterns of the as-prepared and annealed thin films are shown in Figure 11. In the as-prepared

Figure 11. SAED patterns of the as-prepared and annealed SnO2 thin films. (a) as-prepared thin film; (b) 100, (c) 200, and (d) 300 °C. Reprinted with permission from ref 69. Copyright 2009, Elsevier.

thin film, the SAED pattern is a diffuse ring indicating an amorphous structure with a low degree of crystallinity. Figure 10b displays the transmission electron bright-field image of the thin film annealed at 100 °C. The thin film is very smooth and dense. The size distribution of particles is in the range of 24−29 nm. The SAED patterns shown in Figure 11b shows that the degree of crystallinity of the thin film has increased after the 100 °C annealing treatment. The diffraction rings are sharper and the polycrystalline diffraction rings could be seen. Figure 10c shows the transmission electron bright-field image of the thin film annealed at 200 °C. After the 200 °C annealing treatment, the particle size increases, and the smoothness and density decreases. The particle size after the 200 °C annealing treatment is in the range 17−26 nm. Figure 11c shows that the crystalline structure is clearly revealed. The diffraction ring pattern is consistent with typical tetragonal SnO2 with d200 = 2.37 Å, d210 = 2.12 Å, d310 = 1.50 Å, and d202 = 1.32 Å. The transmission electron bright field image of the thin film annealed at 300 °C is shown in Figure 10d. After the annealing treatment at 300 °C, the film becomes rougher and less dense. The particle size is in the range of 25−32 nm. The degree of crystallinity of the thin film is increased further after the 300 °C annealing treatment. The sharpness and clarity of the diffraction rings have increased. The above experimental results suggest that the quality of the thin film can be controlled by suitable adjustment of the annealing temperature. Semiconducting oxides are of major importance in the development of new functional and smart materials. Microstructure and morphology of semiconducting oxide thin films have strong effects on their unique properties. In applications such as lithography, the development of optical coatings, especially in the ultraviolet range, depends on the microroughness of the thin films.151,152 RMS surface roughness is normally used to characterize an optical surface. The RMS roughness affects light scattering. It can provide an indication of the quality of the surface under investigation. The RMS surface roughness could be affected by the annealing process. Figure 12

Figure 10. TEM bright-field images of the as-prepared and annealed SnO2 thin films. (a) as-prepared thin film; (b) 100, (c) 200, and (d) 300 °C. Reprinted with permission from ref 69. Copyright 2009, Elsevier. 7450


Chemical Reviews

Review

Figure 12. RMS roughness versus annealing temperature for the asprepared and annealed SnO2 thin films. Reprinted with permission from ref 69. Copyright 2009, Elsevier.

shows a plot of RMS surface roughness of as-prepared and annealed SnO2 thin films as a function of annealing temperature. It can be seen that the surface roughness depends strongly on the annealing temperature, which correlates with the result of microstructure analysis. At room temperature, the surface is characterized by the presence of very large particles aggregated in amorphous morphology. The surface roughness at room temperature is very high with average RMS value of about 25.6 nm. As the annealing temperature increases, the average RMS surface roughness value of the deposited thin films decreased sharply. Crystal grain formation becomes apparent at annealing temperature of 100 °C, as indicated by the weakening of the diffuse diffraction rings and the appearance of the sharper polycrystalline diffraction rings. Further manifestation of crystallization occurs in the form of appearance of small crystal grains in the thin films when the annealing temperature is increased to 200 °C. However, the RMS surface roughness increases, probably due to the larger grain size or the increased interfacial strain between the thin film and the substrate. In general, increasing annealing temperature would lead to an improvement of the surface mobility of various atomic species, culminating in a smoother surface. With further increase of annealing temperature, grain growth with preferred orientation is observed. RMS surface roughness of the thin film decreases and the surface smoothness increases. As the film becomes thicker with increasing annealing temperature and the surface becomes smoother, the occurrence of defects and strain variations become more probable. Using power spectral density (PSD) function calculations in sputtered SnO2 thin films, Lindström et al. showed that as the film thickens to greater than 300 nm, defects and strain variations would occur.153 Mesoporous SnO2 is a promising candidate material for gas sensors, catalysts, basic material for the transparent conductive oxides, and optoelectronic materials.154−162 SnO2 in the form of mesoporous structures can retain a large specific surface area and can provide efficient gas transport through highly organized pores.163−168 The controlled formation and calcining behavior of highly crystallized cubic and hexagonal mesoporous SnO2 thin films were reported by Lee et al.169 Figure 13 shows TEM images of mesoporous SnO2 thin films calcined at 400 °C. In Figure 13a,b, the SnO2 thin film appears to have cubic mesostructures oriented along the [100] zone axis. The mesopores, with cavity diameter of approximately 9.1 nm, are regularly ordered. Other specimens of mesoporous SnO2 thin

Figure 13. TEM micrographs of cubic mesoporous SnO2 thin films calcined at 400 °C. (a and b) [100] orientation; (c) [111] orientation; (d) high-resolution image, oriented to the [100] zone axis. Reprinted with permission from ref 169. Copyright 2007, American Chemical Society.

film exhibit orientation to the [111] plane, as shown in Figure 13c. Fast Fourier transform (FFT) images (insets of Figure 13b,c) indicate 4-fold and 6-fold symmetries. The above results indicate that the prepared mesostructure is a cubic phase oriented along the [100] and [111] directions. A high magnification image of the [100] oriented mesoporous structure is shown in Figure 13d. At this increased magnification, more details are revealed. The wall of the mesoporous inorganic frameworks appears to consist of crystallized SnO2 grains. The shape of the individual pores is not uniform. The distortion of the pore structure is probably induced by stress generated during the crystallization of the SnO2 grains. SAED (inset of Figure 13d) shows the characteristic diffuse electron diffraction rings for the polycrystalline cassiterite phase. TEM images of the cubic mesoporous SnO2 thin films calcined at 500 and 600 °C are shown in Figure 14. The mesopores are well-organized even at these elevated temperatures. Fast Fourier transform patterns (insets of Figure 14a,b) show that the SnO2 films calcined at 500 °C are in a wellorganized cubic mesophase. However, the corresponding fast Fourier transform patterns could not be obtained for the SnO2 films calcined at 600 °C. These results suggest that, although the mesopores are uniformly aligned at 500 °C, they begin to be distorted with heat treatment at 600 °C, as shown in the High resolution transmission electron microscopy (HRTEM) image of Figure 14d. 2.4. Defect Evolution of Tin Dioxide Thin Films

A substantial amount of recent research work has been focused on the synthesis of SnO2 nanowires,74 nanotubes,74 nanorods,170,171 nanobelts,71,172,173 and exploration of their novel properties. In these studies, nanocrystalline SnO2 has been reported to exhibit some unique properties different from those of the bulk crystal. With the development of higher quality 7451


Chemical Reviews

Review

achieved.73,82 Properties measured on these higher purity SnO2 thin films are less affected by defects and amorphous components. However, HRTEM has revealed that a considerable number of defects still exist in these improved quality SnO2 thin films.185 Microstructural defects of nanocrystalline SnO2 thin films, prepared by pulse laser deposition, have been investigated using TEM, HRTEM, and Raman spectroscopy.61 Defects in nanocrystalline SnO2 thin films are reported to be significantly reduced by in situ annealing at 300 °C for 2 h. Upon in situ annealing, stacking faults and twins are annihilated as revealed by HRTEM. In fact, the interior the SnO2 nanoparticles are found to be perfect lattices free of defects after in situ annealing. This result is also confirmed by Raman spectroscopy. This suggests that defect-free nanocrystalline SnO2 thin films can be prepared in a simple and practical way using in situ annealing. This is a significant development which enhances the properties of SnO2 thin films for applications as transparent electrodes and solid-state gas sensors. Figure 15 shows a typical TEM bright-field image and the corresponding SAED pattern (inset) of the morphology of asFigure 14. Transmission electron micrographs of cubic mesoporous SnO2 thin films: (a and b) calcined at 500 °C; (c and d) calcined at 600 °C. Reprinted with permission from ref 169. Copyright 2007, American Chemical Society.

SnO2 thin films, much attention has been aroused. This is mainly due to the recognition of the bulk-quantity growth mechanism of these thin films. The success in high quality synthesis of SnO2 thin films has stimulated both experimental and theoretical research in this area. SnO2 thin films or nanoparticles have been prepared by a variety of methods, such as sol−gel,174,175 chemical vapor deposition,176,177 magnetron sputtering,178 sonochemical,179 and thermal evaporation.180 Numerous experimental results on SnO2 thin films have been reported,125,126,147,181 including XRD, TEM, electron transport, and Raman spectroscopy. These results form the basis for the development of potential applications of SnO2 thin films, for example, as candidate material for future transparent nanoelectrodes and solid-state gas nanosensors. Properties of as-grown SnO2 thin films would be subject to degradation due to the presence of a high density of defects.182,183 It is a major scientific goal to achieve synthesis of defect-free SnO2 thin films. In view of the diversity of thin films that can be produced by PLD, this technique is often the first choice in the selection of method of thin film production due to its versatility. The rationale of using PLD in preference to other deposition techniques lies primarily in its pulsed nature. This technique offers the possibility of carrying out surface chemistry far from thermal equilibrium. Under favorable conditions, it has the ability to reproduce in thin films the same elemental ratios of even highly chemically complex bulk ablation targets.145 There are a number of reports on the PLD synthesis of nanocrystalline SnO2 thin films.72,98,127,150,184 The SnO2 thin films produced often contain many defects and amorphous components. The presence of such defects and amorphous components can degrade the properties of the SnO2 thin films and substantially limit their practical use as gas-sensing devices. With the recent refinement in PLD growth techniques, bulkquantities of SnO2 thin films of improved purity can be

Figure 15. TEM image of as-grown SnO2 thin film. The inset is a SAED. Reprinted with permission from ref 61. Copyright 2009, Elsevier.

grown nanocrystalline SnO2 thin film. Many small particles of approximately spherical shape are manifested in the TEM bright-field image. The TEM image of the particles shows different contrast in different regions. This is probably related to different grain sizes of these particles. Microstructural characteristics of the typical tetragonal SnO2 thin films (ring 1: d200 = 2.37 Å, ring 2: d210 = 2.12 Å, ring 3: d310 = 1.50 Å, and ring 4: d202 = 1.32 Å) are confirmed by the polycrystalline diffraction rings of the SAED pattern (inset in Figure 15). The crystallites with dark contrast are close to Bragg orientations. Chemical composition measurement using energy dispersive Xray spectroscopy (EDS) reveals 38.3 at. % tin and 61.7 at. % oxygen in the thin film. The Sn:O ratio = 1:1.611 is a departure from that of bulk SnO2 (Sn:O = 1:2). This would be expected to lead to a large number of oxygen defects at the surface due to the oxygen vacancies and nonstoichiometry SnOx (x < 2) in the as-grown nanocrystalline SnO2 thin film. On the other hand, the density of defects at the interior of the in situ annealed thin films is significantly decreased relative to the as-grown condition. Figure 16a shows the HRTEM image of a typical as-grown thin film. The contrast of the SnO2 thin film exhibits complicated features, indicating the presence of many defects. Stacking faults (marked A, B, C and D) in the nanocrystalline SnO2 thin films are apparent. Figure 16b shows 7452


Chemical Reviews

Review

Figure 16. Typical HRTEM images of SnO2 thin films: (a) an asgrown SnO2 thin film with defects and (b) in situ annealed SnO2 thin film without defects. Reprinted with permission from ref 61. Copyright 2009, Elsevier.

Figure 17. Typical HRTEM images of the interior of SnO2 thin films: (a) the interior of an as-grown SnO2 thin film with defects and (b) the interior of in situ annealed SnO2 thin film with a perfect lattice. Reprinted with permission from ref 61. Copyright 2009, Elsevier.

that the density of defects in the nanocrystalline SnO2 thin films after in situ annealing is much lower. In this figure the arrow indicates the lattice fringes of the thin films. At the interior of the SnO2 nanoparticles the defects are annihilated after annealing. Figure 17a shows a HRTEM image of the interior of an as-grown nanoparticle. The interior of the as-grown nanoparticle is generally cotton-like in shape and covered by a relatively thick amorphous layer (indicated by the arrow). The contrast of the amorphous layer is quite uniform and only the SnO2 crystalline structure is observed. Similar to the microstructure in the bulk of the SnO2 thin films, the SnO2 crystal core at the interior also has a high density of defects. Since the presence of a high density of defects, such as dislocations, is known to accelerate crystal growth, these interior defects may be responsible for the fast growth of SnO2 thin films. In fact, these interior defects in SnO2 nanoparticles may be necessary for the growth of SnO2 thin films. On the other hand, such defects are absent at the interior of the in situ annealed SnO2 nanoparticles as shown in Figure 17b (indicated by the arrow). Since structural defects in the SnO2 particles have an important effect on electron transport, which determines sensitivity to gases, the removal of internal defects by annealing is a major step forward in the development of efficient gas sensors. The Raman spectra of as-grown and in situ annealed SnO2 thin films are shown in Figures 18a and b. Raman spectrum of commercial SnO2 bulk material is also shown for comparison (Figure 18c). There is some asymmetry in the A1g peak of the as-grown SnO2 thin film. The degree of asymmetry in this peak is somewhat reduced in the corresponding peak of in situ annealed SnO2 thin film. Asymmetry of Raman peak is thought to be caused by two factors: nanoscale size and defects.186,187

Figure 18. Raman spectra of (a) as-grown SnO2 thin films; (b) in situ annealed SnO2 thin films, and (c) commercial bulk SnO2 material. Reprinted with permission from ref 61. Copyright 2009, Elsevier.

However, in the present case, in situ annealing does not change the size of the nanoparticles. Thus, the decrease in peak asymmetry is attributed to the decreased density of defects in the in situ annealed SnO2 thin film. The density of defects in in 7453


Chemical Reviews

Review

Figure 19. (a and c) Topographic AFM images and (b and d) corresponding grain-like magnetic domain structures revealed by MFM imaging for pure SnO2 film (a and b) and 5% Gd-doped SnO2 film (c and d). Reprinted with permission from ref 194. Copyright 2012, American Chemical Society.

situ annealed SnO2 thin film is probably comparable to that of commercial bulk SnO2 material. Various structural defects such as Sn interstitial (Sni), oxygen vacancy (VO), and Sn vacancy (VSn) are produced during SnO2 thin films preparation, with possible deleterious effects on their magnetic and optical properties. The Sn vacancy (VSn) and singly ionized oxygen vacancy (VO+) have been reported to possess localized magnetic moments.188−190 The other defects (such as Sni or neutral VO or VO2+) are nonmagnetic.188−190 Rahman et al. showed that VSn can generate a large magnetic moment of around 4 μB using predictions based on density functional theory.189 These authors have attributed the experimentally observed giant magnetic moment of Co-doped SnO2 thin films to the VSn defects.191 In general, it is highly unlikely for a large amount of VSn defects in pure SnO2 thin films to form because of the high formation energy of VSn defects. On the other hand, the low formation energy of oxygen vacancy (VO) facilitates its nucleation and growth during SnO2 thin film deposition.192 Unlike ZnO, there are relatively few published research studies on defect magnetism in SnO2 thin films. This situation is particularly acute on the effect of rareearth element (RE)-ion doping on magnetic properties of SnO2 thin films.193 Novel optical and magnetic properties of these thin films could be produced by doping of rare earth element ions in the SnO2 matrix, which would be interesting from the viewpoint of optical and spintronic applications. Trivalent Gd ions have large localized magnetic moments compared to the transition-metal ions because of their 4f7 electron configuration. It is therefore expected that Gd-doped oxides could exhibit a large magnetic moment. The substitution of trivalent Gd ions at

the tetravalent Sn site could also introduce vacancies in the SnO 2 system. This would be favorable to sustaining ferromagnetism coupling.194 Pure and 5% Gd-doped SnO2 thin films have been examined by AFM and magnetic force microscopy (MFM). Figure 19 shows the AFM and MFM images at a lift height of ∼300 nm over a scan size of 2 μm × 2 μm. The difference between the MFM and AFM images of the films is evident. Coarse grain-like domain structures, with bright and dark contrast of domains corresponding to high concentrations of north and south poles, are manifested in the MFM amplitude images. Pure and 5% Gd-doped SnO2 film exhibit different domain configurations. The former has more prominent and vivid grain-like domain structures (Figure 19b) compared to that of the latter (Figure 19d). The bright grain-like domains in the 5% Gd-doped SnO2 film are relatively fewer in number. They appear to occur randomly and are more widely spaced. The implication of this result is that the observed regular distribution of the bright and dark grain-like magnetic domains in pure SnO2 thin film is an indication of long-range magnetic ordering, as opposed to shorter range ordering in the Gd-doped SnO2 film.195−197 Results of field dependent magnetization M (H) measurements of pure and Gd-doped SnO2 thin films are shown in Figure 20. Pure SnO2 thin film exhibits a clear hysteresis loop. This result suggests that pure SnO2 thin film is ferromagnetic, both at room temperature (Figure 20a) and at low temperature (5 K) (inset (ii) of Figure 20a). A significant amount of coercivity (Hc ≈ 110 and 125 Oe at 300 and 5 K, respectively) is observed in the case of pure SnO2 film, which shows its ferromagnetic nature. The saturation magnetization (Ms) and 7454


Chemical Reviews

Review

Figure 20. Field dependent magnetization (M (H)) loops of (a) pure SnO2 thin film at 300 K, (b) 1% Gd-doped SnO2, (c) ferromagnetic component (MTotal − MPara) of the 1% Gd-doped film after subtracting the paramagnetic component and (d) 2% and 5% Gd-doped SnO2 thin films. Insets show (i) the magnetization of pure SnO2 in low field area at 300 K, (ii) the M − H loop of pure SnO2 at 5 K, (iii) the magnetization of Gddoped films (in units of μB per Gd3+ ion), and (iv) the effective magnetic moment (μeff) of the individual Gd ion in units of Bohr magnetron (μB) versus the Gd concentration in the film. (All data presented have been corrected for the diamagnetic substrate contribution.) Reprinted with permission from ref 194. Copyright 2012, American Chemical Society.

retentivity (Mr) at 300 K are 3.4 × 10−3 and 2.2 × 10−4 emu/ cm3, respectively. At 5 K, the corresponding values of Ms and Mr are 4.6 × 10−3 and 9.8 × 10−4 emu/cm3 respectively. Similar ferromagnetic response has been reported in other oxide systems.198−200 For Gd-doped SnO2 films, a linear M-H behavior (Figure 20b,d) is observed. This suggests the presence of dominant paramagnetic ordering in these films. A curvature is observed near the origin (Figure 20b) in the case of the 1% Gd-doped film. This can be explained in terms of a ferromagnetic component superimposed over a strong paramagnetic background. After subtracting the strong paramagnetic component (MPara) from the total moment (MTotal), the residual component (MTotal − MPara) is plotted in Figure 20c. This shows clear hysteresis-type behavior with a coercivity of Hc ≈ 95 Oe. For the 1% Gd-doped film, the saturation magnetization (Ms) is 1.05 × 10−3 emu/cm3, much smaller than that of pure SnO2 film. Analysis of the above results shows that, with 1 at. % Gd doping in SnO2 thin film, ferromagnetism is suppressed and a large paramagnetic moment is established within the system. For the 2% Gd-doped film, the ferromagnetic component (MTotal − MPara) vanishes completely, while the paramagnetic component is markedly enhanced. Ferromagnetism also vanishes in 5% Gd-doped films. Gd doping has profound effects on ferromagnetism in pure SnO2 thin films. It decreases gradually with increasing Gd doping. Above a certain Gd concentration (between 1 and 2 at. %), ferromagnetism completely vanishes. Early magnetic force microscopy studies show evidence of long-range ferromagnetic ordering in pure SnO2 films but not in 5% Gd-doped film. Gd3+ ions have a localized f-shell paramagnetic moment. Although the insertion of Gd3+ ions

increases the associated magnetic moment by an order of magnitude, there is no intrinsic ferromagnetic ordering in Gddoped SnO2 films. Similar results have been reported by Ney et al.,201 Xu et al.,202 and Barla et al.203 in the case Co-doped ZnO thin films and Ghosh et al.204 in Co-doped SnO2. It has been reported recently that doping of transition metal ions would lead to a loss of magnetization in CeO2 thin films.205 Magnetization per Gd3+ ion and the effective magnetic moment (μeff) of individual Gd ion have been estimated by Ghosh et al. in units of Bohr magnetron (μB). These are shown in the insets (iii) and (iv) of Figure 20. Magnetization per Gd ion and μeff value decrease with increasing Gd concentration. This could be due to enhancement of antiferromagnetic superexchange interaction between Gd3+ ions via oxygen ligands. Increase in Gd concentration in the SnO2 matrix would decrease the average distance between Gd3+ ions, culminating in enhancement of antiferromagnetic superexchange interaction. μeff value of the Gd3+ ion is found to be maximum (2.78 μB) at 1% Gd-doped SnO2 film, considerably less than the full moment of the Gd3+ ion (7.94 μB). Similar results have also been observed in the case of Co:ZnO,206 Co:SnO2,204 and Gd:ZnO.207 In these cases the effective moment of the respective transition metal or rare earth ion is observed to decrease significantly. This could be due to the influence of the ionic environment on the magnetic moment of the ions and the ionic symmetry within the host matrix.208,209 The lattice distortion caused by an increase in Gd concentration could lead to a change in the Gd3+ site symmetry and the effective magnetic moment of the Gd3+ ion. 7455


Chemical Reviews

Review

2.5. Optical Properties of Tin Dioxide Thin Films

laser deposition technique has been successfully applied to the growth of good quality tin oxide thin films. A Sn target in an oxidizing oxygen atmosphere or a SnO2 target can be used in this process.72,147 Microstructural properties of SnO2 thin films deposited on glass substrates by pulse laser deposition have been investigated before and after thermal annealing in the temperature range 50 to 550 °C at 50 °C intervals. XRD results confirm that the SnO2 thin films produced consist of nanoparticles with average grain size in the range of 23.7− 28.9 nm as shown in Figure 8. Optical measurements show that these SnO2 thin films have transmittance greater than 75.3%, making them suitable for window material for solar cell applications. From the experimental optical transmittance and reflectance data, various optical parameters such as optical band gap energy, refractive index and optical conductivity have been calculated in the wavelength range 300−2500 nm. The optical band gap energy and refractive index of SnO2 thin films show approximately linear relationship with the annealing temperature. The optical conductivity of these SnO2 thin films is high. Such results confirm that SnO2 thin films annealed up to 400 °C are good window materials for solar cell applications. In the evaluation of the optical performance of semiconducting oxide thin films the information on optical transmittance is important.219 Optical transmittance spectra of the as-prepared thin film and those after annealing at various temperatures are shown in Figure 22. The annealing temperature varies from 100 to 500 °C at 100 °C intervals. The annealing time is 30 min. All thin films exhibit similar optical transmission behavior in the wavelength range 300−2500 nm,

SnO2 is an important strategic material and has technological applications in diverse areas such as optical and nanoelectronics,116,117,210 gas sensing,211 energy storage and conversion,212,213 catalyst supports,214 etc. The properties of devices such as those for gas sensing and optoelectronic applications are influenced by surface area and grain size of the particles. Improvement of properties of these devices is feasible using nanocrystalline particles with a high surface-to-volume ratio. Therefore, the production of nanoparticles with specific size and microstructure is of crucial significance for such applications. Consequently, synthesis of nanoparticles with precise control of size, shape, and microstructure is one of the most challenging research topics. In particular, the properties of thin films are strongly dependent on their microstructure, composition, crystal defects, and postprocessing operating environment. These issues have aroused much research interest in recent years. Cavaleiro et al. has investigated the effect of annealing temperature on the optical transmittance and band gap of SnO2 thin films.215 The optical transmittance spectra in the wavelength range 250−1000 nm for SnO2 thin films at various annealing temperatures are shown in Figure 21. With increasing

Figure 21. Optical transmittance spectra of SnO2 films with different annealing temperatures. Reprinted with permission from ref 215. Copyright 2012, Elsevier.

annealing temperature, the absorption edge is shifted toward shorter wavelengths and the average transmittance increases. After annealing at 500 °C, the maximum optical transmittance in the visible region improves from 79% to 90%. The increase in transmittance with annealing temperature is attributed to decreasing optical scattering caused by the densification of grains followed by grain growth.216 With increasing annealing temperature, defect density is reduced and the stoichiometry of the thin film is improved. Surface morphology also has an effect on optical properties of these films. Higher transmittance is associated with smoother surface with less grain boundaries.217 Higher annealing temperature leads to better structural homogeneity and crystallinity, resulting in an increase in optical transmittance.218 The use of laser light and the control of composition through in situ doping have enabled the pulse laser deposition technique to produce thin films with reduced contamination.142 The pulse

Figure 22. Optical transmittance spectra of SnO2 thin films annealed at different temperatures for 30 min. (a) as-prepared thin film; (b) 100, (c) 200, (d) 300, (e) 400, and (f) 500 °C. Reprinted with permission from ref 219. Copyright 2011, Elsevier. 7456


Chemical Reviews

Review

temperature (Figure 24a). Because of the temperature dependence of electron phonon interactions, annealing may

which almost covers the entire solar energy spectrum. However, in the wavelength range 300−750 nm, the values of optical transmittance for the as-prepared thin film and the thin films annealed up to 400 °C are larger (>75.3%) than that for the thin film annealed at 500 °C (1000. In view of this, they are often referred to as one-dimensional materials. Nanowires have many novel properties that bulk materials do not possess. The narrow width of nanowires leads to quatum confinement of the electrons in the lateral direction. Consequently, electron energy levels in nanowires are different from those in bulk materials.235−238 In some nanowires, quantum confinement in the lateral direction leads to discrete values of the electrical conductance. Such discrete values of the electrical conductance arise because the number of electrons that can travel through the wire at the nanometer scale is restricted due to this quantum confinement.239−243 The number of potential applications for nanowires is numerous. Examples include components in electronic, optoelectronic and nanoelectromechanical devices, additives in advanced composites, metallic interconnects in nanoscale quantum devices, field-emitters and leads for biomolecular nanosensors.244−248 In view of their novel properties and potential applications, one-dimensinal nanostructures have attracted increasing interests among materials scientists in recent years.249−254 Nanostructured materials include nanoparticles, nanowires, nanorods, nanotubes, nanoribbons (or nanobelts), and nanorings. These are candidate materials for the construction of nanoscaled electronic and optoelectronic devices, gas sensors, catalysts, and thin films.72,255−259 Among the various semiconducting nanowires, SnO2 nanowire has the unique characteristics of a wide-band gap (Eg = 3.6 eV at 300 K) and large surface-to-volume ratio. It has been extensively studied for applications as lithium-ion batteries, varistors, gas sensors, and transparent conducting electrodes.71,260−262 Control of microstructure is a critical step in the development of methods of synthesis of nanostructures for functional oxides for scientific and technological applications.263−267 The synthesis of metal oxide based nanostructures has been achieved by various workers.74,172,268−270 Early techniques of synthesis of one-dimensional nanomaterials include the use of photolithography and scanning tunneling microscopy.271,272 However, nanomaterials produced by these techniques are not suitable for industrial applications. Two established mechanisms by which materials can grow one-dimensionally into wires, whiskers, or rods involve screw-dislocations and vapor− liquid−solid (VLS).273,274 SnO2 nanowires will need to be produced on a large scale with relative ease for industrial applications. The growth behavior of SnO2 nanowires prepared by PLD method is not yet fully understood. In the following sections, the micro/nanostructures and properties of SnO2 nanowires will be described. Research efforts aimed at understanding the relationship between synthesis strategies and micro/nanostructures and properties of SnO2 nanowires will be reviewed.

Figure 30. XRD patterns of the as-prepared SnO2 nanowires prepared by pulsed laser deposition method. Reprinted with permission from ref 275. Copyright 2009, Elsevier.

3.2. Preparation and Characterization of Tin Dioxide Nanowires

the diffraction peaks are in nearly perfect agreement with those of the rutile SnO2 structure in terms of peak positions and relative intensities. Analysis of the XRD data reveals SnO2 lattice constants of a = 4.737 Å and c = 3.185 Å for as-prepared nanowires. These are consistent with those of bulk rutile SnO2 tetragonal structure.274 Thus, the crystallographic planes identified on the XRD patterns of SnO2 nanowires are, in terms of peak positions and relative peak intensities, typical of pure crystalline SnO2 phase. The absence of XRD peaks

The details of fabrication procedure for SnO2 nanowires are described below.127 A sintered cassiterite SnO2 bulk is used as the target. Cleaning of the target with methanol in an ultrasonic bath before installation is essential in order to minimize contamination. Typically, a KrF excimer laser producing pulse energies of 150 mJ at a wavelength of 248 nm could be used for pulse laser ablation. The duration of each excimer laser pulse is 7460


Chemical Reviews

Review

lattice parameters of a = 4.737 Å and c = 3.185 Å, in agreement with the XRD results. The single rutile SnO2 structure of the nanowire is also confirmed by SAED. HRTEM image of the single SnO2 nanowire of Figure 31b is shown in Figure 31c. The growth direction of the nanowire is confirmed to be [110] (indicated by an arrow parallel to the long axis of the nanowire). The nanowire is structurally uniform, and the clear lattice fringes suggest that the nanowire is a single crystal. The interplanar spacing is about 0.334 nm corresponding to the {110} planes of rutile crystalline SnO2. The above results reveal that growth direction of SnO2 nanowires is along the {110} planes, and that the nanowires are single pure crystals of SnO2 rutile phase. The crystalline nanowires are deposited at ambient temperature at a pressure of about 3 × 10−4 mbar. The mean free path of the nanoclusters is about 0.3 m, approximately equal to the dimensions of the deposition chamber. Therefore, the targetsubstrate distance and the power density on the target are critical for this process. Since no catalytic droplets have been observed on any end of SnO2 nanowires, the vapor−liquid− solid mechanism is not supported by experimental observations. It is probable that a vapor−solid growth process is responsible for the SnO2 nanowire’s formation process. The SnO2 target is etched by the high-temperature laser plumes during pulse laser ablation. The SnO2 vapor originating from the starting material in the high temperature zone directly deposits at the low temperature region. Due to the absence of a carrier gas in the chamber, the SnO2 vapor deposits in the form of nanoclusters through aggregation of SnO2 molecules in the deposition chamber near the target. As the pulse laser ablation process continues, more SnO2 vapor is generated and more SnO2 nanoclusters are formed. The as-formed SnO2 nanoclusters are energetically favorable sites for rapid adhesion of additional SnO2 molecules. This results in the formation of SnO2 nanowires. The current/voltage relationship for a SnO2 based field effect transistor device is shown in Figure 32. The gate voltage ranges from −1 to +5 V and the drain voltage ranges from 0 to 1 V. Increasing the back gate voltage would lead to higher conductance, indicating n-type semiconductor characteristics

corresponding to impurities, such as other forms of tin oxides, indicates that the nanowires are pure phase of rutile SnO2. The TEM image showing the morphology of as-prepared nanowires is shown in Figure 31a. The SnO2 crystals exhibit

Figure 31. (a) Typical TEM bright-field image of SnO2 nanowires. (b) TEM bright-field image of a single SnO2 nanowire; the inset at the upper left-hand corner shows the SAED patterns of the single SnO2 nanowire; (c) the HRTEM image of the single SnO2 nanowire shown in panel b. Reprinted with permission from ref 275. Copyright 2009, Elsevier.

wire-like shapes with diameters of 10−30 nm and lengths from several hundreds nanometers to a few micrometers. No spherical droplets were observed at the top ends of these nanowires. This observation is inconsistent with the vapor− liquid−solid or solution-liquid−solid mechanism proposed for nanowire growth by a catalytic-assisted process. In this process a metal liquid droplet forms at the growth front of the wire and acts as the catalytic active site. The absence of such droplets fails to confirm this theory. A typical TEM bright-field image of a single SnO2 nanowire is shown in Figure 31b. The nanowire is about 15 nm wide and exceeds 380 nm in length. It is very smooth, straight and uniform, with almost structurally perfect geometrical shape. The selective area electron diffraction pattern of a single nanowire is shown at the inset at the upper left-hand corner of Figure 31b. The selective area electron diffraction pattern, taken along the [1̅11] direction, can be indexed using the d-spacings of the (110) and (211) crystal planes. It is consistent with a tetragonal unit cell with

Figure 32. Current−voltage (IDS - VDS) curves of field-effect transistor devices for the back-gated single SnO2 nanowire with VGS = +5 to −1 V in −2 V steps from top to bottom. Reprinted with permission from ref 275. Copyright 2009, Elsevier. 7461


Chemical Reviews

Review

indicate that pulse laser ablation technique may be used to manipulate other one-dimension nanostructural materials. This may provide new opportunities on nanoscale electronic and optoelectronic device applications.

of the SnO2 nanowires. This might be due to oxygen vacancies and extra tin interstitial atoms in the lattice during the pulsed laser ablation process. At VG from −1 to +5 V with VDS = 15 mV, an on/off current ratio as high as 100 000 has been achieved. Mobility and subthreshold voltages are determined using ideal field effect transistor equations. The total charge on the nanowire is Q = CVg,t, where C is the nanowire capacitance and Vg,t is the threshold voltage necessary to completely deplete the wire. The capacitance of the nanowire is calculated from the experimental results using the equation249,276 2πεε0L C= 2h ln r (1)

3.3. Microstructure and Growth of Tin Dioxide Nanowires

Vertically grown, highly aligned, single crystalline SnO2 nanowires have potential applications as vertical surroundgate FET, light-emitting diodes, highly efficient dye-sensitized solar cells and nanogenerators.279−283 Uniquely aligned or oriented SnO2 nanowires are expected to have novel optical and electrical properties compared with those of epitaxial films. Hong et al. reported that highly uniform, well-aligned, single crystal SnO2 nanowires can be grown on TiO2 (101) substrates by combining carbothermal reduction with Au-catalyzed heteroepitaxial growth.284 The crystallographic growth directions were analyzed in detail by these authors. A low magnification TEM bright-field image of an individual SnO2 nanowire is shown in Figure 34a. The nanowire shown has a

( )

where ε = 3.9 is the relative dielectric constant, h = 500 nm is the thickness of the silicon dioxide layer, r = 15 nm is the nanowire radius and L = 1.5 μm is the channel length. The IDS − VGS curve (in log scale) of the FET device constructed from the single SnO2 nanowire at VDS = 15 mV is shown in Figure 33. The subthreshold slope is independent of the bias voltage

Figure 33. IDS − VGS curve of field-effect transistor device of the single SnO2 nanowire in log scale at VDS =15 mV. Reprinted with permission from ref 275. Copyright 2009, Elsevier.

and is estimated to be about 0.12 V/decade. An on/off current ratio as high as 105 was achieved when VDS = 15 mV. This value is consistent with those obtained in planar single-crystalline SnO2 nanowires and other mental oxide nanowires.277,278 The one-dimensional electron density is n = Q/eL ≈ 3.86 × 105 cm−1. Electron mobility can also be deduced from the transconductance of the field effect transistor. In the linear region, the trans-conductance gm is

gm =

dIDS dVDS

Figure 34. (a) Low magnification TEM image of a SnO2 nanowire. (b) SAED pattern of the SnO2 nanowire in panel a, and (c) HRTEM image of the SnO2 nanowire in panel a. Reprinted with permission from ref 284. Copyright 2010, American Chemical Society.

(2)

and ⎛C ⎞ gm = μe ⎜ 2 ⎟VDS ⎝L ⎠

constant diameter and a straight, smooth surface. The average diameter of nanowires similar to that shown in Figure 34a is about 22 nm with a relatively narrow distribution that increases slightly with increasing growth time. The SAED pattern depicted in Figure 34b shows that the nanowire is single crystal rutile SnO2 with a zone axis of [010]. Figure 34c shows a HRTEM image of SnO2 nanowire. Clear lattice fringes are manifested and there are no obvious defects or dislocations, providing further evidence of its high single crystallinity. The

(3)

The data presented in Figure 33 show that the mobility of channel stoichiometry of the field effect transistor is about 191 cm2 V−1 s−1. Field effect transistor devices constructed from SnO2 nanowires have better electrical properties for the control of the on/off ratio and threshold voltage. These findings 7462


Chemical Reviews

Review

inset of Figure 34c shows the fast Fourier transformation (FFT) pattern taken from the image marked by the box in Figure 34c. The fast Fourier transform pattern could be completely indexed to the rutile structure when viewed along the [010] zone axis. The results show that the sidewall of the nanowire is straight and sharp on the atomic scale. It is parallel to the (1̅01) plane. The fast Fourier transform pattern in the inset of Figure 34c shows only the {101} and {200} families of planes. The corresponding interplanar spacings are 0.259 and 0.347 nm, respectively. These are in good agreement with the (101) and (200) interplanar spacings of rutile SnO2. In a tetragonal structure with lattice constants (a = b = 4.738 Å, c = 3.187 Å), the (101) plane and [101] direction are not perpendicular but are inclined at an angle of 68° with each other. The growth direction of nanowire, angled 68° to the (1̅01) plane, is [101]. Thus, the SnO2 nanowires grow along the [101] direction at an angle 68° to the (101) plane. The structural characteristics of the SnO2 nanowires are similar to those of SnO2 nanobelts synthesized by the vapor−solid method, where the growth direction, top surface planes, and side surface planes are [101], ± (101̅), and ± (010), respectively.285 Figure 35 shows a HRTEM picture of the nanowire/ substrate interfacial region. The (101) plane of the SnO2 nanowire and the (101) plane of the titanium dioxide substrate are parallel. The nanowire/substrate interface is clean and flat on the atomic scale under high magnification (Figure 35b). Both (1̅01) lattice fringes are connected at the interface with an approximately 2° tilt resulting from the lattice mismatch between SnO2 and titanium dioxide. Figure 35c shows a crystal structure model schematic diagram of the tilting of the (1̅01) lattice fringes at the interface. The SnO2 (101) nanowires are heteroepitaxially grown on the titanium dioxide (101) substrate without a distinct buffer layer. In Figure 35a, the (1̅01) lattice fringes in the broadened base of the nanowire is enclosed by the circle. The fast Fourier transform pattern of this area exactly matches that of the nanowire. This indicates that this region has the same crystallographic orientation as the nanowire. Planar defects could be observed between broadened base and the nanowire at some interfaces. However, the growth direction and epitaxial relationships are still conserved. In some other situations, the interface could not be clearly identified due to overlapping with neighboring particles. The larger diameter of the nanowire base could be caused by surface tension acting on the droplet during VLS growth or contact between two nanowires culminating in growth termination of one of the nanowires.286 While there have been many reports on the controlled synthesis of metal oxide nanowires,287 the detailed growth mechanism of SnO2 nanowires is still not fully understood. The physical properties of nanowires such as thermal and electrical conductivity, refraction index, piezoelectric polarization, and band gap energy can be modulated by tuning the growth direction of nanowires.288 Therefore, control of the growth direction of nanowires is highly desirable. The fabrication of single crystalline SnO2 nanowire arrays have been reported.284,289,290 However, there are no in-depth studies to fully elucidate their growth mechanism. Leonardy et al. synthesized aligned SnO2 nanowires on sapphire substrates (100) and (110)290 and Kim et al. studied the orientation relationship and interface structure of SnO2 nanowires on titanium dioxide (101) substrate.284 In these studies, a reasonable growth mechanism of ordered SnO2 nanowires

Figure 35. (a) HRTEM image of a SnO2 nanowire on TiO2 (101) substrate. Incident beam is in the [010] direction. (b) High magnification image of the interfacial area marked by the box in panel a. (c) Schematic representation of the atomic configuration of the SnO2/TiO2 interface viewed along the [010] direction of TiO2. Reprinted with permission from ref 284. Copyright 2010, American Chemical Society.

has not been proposed. In view of the technological importance of SnO2 nanowires, for instance, for use as a prototype material in gas sensing applications, the growth mechanism of SnO2 nanowires with controlled orientation and surface activity is of significant interest.291 Distinct morphologies are exhibited by SnO2 nanowires grown on polycrystalline Al2O3 and single-crystalline titanium dioxide (001) substrates. Figure 36a shows an oriented meshlike network of SnO2 nanowires on titanium dioxide (001). However, on Al2O3 substrates the morphology of SnO2 nanowires is very different (Figure 36b). The SnO2 nanowires are very long (about 50 μm) and disorderly entangled. SAED reveals that the particle at the nanowire tip is gold, indicating a VLS type of growth pattern (Figure 36c).274 The different growth directions of SnO2 nanowires may be caused by the influence of crystallographic orientation of substrates and coherent grain boundaries at the liquid−solid growth front. The control of the diameter of SnO2 nanowires by adjusting the size of gold catalyst has been reported.292 The oriented SnO2 nanowires on titanium dioxide substrate have narrower diameter distribution in comparison to that of random SnO2 nanowires on Al2O3 substrate. This is probably due to the heteroepitaxy and lowering of the nucleation barrier for SnO2 crystallization by gold particles on titanium dioxide substrates. 7463


Chemical Reviews

Review

Figure 36. SEM images of SnO2 nanowires grown on (a) TiO2 (001) and (b) polycrystalline Al2O3 substrates. (c) VLS growth mechanism: influence of substrate. Reprinted with permission from ref 291. Copyright 2011, American Chemical Society.

Al2O3 differs from titanium dioxide as a substrate material in that the former has complex surface morphologies and random grain orientation. Thus, polycrystalline Al2O3 substrate provides multiple sites for different nucleation rates, culminating in wider size distribution. The tilt angles of the nanowires from the surface of the titanium dioxide (001) substrate could be determined from SEM images taken at different electron incident angles to measure the projected angles (Figure 37a). The characteristic angular orientation of the nanowires could also be determined from TEM images (Figure 37b). In the case shown in Figure 37, the tilt angles are in the range 33 ± 2° determined by both methods. The various electron incident angle views show that the orientations of the SnO2 nanowires grown on titanium dioxide (001) substrate are quite uniform.293 Because of the mirror images of the unit cell and 4-fold axes of symmetry along the c-axis of the titanium dioxide (001) substrate, there are four equivalent growth directions: [011], [011̅ ], [101], and [10̅ 1]. Epitaxial growth of SnO2 nanowires on titanium dioxide (001) substrate has been observed by HRTEM. The observations of epitaxial growth and single crystalline nature of SnO2 nanowires grown on titanium dioxide (001) substrate are also confirmed by fast Fourier transform patterns. Figure 38 shows a focused ion beam assisted SEM tomographic study of ordered SnO2 nanowires on titanium dioxide (001) and (101) samples. This method has the advantage over conventional electron microscopy in that, instead of a two-dimensional projection of three-dimensional objects, it can provide a three-dimensional picture of the nanostructures grown on a flat surface.294 The methodology is based on well established procedures in thin film sectioning and computer-aided combination of images commonly used in the semiconductor industry for testing silicon based microchips.295 The sample preparation procedure for focused ion beam tomographic analysis involves serial cuts through the nanowires (Figures 38a and b). The focused ion beam scanning electron microcopy tomographic images of ordered SnO2 nanowires on titanium dioxide (001) and (101) substrates are shown in Figures 38c and d. For ordered SnO2 nanowires grown on titanium dioxide (101) substrate there is a preferred single orientation direction [101] with a tilt angle of about 69 degrees.

Figure 37. (a) SEM images of ordered SnO2 nanowires on titanium dioxide (001) substrate recorded at different incident angles of electron beams. Fast Fourier transform patterns are obtained from the corresponding SEM images. (b) Cross-sectional TEM image of ordered SnO2 nanowires grown on titanium dioxide (001) substrate. The short nanowires are obtained by short deposition time. Reprinted with permission from ref 291. Copyright 2011, American Chemical Society.

On the other hand, for ordered SnO2 nanowires grown on titanium dioxide (001) substrate there are four equally probable directions: [101], [1̅01], [101̅], and [1̅01̅]. Since the gold catalyst on titanium dioxide [101] substrate meets the criteria of the incubation growth model, tilting of liquefied catalyst droplets and incubation time in the SnO2/titanium dioxide (001) sample have not been observed. Epitaxial growth of SnO2 nanowires on titanium dioxide (101) substrate has recently been confirmed by detailed structural analysis performed by Kim et al.284 The incubation model can explain why the [101] direction has dominated the growth of SnO2 nanowires even though the growth directions [101], [011], and [011̅] are also crystallographically equivalent. 3.4. Defect and Properties of Tin Dioxide Nanowires

The lattice mismatch at the interface between the growth product and substrate would introduce various types and degrees of strains which could have a profound effect on its growth behavior.296,297 This could result in different densities of surface or defect states.298,299 Various reports have shown that surface and defect states have critical effects on luminescence properties, although the relationship between the density of the defect states and luminescence properties is still not fully understood.300,301 This surface and defect related effect on optical properties of SnO2 nanowires has not been fully investigated and is not yet well understood. An improved understanding of the electronic structure of SnO2 nanostructures is necessary if further progress on this topic is to be made. The electronic structure of SnO2 nanowires could be investigated using synchrotron radiation to perform X-ray 7464


Chemical Reviews

Review

Figure 38. (a and b) Sample preparation (serial cuts through the nanowires) for focused ion beam tomographic analysis. Focused ion beam tomographic images of ordered SnO2 nanowires on (c) titanium dioxide (001) and (d) titanium dioxide (101) substrates. Reprinted with permission from ref 291. Copyright 2011, American Chemical Society.

octahedral tip structure as those grown on stainless steel substrate. The only difference is that they are less uniformly distributed (Figure 39c,d). The X-ray absorption near edge structure spectra for SnO2 nanowires grown on different substrates are shown in Figure 40. The total electron yield (TEY), fluorescence yield (FLY), and photoluminescence yield (PLY) are plotted as a function of photon energy.302 Two sets of triplet peaks appear at about 490 and 498 eV for the Sn M5,4 edge X-ray absorption near edge structure spectra in Figure 40a,c. These peaks correspond to the 3d5/2 and 3d3/2 to 5p transitions. They are characteristic features of rutile SnO2.300 Similar behavior is exhibited by the Sn M5,4 edge tracked by the three yields (TEY, FLY, and PLY). Positive edge jumps appear in all three cases. The inset in Figure 40a shows that, for SnO2 grown on stainless steel substrate, there is a pre-edge resonance at 486.5 eV for the total electron yield (TEY) and fluorescence yield (FLY) spectra just before the Sn M5,4 edge. This is attributed to surface states caused by unsaturated coordination of surface tin ions due to oxygen vacancies.301 The third peak of the first triplet (M5) is very intense. This is caused by contribution from the surface state of the M4 edge underneath, providing further support of the influence of surface states. The surface sensitive total electron yield, bulk sensitive fluorescence yield and optically sensitive photoluminescence yield of the tin M edge are similar for copper substrate. There is no pre-edge resonance peak at 486.5 eV before the tin M-edge for SnO2 grown on copper substrate. The third peak of the first triplet (M5) is less intense compared to that of SnO2 grown on stainless steel substrate. This indicates that defect/surface states are less numerous for SnO2 grown on copper substrates. Figure 40a shows the preedge and the very intense third peak of the first triplet (M5) in photoluminescence yield. This implies that selective excitation of tin associated with defects has a profound effect on luminescence yield. For SnO2 grown on copper substrate, the photoluminescence yield spectrum at resonance (Figure 40b)

absorption near edge structure (XANES) and X-ray excited optical luminescence (XEOL) measurements. This approach has been used by Wang et al. on SnO2 nanowires prepared on different substrates.302 As synthesized SnO2 nanowire arrays vertically grown on stainless steel substrate have been examined by these authors using SEM. The top and side views are shown in Figure 39a,b. The nanowires have diameters ranging from 100 to 150 nm, and lengths of about 500 nm with octahedral tips. They are uniformly distributed on the substrate. SnO2 nanowires grown on copper substrate have similar shape and

Figure 39. SEM images of SnO2 nanowires growing on different substrates: stainless steel substrate, (a) top view and (b) side view; copper substrate, (c) top view and (d) side view. Reprinted with permission from ref 302. Copyright 2012, American Chemical Society. 7465


Chemical Reviews

Review

Figure 40. O K-edge and Sn M5,4 edge XANES spectra of SnO2 nanowires grown on stainless steel (a) Sn M5,4 edge and (b) O K-edge, on copper (c) Sn M5,4 edge, and (d) O K-edge. Reprinted with permission from ref 302. Copyright 2012, American Chemical Society.

Figure 41. X-ray excited optical luminescence of SnO2 nanowires on different substrates: (a and b) stainless steel; (c and d) copper. Reprinted with permission from ref 302. Copyright 2012, American Chemical Society. 7466


Chemical Reviews

Review

Figure 42. SEM micrographs of (a) as-deposited (TO) and plasma-treated (TTO) SnO2 nanowire samples at (b) 10, (c) 40, and (d) 80 W. Reprinted with permission from ref 311. Copyright 2010, American Chemical Society.

metal−semiconductor interfaces and modulating the crystal fields near the interfaces.304,305 Partial reduction of Sn(IV) species could be achieved by using capacitively coupled rf plasma to modify tin oxide films after their growth. This discovery was made by Mathur et al.292,306 The capacitively coupled rf plasma creates a valence dynamics (Sn2+, Sn3+, and Sn4+) on the surface, which substantially enhances their ethanol-sensing properties. Preferential etching of bridging oxygen species in the lattice produces a partial reduction of SnO2 deposits. This increases the sensor performance through higher sensitivity with lower response and recovery times, accompanied by a reduction in the operating temperature. Surface chemical composition could be changed homogeously by plasma assisted modification of the nanowires. This could be achieved without losing the single-crystalline nature and the characteristic high aspect ratio of the nanowires. Mathur et al. believe that the sensitivity of the tin oxide nanowire-based gas sensor can be enhanced by controlling the gas species and power for plasma treatment.292,306 Surface etching is stable over several cycles indicating an equilibrium state driven by redox reactions. While there have been many studies on plasmatreated tin oxide films307,308 and nanorods,309,310 little is known on the plasma-treatment of SnO2 nanowires. Although plasma treatment has been used to fabricate one-dimensional nanostructures or as a surface modification method, the mechanistic aspects of plasma treatment and their correlation to enhanced gas-sensing performance have not been fully elucidated. SEM pictures of as deposited and plasma treated SnO2 nanowires are shown in Figure 42. The nanowires exhibit uniform one-dimensional structure with diameters ranging from

shows a less intense peak due to lower densities of surface or defect states. The luminescence properties of SnO2 nanowires with different densities of surface or defects states have been investigated by X-ray excited optical luminescence performed at room temperature.302 In Figure 41, each sample has been excited with soft X-ray photons at both below and above the threshold of the O K-edge and Sn M5,4-edge. This is equivalent to X-ray excited optical luminescence obtained by using selected excitation energy in the X-ray absorption near edge structure photoluminescence yield. There is an intense emission band in the range 400 to 900 nm (3.1−1.4 eV) for all four spectra. They are centered at 498 nm (yellow-green luminescence), in agreement with X-ray excited optical luminescence studies of SnO2 nanoribbons originating from intrinsic surface states.300,301 Very weak emission has been observed originating from the band gap of SnO2 (3.6 eV). This would produce a peak at 344 nm. The weak emission at 344 nm suggests that near band gap emission is quenched by the surface or defect states in the band gap of SnO2. A weak band gap feature is not revealed by a log plot of the X-ray excited optical luminescence data. For stainless steel substrate, the luminescence intensity increases across the Sn M5,4 edge but that the intensity decreases across the O K-edge. 3.5. Modified Tin Dioxide Nanowires and Composites

Due to their marginal transducing property and high operation temperatures, stoichiometric tin oxide nanowires do not satisfy the requirements for gas detection.303 Doping the pure base material with metals and/or metal oxides may overcome these inherent limitations. Tunable physical and chemical properties can be produced, for instance, by creating Schottky barriers at 7467


Chemical Reviews

Review

20 to 80 nm. At plasma energy of 10 W, there is no substantial modification of the nanowire morphology. At higher energies (>20 W), there is surface erosion accompanied by enhanced surface roughness (Figure 42c,d) in the samples. HRTEM analysis shows a uniform and well-ordered surface structure (Figure 43a) for the untreated as-deposited (TO) sample. For

Figure 44. HRTEM images of plasma-treated SnO2 nanowires after (a and b) short and (c and d) long e-beam radiation showing the e-beam influence on the plasma-treated SnO2 nanowires. Reprinted with permission from ref 311. Copyright 2010, American Chemical Society.

increase electron-transport capability and improve fieldemission properties. Photoluminescence (PL) is also one of the most fascinating properties of SnO2 nanomaterials which have many applications in optoelectronic devices, such as UVlight emitting diodes or laser diodes.333,334 However, while the photoluminescence properties of SnO2 nanomaterials have been investigated, the corresponding properites in grapheneSnO2 composite nanostructures have not been as extensively studied. Field emission SEM has been employed to examine SnO2 and graphene-SnO2 nanostructures.335 The results of one such study are shown in Figure 45. The sputtering times of tin nanoparticles to produce these nanostructures are 5 and 10 min. The corresponding low magnification pictures are shown in the insets. The high magnification pictures show a “spherical flower” appearance with each spherical flower comprising a number of nanorods protruding radially from the center. The SnO2 and graphene-SnO2 nanostructures have somewhat different morphorlogies, with the latter being more compact and uniform. Such differences in morphology are caused by the influence of the graphene buffer layer during the growth of SnO2. This shows that the graphene buffer layer has an effect on the properties of the graphene-SnO2 composite. The field emission current density vs field strength data on SnO2 and graphene-SnO2 composite nanostructures are shown in Figure 46. The sputtering times for tin nanoparticles are 5 and 10 min. For a sputtering time of 5 min, Eto of SnO2 nanostructure and graphene-SnO2 composite nanostructure are 7.21 V/μm and 3.86 V/μm respectively, and the corresponding Ethr are 15.9 and 8.6 V/μm, respectively. For a sputtering time of 10 min, Eto of SnO2 nanostructure and graphene-SnO2 composite nanostructure are 10.0 and 5.39 V/μm respectively, and the corresponding Ethr are 19.7 and 10.2 V/μm, respectively. The data show that field emission properties of graphene-SnO2 composite nanostructure are better than those

Figure 43. HRTEM images of (a) as-deposited (TO) and (b) plasmatreated (40 W, TTO4) SnO2 nanowires. Reprinted with permission from ref 311. Copyright 2010, American Chemical Society.

sample TTO4 plasma treated at 40 W, enhanced disorder and depletion of material due to the influence of plasma treatment and bombardment of energetic ions is evident (Figure 43b). The TEM picture shown in the inset of Figure 43b reveals that sample TTO4 has one-dimensional heterostructure with a crystalline SnO2 core and an amorphous outer layer. This structural change appears to be caused by energetic ions impinging on the nanowire surface. The rough amorphous outerlayer in the plasma-treated sample increases the density of dangling bonds and free lattice sites, thereby favoring surface adsorption phenomena. The existence of an amorphous outerlayer is confirmed by high resolution TEM (Figure 44b). Fast Fourier transform of the outerlayers shows the amorphous feature. After short (Figure 44a,b) and long (Figure 44c,d) e-beam radiation, HRTEM shows that crystallization and ordering processes begin to appear on the surface to form crystalline nanoclusters.311 Researchers over the world have used various methods to prepare graphene-metal oxide nanocomposites such as graphene-CuO, 312−314 graphene-Co 3 O 4 , 315 grapheneZnO, 3 1 6 − 3 1 8 graphene-TiO 2 , 3 1 9 − 3 2 3 and grapheneSnO2.324−332 However, there is a scarcity of data on the effect of graphene buffer layers on field-emission properties of SnO2. SnO2 composite nanostructures with graphene buffer layer may 7468


Chemical Reviews

Review

Figure 45. Field emission SEM images of SnO2 nanostructrues when sputtering time of tin nanoparticles is (a) 10 and (c) 5 min and graphene-SnO2 composite nanostructrues when sputtering time of tin nanoparticles is (b) 10 and (d) 5 min. The insets show the corresponding low magnification images. Reprinted with permission from ref 335. Copyright 2011, American Chemical Society.

in SnO2 layer thickness, the field emission properties of the graphene-SnO2 composite nanostructure are also improved. Photoluminescence spectra of SnO2 and graphene-SnO2 nanostructures are shown in Figure 47. The sputtering times

Figure 46. Field emission current density-field strength characteristics of the SnO2 nanostructrues when sputtering times of tin nanoparticles are (a) 10 and 5 min (inset) and graphene-SnO2 composite nanostructrues when sputtering times of tin nanoparticles are (b) 10 and (c) 5 min. Reprinted with permission from ref 335. Copyright 2011, American Chemical Society.

Figure 47. Photoluminescence spectra of the SnO2 nanostructrue when sputtering times of tin nanoparticles are (a) 10 and (c) 5 min and graphene-SnO2 composite nanostructrue when sputtering times of tin nanoparticles are (b) 10 and (d) 5 min. Reprinted with permission from ref 335. Copyright 2011, American Chemical Society.

of pure SnO2 nanostructure. A typical metal−semiconductor ohmic contact without a contact barrier between zinc oxide and graphene has recently been reported by Hwang et al.336 In their study, the graphene is deposited on silicon substrate to act as a buffer layer in the zinc oxide/graphene system similar to that in the tin dioxde/graphene system described above. The graphene-SnO2 composite nanostruture has improved field emission properties due to its high electron-transport and low contact barrier. Moreover, for the graphene-SnO2 composite nanostructure, when the sputtering time is decreased from 10 to 5 min, Eto decreases from 5.39 to 3.86 V/μm and Ethr decreases from 10.2 V/μm to 8.6 V/μm. Decreasing the sputtering time would decrease the thickness of the SnO2 layer. Since electron transport would be enhanced by this reduction

of tin nanoparticles are 5 and 10 min. There are six peaks located at 403, 422, 447, 485, 527, and 600 nm. The positions of these peaks are not changed by the deposition of the graphene buffer layer or the variation in sputtering time for tin nanoparticles. The band to band emission peaks of the SnO2 and graphene-SnO2 nanostructures are not present. This is because the energy gap of bulk SnO2 is 3.62 eV, outside the limits of the photoluminescence detection range.337 Similar result on SnO2 nanograss has been reported.338 7469


Chemical Reviews

Review

cells.223−227 Synthesis of one-dimensional SnO2 nanorods is more difficult than that of two-dimensional SnO2 thin films and zero dimensional SnO2 particles. Well established methods such as sputtering can be used to synthesize SnO2 thin films, and sol gel technique can be used to synthesize SnO2 nanoparticles. On the other hand, synthesis of one-dimensional nanostructures such as SnO2 nanorods, nanowires, nanotubes, and nanobelts requires development of novel methods. Research on the synthesis techniques of one-dimensional nanostructures such as SnO2 nanorods involves addressing some basic issues about dimensionality and space confined transport phenomena.357 There is, however, significant incentives to develop these techniques as one-dimensional nanostructured systems possess novel properties different from those of the bulk crystals with potential applications in numerous areas such as nanoscale electronics and photonics.268,358−360 The synthesis and growth, doping and modification of SnO2 nanorods and nanorods planted graphites will be described in the following sections. The relationship between synthesis strategies, micro/nanostrutures and properties will be reviewed.

4. TIN DIOXIDE NANORODS 4.1. Overview on Tin Dioxide Nanorods

Of the various nanoscale objects encompassed by the field of nanotechnology, nanorods possess a specific morphology with an aspect ratio of about 3 to 5. Nanorods are produced by direct chemical synthesis from metals or semiconducting materials. Synthesis procedure usually involves a combination of ligands acting as shape control agents which bond to different faces of the nanorod seed with different strengths. As a result, different faces of the nanorod seed grow at different rates, culminating in an elongated shape.339−345 Nanorod is a one-dimensional nanostructure. This family also includes nanowires, nanobelts, nanoribbons, nanoneedles, and nanotubes. They have attracted considerable research interest due to their unique and fascinating properties as well as their potential technological applications.346 In particular, some one-dimensional nanostructures have high luminescence efficiency. They are emerging as potentially powerful building blocks for nanoscale photonic devices such as light-emitting diodes, photodiodes, lasers, active waveguides, and integrated electrooptic modulator structures.347,348 It is reported that the creation of core/shell heterostructures in these one-dimensional nano objects could further improve their luminescence properties.349,350 For example, the ultraviolet photoluminescence properties of one-dimensional zinc oxide could be enhanced by adding an outer sheath of a larger band gap material such as zinc sulfide,351 SnO2,352 magnesium oxide,353−355 or aluminum oxide.356 The emission intensity could be increased by a few times due to the quantum confinement effect of the photogenerated carriers in the zinc oxide core. The reflectivity of nanorods can be altered by changing their orientation with an applied electric field. This special property offers significant potential for applications in display technologies. In view of the small size of nanorods, direct miniaturization of micro/nanodevices may be achieved, facilitating another potential application in microelectromechanical systems. There are extensive research interests on the influence of nanorod dimensions, surface and interface morphology on the materials’ chemical and physical properties. Nanorods, similar to other noble metal nanoparticles, can also be used as theragnostic agents. Nanorods absorb photons in the near IR, and generate heat when excited with IR radiation. This special property of nanorods has led to their use as cancer therapeutics. Nanoscale objects in the size range 20 to 300 nm tend to preferentially accumulate in tumor cells. This is because tumor cells need to develop new blood vessels to fuel their growth. These new blood vessels are different from regular blood vessels in terms of poor lymphatic drainage and disorganized, leaky vasculature. Thus, when nanoscale objects such as nanorods are injected into the patient’s bloodstream, they tend to accumulate at the tumor cells. When the patient is exposed to IR radiation of the appropriate wavelength, the IR photons will pass through body tissue and activate the nanorods selectively taken-up by the tumor cells. The activated nanorods will generate heat locally at the tumor cells, destroying the cancerous tissue while leaving healthy cells unharmed. Nanorods based semiconducting materials have potential applications as energy harvesting and light emitting devices. SnO2 nanorods, in particular, have attracted substantial research interest in view of their potential applications in transparent electrodes, far-IR detectors, and high-efficiency solar

4.2. Synthesis and Growth of Tin Dioxide Nanorods

Unlike the preparation of SnO2 thin films and nanoparticles, the preparation of SnO2 nanorods cannot be done using well established techniques such as sputtering or sol gel. SnO2 nanorods preparation involves annealing precursor powders in which sodium chloride, sodium carbonate and tin chloride are homogeneously mixed. The method is somewhat similar to the molten salt synthesis (MSS) mechanism170,361,362 which is one of the simplest techniques for preparing ceramic powders with whisker-like,363,364 needle-like,365 and plate-like366−368 morphologies. In order to obtain the microemulsion, cyclohexane is used as the oil phase. A mixture of analytical grade poly(oxyethylene)-5-nonyl phenolether (NP5), poly(oxyethylene)-9-nonyl phenolether (NP9) and p-octyl-poly(ethylene glycol) phenylether (OP) with weight ratio 1:1:1 is used as nonionic surfactant. After stirring the mixture until it becomes transparent, 8 mL of 1 M Na2CO3, 32 mL of 2 M NaCl, and 8 mL of 0.5 M SnCl4 aqueous solutions are added to 60, 100, and 60 mL of the mixture followed by further stirring. The mixture with Na2CO3 (μE1) is then mixed with the mixture with NaCl (μE2), followed by adding the mixture with SnCl4 (μE3) to the system (μE1 + μE2). Vigorous stirring conditions are maintained throughout the process. Precursor powders are formed from drying the solution μE4 (μE1 + μE2 + μE3) containing Na2CO3, NaCl, and SnCl4 after washing with acetone. The dried precursor powders are calcined in a tube furnace operating at 800 °C for 2 h in an oxygen partial pressure of 3 × 10−2 Pa atmosphere, followed by natural cooling to room temperature. Distilled water is then used to remove the salts from the resulting mixtures. When the precursor powders are calcined in an oxygen environment, they exchange electrons with the oxygen species. The oxygen atoms enter the lattice of the powder material through rapid grain boundary diffusion. This results in a reduction of oxygen vacancies which promotes the growth of defect free SnO2 nanorods. The resulting SnO2 nanorods can then be used to fabricate various device structures. The electrical properties of the nanorods can be characterized by constructing a single SnO2 nanorod FET. The fabrication process involves thermally growing SnO2 films with 500 nm thickness on top of n-type silicon wafers, followed by deposition of gold metal using 7470


Chemical Reviews

Review

A typical TEM picture of clusters of SnO2 nanorods is shown in Figure 50. The nanorods were prepared by calcining

sputtering. Standard optical lithography technique is used to pattern parallel gold-electrode pairs, with a length of 15 μm and a thichness of 90 nm. The parallel gold electrode pair forms the source and drain electrodes of the field effect transistor. The separation between the parallel gold electrode pair is typically 3 μm with an n-type silicon layer serving as the back gate. The SnO2 nanorods are agitated into a suspension in an ultrasonic bath of ethanol. The suspension is then dispersed onto an ntype silicon chip coated with 500 nm of SnO2 with predefined gold electrodes. The SnO2 field effect transistor samples are rapidly transferred to a sealed oven with an argon atmosphere for annealing. The argon gas minimizes the oxidation of the SnO2 nanorods during annealing. Typical annealing condition for improving the quality of gold/SnO2 contact is 500 °C for 10 min.369 A typical XRD pattern of SnO2 nanorods synthesized at 800 °C for 2 h in an atmosphere with oxygen partial pressure of 3 × 10−2 Pa is shown in Figure 48. All diffraction peaks can be

Figure 50. Typical TEM bright-field image of clusters of SnO2 nanorods. Reprinted with permission from ref 369. Copyright 2009, Elsevier.

precursor powders at 800 °C for 2 h in an atmosphere with oxygen partial pressure of 3 × 10−2 Pa. The nanorods have lengths ranging from several hundred nanometers to a few micrometers, and diameters ranging from ten to 25 nanometers. A typical TEM bright field image of a single SnO2 nanorod is ahown in Figure 51a. The nanorod is about 20 nm in width

Figure 48. Typical XRD pattern of the SnO2 nanorods obtained using Cu Kα radiation (1.5406 Å). Reprinted with permission from ref 369. Copyright 2009, Elsevier.

indexed to the tetragonal rutile structure of SnO2 with lattice constants a = 4.738 Å and c = 3.187 Å, which are consistent with the standard values for bulk SnO2.31 Purity of the SnO2 nanorods is confirmed by the absence of characteristic peaks of impurities such as other forms of tin oxides. The presence of the copper peaks is due to the use of copper grid for the sample. Result of energy dispersive X-ray analysis (EDS) of the SnO2 nanorods is shown in Figure 49. The chemical composition of the nanorods analysized by energy dispersive X-ray spectrometry is 33.5 at. % tin and 66.5 at. % oxygen. The ratio Sn:O = 1:1.985 is in good agreement with that of bulk SnO2 (Sn:O = 1:2).

Figure 51. (a) Typical TEM bright-field image of a single SnO2 nanorod with [110] growth direction. Its growth direction is shown by the arrow. The inset at the upper left-hand corner shows the end of a single SnO2 nanorod. The inset at the bottom right-hand corner shows the SAED pattern of the single SnO2 nanorod. (b) A HRTEM image of the single SnO2 nanorod shown in panel a, where the arrow shows its [110] growth direction. Reprinted with permission from ref 369. Copyright 2009, Elsevier.

and larger than 550 nm in length. It is very smooth, straight and uniform, with almost perfect geometrical shape. Energy dispersive X-ray analysis shows the chemical composition of the nanorod to be close to SnO2. The end of a single SnO2 nanorod is shown at the inset at the upper left-hand corner of Figure 51a. The growth directions of the nanorods are determined by SAED and HRTEM. The inset at the bottom

Figure 49. Energy-dispersive X-ray spectrum (EDS) of SnO2 nanorods. Reprinted with permission from ref 369. Copyright 2009, Elsevier. 7471


Chemical Reviews

Review

right-hand corner of Figure 51a is the SAED pattern of the single SnO2 nanorod shown in this figure. The d-spacing of (110) crystal plane is 3.35 Å, (200) crystal plane is 2.37 Å, (101) crystal plane is 2.64 Å, and (310) crystal plane is 1.50 Å. A more-detailed analysis can be made based on the highly magnified HRTEM image shown in Figure 51b. The growth direction of the nanorod is confirmed to be [110] as indicated by arrow parallel to the long axis of the nanorod. The field of view mainly consists of a single-domain crystallite in the high resolution electron microscopy image. There are recurrent values of separation distance between lattice layers, in particular, 0.33 nm, as shown in Figure 51b. This corresponds to the lattice parameter of the rutile structure of the SnO2 cassiterite phase arising from [110] reflections. These results are also confirmed by analysis of SAED patterns obtained by TEM. The data indicate that the nanorod segregates as a single phase belonging to the SnO2 cassiterite structure. It is a structurally uniform single crystal with a growth direction of [110]. There are no dislocations or other planar defects in the nanorod. The Raman spectrum of SnO2 nanorods is shown in Figure 52. There are three fundamental Raman scattering peaks at

Figure 53. Current−voltage (IDS − VDS) curves of field-effect transistor devices for the back-gated single SnO2 nanorod with VGS = +5 to −1 V in −2 V steps from top to bottom. Reprinted with permission from ref 369. Copyright 2009, Elsevier.

by the increase in conductance when the back gate voltage is increased. This might be due to oxygen vacancies and extra tin interstitial atoms in the lattice during the annealing of precursor powders. At VG from −1 to +5 V with VDS = 15 mV, an on/off current ratio (Ion/Ioff) as high as 107 has been achieved. The mobility and subthreshold voltages can be calculated from ideal FET equations. The total charge on the nanorods is Q = CVg,t, where C is the nanorod capacitance and Vg,t is the threshold voltage required to completely deplete the charge on the rod. The capacitance of the nanorod is given by eq 1, where ε = 3.9, ε0 = 1, and h = 500 nm are the relative dielectric constant, the vacuum dielectric constant, and the thickness of the silicon dioxide layer, respectively. r = 10 nm is the nanorod radius and L = 3 μm is the channel length. The IDS − VGS relationship of the single SnO2 nanorod FET device in log scale at VDS = 15 mV is shown in Figure 54. The subthreshold slope is about 0.06 V/decade and is independent of the bias voltage. An on/off current ratio as high as 107 is achieved at VDS = 15 mV, in agreement with that obtained in planar single-crystalline SnO2 nanorods and other mental oxide nanorods.106,107 The one-dimensional electron density is n =

Figure 52. Room temperature Raman spectrum of SnO2 nanorods. Reprinted with permission from ref 369. Copyright 2009, Elsevier.

477.8, 638.6, and 774.8 cm−1, in good agreement with those of rutile SnO2 single crystal. These three peaks are identified to the Eg mode, the A1g mode and the B2g mode, respectively. These are the typical features of the rutile phase of synthesized SnO2 nanorods.116,370 Sodium chloride and the surfactant have major effects on the nucleation and growth of SnO2 nanorods. The viscosity of the melt is significantly reduced by sodium chloride, resulting in an increase in the mobility of components in the flux. This provides a favorable environment for the nucleation and growth of nanorods. The surfactant (NP9/NP5/OP) is favorable to the formation of fine particles. It produces a surrounding shell to prevent the particles from aggregation during the formation of the precursor. The surfactant acts as a template, with the template action resulting in the expitaxial growth of the products. With this knowledge of the formation process of SnO2 nanorods, the preparation route can be extended to producing other nanostructured semiconducting metal oxides by suitable choice of the appropriate surfactant. The current−voltage (IDS − VDS) relationship of the SnO2 nanorod FET device is shown in Figure 53. For the room temperature samples, the applied gate voltage varies from −1 to +5 V and the drain voltage from 0 to 1.0 V. The n-type semiconductor characteristic of the SnO2 nanorods is revealed

Figure 54. IDS − VGS curve of field-effect transistor device for the single SnO2 nanorod in log scale at VDS = 15 mV. Reprinted with permission from ref 369. Copyright 2009, Elsevier. 7472


Chemical Reviews

Review

Q/eL ≈ 1.17 × 107 cm−1. The mobility of the electrons can also be deduced from the trans-conductance of the single SnO2 nanorod FET. The trans-conductance gm is given by eqs 2 and 3 in the linear region. The mobility of the channel stoichiometry in the single SnO2 nanorod FET device is about 82 cm2 V−1 s−1. Field effect transistors made from SnO2 nanorods have good electrical properties in terms of the on/off ratio and threshold voltage of the device. The SnO2 nanorods are n-type semiconductors, as confirmed by electrical measurements. Since other one-dimension nanostructural materials may be manipulated using this simple technique, there are new opportunities in applications as building blocks for nanoelectronics and active sensing materials. 4.3. Doping and Modification of Tin Dioxide Nanorods

Many metal oxide semiconductor nanomaterials have unique properties and potential applications in various nanodevices. These include ZnO, In2O3, Ga2O3, WO3, and SnO2. SnO2 is a stable and large band gap (n-type) semiconductor. It has been widely used in gas sensors due to its low cost, high sensitivity, fast gas response and high stability.371 Doping with selected dopants or modulating the morphology could enhance the gassensing properties of SnO2 nanomaterials. The synthesis of SnO2 nanomaterials with various morphologies such as nanorods, nanowires, nanotubes, nanobelts, nanosheets, and assembled hollow nanospheres has been extensively reported in the past decade.173,262,372−378 They have high sensitivity in gas detection due to the ultrahigh surface to volume ratio of nanomaterials. Doping with extrinsic dopants is a facile and effective way to improve gas detection sensitivity.379−382 In particular, formaldehyde of concentration as low as 30 ppb could be detected by Pd-doped SnO2 microgas sensors.383 The morphology and microstructure of doped, or loaded, nanorods could be characterized by field emission SEM.384 Figure 55a shows a panoramic field emission SEM image of the hierarchical unloaded SnO2 nanorods obtained after hydrothermal treatment at 200 °C for 24 h. The picture shows flower-like structures with an average diameter of about 2 μm for each “flower”. There are no other morphologies, indicating a high yield of this three-dimensional flower-like microstructure. Figure 55b shows that each “flower” is made up of uniform nanorods with diameters of 100−150 nm and lengths of 400−500 nm. The inset of Figure 55b shows an individual SnO2 nanorod. There is a sharp tip at the end of the nanorod. Images of samples doped with Pd are shown in Figure 55c−h. The morphology of the Pd loaded samples is similar to that of the unloaded ones. However, there is a decrease in the diameter of the as-prepared nanorods as the amount of Pd loading increases. This may be due to suppression of grain boundary migration and increasing energy barrier for grain growth.385 Sensors based on unloaded and hierarchical Pd-loaded SnO2 nanorods have been constructed. The response of the sensor using hierarchical Pd-loaded SnO2 nanorods to 1000 ppm butanone is about ten times higher than that of the sensor based on unloaded SnO2. The chemical and electrical contribution of Pd loading appears to be the cause of the enhanced gas sensing performance. A rapid, continuous and scalable method of synthesis of single crystalline SnO2 nanorods has been reported by Liu et al.386 The method involves a flame approach by introducing an iron dopant and adjusting the flame residence time (equal to the flame height, FH) to facilitate the anisotropic growth of SnO2. The morphology of SnO2 products for various

Figure 55. FESEM images (a and b) of unloaded SnO2 nanorods. Pdloaded SnO2 nanorods obtained with different amounts of Pd (c and d) 1.3 wt %, (e and f) 3.8 wt %, and (g and h) 7.4 wt %. Reprinted with permission from ref 384. Copyright 2011, Elsevier.

concentrations of iron dopant is shown in Figure 56a−e. At a flame height of 50 cm, there are only a small number of nanorods formed at 1.0 at. % iron doping. The number of SnO2 nanorods increases and become more elongated as the iron doping concentration is increased to 1.5 at. %. As the iron doping concentration is increased to 2.0 and 2.5 at. %, there are well-defined nanorods with diameters of about 20−40 nm and lengths up to 200 nm. When iron dopant is incorporated into the SnO2 lattice, it selectively disrupts a specific SnO2 crystal plane, promoting further crystal oriented growth into nanorods. On the other hand, lithium and zinc have lower valence than iron. Doping with these elements does not promote nanorod growth (Figure 56h,i. Flame height has a major effect on SnO2 morphology. The size of the nanoparticles is small (about 20 nm) when the flame height is 15 cm (Figure 56g). At a flame height of 30 cm, some nanoparticles have disappeared and larger nanorods are formed (Figure 56f). When the flame height is 50 cm, well-defined nanorods are formed and the small nanoparticles have completely disappeared (Figure 56e). Increase in flame height prolongs the residence time of the reagents in the high-temperature zone and promotes nanorod formation. The presence of crystal defects and rodlike morphology are probably the reasons for the photoluminescence (PL) spectrum of SnO2 nanorod to exhibit a broad, strong orange-emission peak at around 620 nm. This property 7473


Chemical Reviews

Review

Figure 56. Nanoparticles and nanorods images: (a) pure SnO2 with FH 50 cm, (b−e) 1.0−2.5 at. % Fe-doped SnO2 with FH 50 cm, showing that the morphologies of the Fe-doped SnO2 clearly change from cuboid to nanorod as Fe dopant concentration increases, (f and g) 2.5 at. % Fe-doped SnO2 with FH 30 and 15 cm, and (h and i) 2.5 at. % Li, Zn-doped SnO2 with FH 50 cm. Reprinted with permission from ref 386. Copyright 2010, American Chemical Society.

inhibit charge recombination at the FTO-electrolyte. This will improve photocurrent density, photovoltage, fill factor, and power conversion efficiency. Moreover, the compact titanium dioxide layer improves adherence between the SnO2 nanorods and FTO surface and provides more electron pathways from the SnO2 film to the FTO glass. This is beneficial to electron transfer. The electron-transfer efficiency is enhanced as a consequence. After TiCl4 post-treatment of the SnO2 photoelectrodes, the η of DSSC-3 has improved by 231% compared to that of DSSC-1. This is due to the increase in JSC, VOC, and FF. The JSC, VOC, FF, and η of DSSC-3 (VOC = 0.689 V, JSC = 10.12 mA/cm2, FF = 57.3%, and η = 4.00%) are considerably higher than those of DSSC-1. The conduction band edge of SnO2 is more positive than that of titanium dioxide by about 300 mV. As a consequence, the open circuit photovoltage of SnO2 photoelectrode is greatly reduced. Moreover, the energy barrier formed at the SnO2 matrix surface may be the reason for the higher JSC and FF of the DSSC-3 achieved by the titanium dioxide coating. The rate of the recombination process of the photoinjected electrons returning to the oxidized dye or ions is reduced by this energy barrier. DSSC-4 shows the highest efficiency of 4.67% after the introduction of both titanium dioxide compact layer and TiCl4 post-treatment. The dark current of the DSSCs as a function of the applied potential is shown in Figure 58. It is in the following decending order: DSSC-1 > DSSC-2 > DSSC-3 > DSSC-4. The dark current originates from the recombination between photogenerated electrons in the anode and I3− ions in the electrolyte. Since the conduction band energy level of SnO2 is lower than that of

may be exploited for potential applications in optoelectronics.386 A simple hydrothermal method for the synthesis of SnO2 nanorods for application in dye-sentitized solar cells has been reported by Shang et al.387 A SEM image of the SnO2 nanorods is shown in Figure 57a. The picture shows numerous welldefined nanorods. At higher magnification (Figure 57b), the diameters of the SnO2 nanorods are around 40−80 nm with lengths of about 200 nm. There are also SnO2 nanoparticles of about 50 nm diameter (Figure 57c). A top view SEM image of a SnO2 nanorod photoanode is shown in Figure 57d. The preparation process of making the slurry, coating on FTO glass and sintering of the paste does not appear to change the morphology of the SnO2 nanorods. SEM shows that the SnO2 nanorod structure is maintained after coating with titanium chloride treatment (Figure 57e). However, the surface of the nanorods is not as smooth as the uncoated ones. A SEM image of the cross section of a SnO2 nanorod photoanode with a thickness of about 10 μm is shown in Figure 57f. The current density−voltage (J−V) characteristics of dye sentitized solar cells (DSSC) based on four types of films (DSSC-1, bare SnO2 nanorods; DSSC-2, TiO2 compact layer + SnO2 nanorods; DSSC-3, SnO2 nanorods + TiCl4 posttreatment; DSSC-4, TiO2 compact layer + SnO2 nanorods +TiCl4 post-treatment), and the detailed photovoltaic parameters are tabulated in Table 2. The JSC, VOC, FF, and η of DSSC-2 (VOC = 0.575 V, JSC = 8.06 mA/cm2, FF = 52.2%, and η = 2.42%) are considerably higher than those of DSSC-1 (VOC = 0.530 V, JSC = 4.58 mA/cm2, FF = 49.7%, and η = 1.21%). The introduction of titanium dioxide compact layer could 7474


Chemical Reviews

Review

titanium dioxide,388 the former establishes a bridge between titanium dioxde and FTO glass for the transfer of photogenerated electrons. This reduces the recombination between photogenerated electrons in the anode and I3− ions in the electrolyte. Thus, the dark current for DSSC-2 is smaller than that for DSSC-1. The adherence between the SnO2 and FTO surface is thought to be improved by the titanium dioxide compact layer. Electron transfer and subsequent suppression of dark current is assisted since there are more electron pathways from SnO2 film to FTO. Thus, the dark current for DSSC-3 is smaller than that for DSSC-2. Moreover, the TiCl4 post-treatment establishes an energy barrier at the SnO2-electrolyte interface. This retards the back electron transport, reduces the rate of interfacial recombination, and suppresses dark current. Therefore, DSSC-4 has the lowest dark current among the four cells. Larger dark current implies smaller dark-voltage start point. Yang et al. and Kang et al. have reported similar phenomena,389,390 although the dark-voltage start point is not the same as the light-voltage drop point. 4.4. Tin Dioxide Nanorods Planted Graphites

In view of its high energy density and excellent cycle life, lithium ion battery has dominated the power source market for portable electronic devices, power tools, and electric vehicles.391 The anodes of most lithium ion batteries are still graphite based even though the theoretical capacity of 372 mAh/g is not sufficient for modern equipment.392 The development of higher capacity anode materials is currently being pursued in earnest. The search for better anode materials include various metal oxides,393−398 polymers,392 carbon materials,399 and their composites.327,400−402 Among them, SnO2-based materials are promising substitutes for the commercial graphite anodes because of their low cost, safety, and high theoretical lithium storage capacity (about 782 mAh/ g). However, there are drawbacks. The Li+ insertion and extraction process result in large volume expansion/contraction and severe particle aggregation. This leads to pulverization and loss of interparticle contact, culminating in a large irreversible capacity loss and poor cycling stability.403−405 The SEM pictures of SnO2 nanrod planted graphite material are shown in Figure 59. The nanorods have different diameters and lengths. They are rectangular in shape and are densely distributed throughout the graphite surface. The growth of SnO2 nanorods occurs according to the following reactions406,407

Figure 57. SEM images of SnO2 nanorods at different magnifications (a, b), tin dixide nanoparticles (c), top view SEM images of the SnO2 nanorod photoanode without (d) and with (e) TiCl4 treatment, and cross-sectional SEM image of the SnO2 nanorod photoanode with TiCl4 treatment (f). Reprinted with permission from ref 387. Copyright 2013, American Chemical Society.

Table 2. Photovoltaic Parameters of the DSSCs Based on SnO2 Nanorodsa DSSCs

VOC (V)

JSC (mA/cm2)

FF (%)

η (%)

DSSC-1 DSSC-2 DSSC-3 DSSC-4

0.530 0.575 0.689 0.757

4.58 8.06 10.12 10.51

49.7 52.2 57.3 58.8

1.21 2.42 4.00 4.67

a

Reprinted with permission from ref 387. Copyright 2013, American Chemical Society.

Sn 4 + + 4OH− → Sn(OH)4

Sn(OH)4 + 2OH− → [Sn(OH)6 ]2 − [Sn(OH)6 ]2 − → SnO2 + 2H 2O + 2OH−

The [Sn(OH)6]2− concentration and molar ratio of SnCl4· 5H2O to NaOH play a crucial role in obtaining rectangular SnO2 nanorods.408 Kim et al. reported that the appropriate [Sn(OH)6]2− concentration for SnO2 nanorod growth is above 0.05 M.405 A molar ratio of SnCl4·5H2O to NaOH of around 1:10.5−1:24 is suitable for the successful synthesis of SnO2 nanorod planted graphite. The effect of [Sn(OH) 6 ] 2− concentration and reaction time on the morphology of synthesized SnO2 nanorod-planted graphite is shown in Figure 59. In this figure, the samples are denoted as S1, S2, S3, S4, S5, and S6 under the specified synthesis condition. The average

Figure 58. J−V curves of the DSSCs and dark J−V curves of the DSSCs. Reprinted with permission from ref 387. Copyright 2013, American Chemical Society.

7475


Chemical Reviews

Review

Figure 59. Typical SEM images of nanostructures of as-prepared SnO2 nanorod planted graphite materials (with cross-sectional images in the insets) as a function of [Sn(OH)6]2− concentration and reaction time. Reprinted with permission from ref 405. Copyright 2013, American Chemical Society.

Figure 60. (a) Voltage profiles of SnO2 nanorod-planted graphite (S1) at a current density of 72 mA/g between 0.01 and 1.5 V. The inset shows an SEM image of SnO2 nanorod arrays after cycling 25 times. (b) Capacity cycle number curves of as-obtained SnO2 nanorod-planted graphite (S1). The inset shows the differential capacity versus voltage plot. (c) Relations of the capacity fade and Coulombic efficiency with the nanorod length of SnO2 nanorod-planted graphite samples S1−S6. (d) Relations of the SnO2 contents and first discharge capacity over the different nanorod length of SnO2 nanorod-planted graphite samples S1−S6. Herein, the error limits of all the samples were determined statistically from the galvanostatic charge/discharge measurements. Reprinted with permission from ref 405. Copyright 2013, American Chemical Society.

high initial discharge capacity of about 1010 mAh/g. This value is between that of SnO2 and graphite. This can be explained by the efficient electron transport, large interfacial area and improved kinetic properties of the one-dimensional SnO2 nanorod.409,410 Moreover, there is stable capacity retention after the first cycle, indicating homogeneous dispersion of electroactive composites in the electrode film without aggregation. The SnO2 nanorod planted graphite composite electrode has higher initial Coulombic efficiency (59.2%) than the theoretical value (52%) for SnO2 electrode under full Li

diameter of the prepared SnO2 nanorods is controlled between 28 and 84 nm, and the length from 123 to 646 nm. Larger diameter and longer nanorods could be obtained with a larger [Sn(OH)6]2− concentration. When the [Sn(OH)6]2− concentration is below 0.1 M, e.g. 0.05 M, the time required to produce SnO2 nanorods is extended from 24 to 48 h. The voltage profiles of SnO2 nanorod planted graphite (S1) cycled at a current density of 72 mA/g in the potential range of 0.01−1.5 V (versus Li/Li+) for up to 25 cycles are shown in Figure 60a. The SnO2 nanorod planted graphite delivers a very 7476


Chemical Reviews

Review

alloying/dealloying.411 It is also higher than the values of Coulombic efficiencies typically in the range of 40 to 50% of SnO2 based materials previously reported.409,412 The initial irreversible capacity loss is mainly due to electrolyte decomposition on electro-active materials. As irreversible side reactions are suppressed on carbon compared with other materials, such as silicon, Fe2O3 and Co3O4,413−415 SnO2 nanorod graphite composites could have a higher Coulombic efficiency. The electrochemical impedance spectroscopy films covering the surface of SnO2 nanorod planted graphite prevent the electrolyte from being further decomposed after the first discharge process. Consequently, the Coulombic efficiency increases to 94.2% in the second cycle. The inset SEM image shows the preserved SnO2 nanorod arrays on the graphite core after 25 cycles. In spite of volume variations, structural integrity of the nanorod arrays is maintained. This indicates that the loss of electrical contact with the SnO2 nanorods has declined. The galvanostatic cycling profiles of a SnO2 nanorod planted graphite (S1) electrode are shown in Figure 60b. The cycling performance is enhanced and reversible capacity is maintained over 25 cycles. The fading of average capacity of SnO2 nanorod planted graphite is observed to be 0.85% per cycle after the second cycle. There is good capacity retention as shown in Figure 60c. It is much smaller than the previously reported values of SnO2 nanoparticles, nanowires, and nanorods when they are used alone in the anode system.409,412,416 The elasticity of carbon is larger than that of SnO2.417 Therefore, the elastic graphite intersp aced between nanorods can accommodate strain energy effectively when SnO2 nanorods and graphite react with Li+. Thus, SnO2 nanorod planted graphite has good cyclability. Nonetheless, in the composites with excessively dense SnO2 nanorod arrays, such as S5 or S6, the performance is worse than that of sparsely grown SnO2 nanorod arrays, such as S1. Limited space between nanorods might lead to a lack of electrolyte permeability and strain relaxation, culminating in poor performance. Figure 60d shows that the first discharge capacity increases as the SnO2 nanorod length increases. The SnO2 contents inferred from the theoretical capacity values (2Li2O and Li4.4Sn, 1494 mAh/g; LiC6, 372 mAh/g) at the first discharge capacity are in accordance with the previous TGA measurements. This suggests that both SnO2 nanorods and graphite contribute to the total capacity of the electrode. The rate performance of SnO2 nanorod planted graphite has been investigated by Kim et al.405 They cycled the cell from a current density of 72 mA/g to 288 mA/g at intervals of 72 mA/ g (Figure 61). Even at a high current density of 288 mA/g, the sample S1 is still able to deliver a substantial amount of capacity of 257.7 mAh/g. For S3 and S5, the delivered capacity is 249.3 and 248.2 mAh/g, respectively. For S1, 81.7% of the initial capacity is recovered again when the current density is reduced back from 288 to 72 mA/g. These authors compared the rate capability of individual SnO2 nanowires and nanoparticles under identical test conditions in order to discover the reasons for improved performance of the SnO2 nanorod planted graphite. Synthesized SnO2 nanowires have diameters of about 80 nm and micrometer-sized length, and SnO2 nanoparticles have diameters of about 100 nm. Tin dioxide nanowires and nanoparticles deliver discharge capacities of 242.5 and 192.9 mAh/g respectively, at a current density of 288 mA/g. The respective recovered capacity ratios are 58.8 and 34.4%. The normalized capacity of SnO2 nanowires (81.8%) is higher than that of SnO2 nanorod planted graphite (S1, 79.5%) at a low current density of 144 mA/g, but the capacity retention of

Figure 61. Cycling performance at various current densities between 0.01 and 1.5 V of SnO2 nanorod planted graphite (S1, S3, and S5). The inset shows the normalized capacity at each step by the average capacity values under 0.72 mA/g current density of the first step. Reprinted with permission from ref 405. Copyright 2013, American Chemical Society.

SnO2 nanorod planted graphite (S1, 46.2%) is improved significantly compared to that of SnO2 nanowires (34.5%) at a high current density of 288 mA/g. Thus, the rate capacity fading of SnO2 nanorod planted graphite decreases compared to that of SnO2 nanowires as the current rate increases. The rate performance of SnO2 nanorod planted graphite (S1, S3, and S5) is better than that of SnO2 nanoparticles and nanowires, indicating stable capacity retention and a high recovered capacity ratio. This shows that one-dimensional SnO2 combined with graphite can improve the rate capability of SnO2 based electrodes. This novel SnO2 nanorod anode material has a better rate capability than SnO2 nanowires and nanoparticles at relatively high current rates. The enhanced electrochemical properties may originate from the effectiveness of the novel nanostructure of SnO2 nanorod planted graphite. SnO2 nanorods with appropriate interspacing could prevent aggregation of the electroactive material and enhance the Li+ transfer rate during cycling. The graphite, as a buffer matrix, has high electric conductivity. This can improve electron transport and the Coulombic efficiency of SnO2 based materials.

5. CONCLUSIONS 5.1. Summary

This article has reviewed important advances in the field of semiconductor SnO2 thin films, nanowires, and nanorods in recent years. The topics included the preparation, characterization, defects, growth, micro/nanostructure, mechanism, doping, modification, composites, properties, and applications during the formation processes of SnO2 thin films, nanowires, and nanorods. Nanostructured SnO2 semiconductor materials have widely attracted enormous interests and have become one of the hottest research pursuits in the past decade due to their unique geometry, morphology dependent chemical and physical attributes, and potential applications. Major advancements have been achieved in the synthesis, characterization, growth, micro/nanostructure eveolution, properties and applications for SnO2 nanomaterials. As outlined in this review, many novel SnO2 preparation methods, micro/nanostructural features, and related applications are continually being developed and studied. Improved insights on the growth 7477


Chemical Reviews

Review

*E-mail: [email protected]. *E-mail: mhwu@staff.shu.edu.cn. *E-mail: [email protected].

mechanisms of different SnO2 morphologies, including thin films, nanowires, and nanorods, have been achieved. Such insights have made significant contributions to the progressive development of advanced semiconductor SnO2 nanostrutures with defined morphological structures (thin films, nanowires, and nanorods) with unique chemical and physical properties. At the same time, some new applications of these advanced SnO2 nanostructures are enthusiastically proposed and demonstrated. In fact, the rise of low dimensional SnO2 nanostructured materials has opened up new routes for their use in various materials engineering applications. Significant impacts have also been made along the way. Of course, investigation in the SnO2 based applications still poses many critical problems that need to be addressed and understood. In particular, these crucial problems include: (i) Development of improved methods of structurally controlled growth of semiconductor SnO2 nanomaterials based on adjustment of structural parameters. (ii) Achievement of better understanding of the interaction between low dimensional SnO2 nanomaterials and other elemental species arising from the process of doping, modification, and formation of composites, and associated micro/nanostructure changes caused by composition modulation. (iii) Improvement in the sensitivity, selectivity, and stability of SnO2 based sensors through better understanding of the fundamental mechanisms of signal transduction and detection. (iv) Development of cost-effective mass production methods for functional semiconductor SnO2 nanomaterials and the resultant SnO2 based micro/nanodevices (e.g., mammalian olfactory system electronic nose) for commercialization.418 There is still a lot of room for development in semiconductor SnO2 nanomaterials to realize its full potential. Future work in this field should focus on the above-mentioned areas to enable commercial SnO2 micro/nanodevices to be available for practical applications.

Notes

The authors declare no competing financial interest. Biographies

Zhiwen Chen was born in Hefei, Anhui, China, and received his B.Sc. degree in Chemistry from the Heifei Normal University, M.Sc. degree in Inorganic Chemistry from the University of Science and Technology of China (USTC), and Ph.D. degree in Materials Physics and Chemistry from the City University of Hong Kong (CityUHK). Since 1993, he has worked at USTC as an Assistant Professor, Associate Professor. He joined the Teikyo University in Japan as a Research Fellow in 1999 and the CityUHK as a Research Fellow and Senior Research Fellow in 2001 and 2007, respectively, and is now a Chartered Professor at the Shanghai University, China. His research interests mainly focus on the controllable-synthesis, characterization, and properties of advanced functional materials and oxides nanomaterials. He has published over 160 papers in refereed journals. He served as a Guest Editor, the member of Editorial Board, and Reviewer or Referee for many international journals, and is now Editor-in-Chief of Advances in Materials Physics and Chemistry.

5.2. Concluding Remarks

This work is interdisciplinary in nature. It integrates the areas of chemistry, physics, and materials science. It is intended that this review will provide the reader an overview of the tremendous potential of low dimensional SnO2 functional materials with novel micro/nanostructures. Future research on this topic is likely to lead to the development of further exciting techniques and powerful combinations of existing ones. It is envisaged that these low dimensional SnO2 functional materials with fascinating micro/nanostructures may offer vast and unforeseen opportunities in oxide semiconductor devices as well as in other fields of science and technology. Some challenges still remain in the clarification of the intricate aspects and potential applications of low dimensional SnO2 functional materials. Future research may enable these novel SnO2 functional materials with appropriate micro/nanostructures to be tailormade for a large number of applications. It may also provide new opportunities in the development of SnO2 architectures with the goal of optimizing functional material properties for specific applications.

Dengyu Pan is now a Professor of Chemistry in the Institute of Nanochemistry and Nanobiology, Shanghai University, China. He received his Ph.D. degree from the University of Science and Technology of China in 2003 under the supervision of Professor J. G. Hou. From 2003 to 2006, he pursued postdoctoral research in the group of Professor G. H. Wang in Nanjing University, China. His research interests are centered on the synthesis and applications of nanomaterials.

AUTHOR INFORMATION Corresponding Authors

*Tel.: +86 21 66137503. Fax: +86 21 66137787. E-mail: [email protected]. 7478


Chemical Reviews

Review

environmental chemistry, nanostructures of sensors, and biomedical materials. So far she has authored more than 130 publications and patents.

Zhen Li was born in 1972 in Hunan, China. She received her M.S. and Ph.D. degrees form the Shanghai University in 2001 and 2008, respectively. She was appointed as a Professor at the Shanghai University in 2010. Her research interests focus on the semiconductor nanostructured materials. She has published more than 30 papers in international journals.

Chan-Hung Shek got his Ph.D. in 1994 from the Department of Mechanical Engineering at the University of Hong Kong. He then joined City University of Hong Kong as an Assistant Professor and became an Associate Professor in 2001. He is currently a Professor in the Department of Physics and Materials Science, City University of Hong Kong. Since 2007, he has been appointed as Assistant Head of Department of Physics and Materials Science, and in 2011 he was concurrently appointed as Assistant Dean of College of Science and Engineering. His major research interests are the correlation among processing, microstructures and properties of materials, including

Zheng Jiao was born in 1972 in Peking, China. After he got his Ph.D. degree in Inorganic Chemistry from the University of Science and Technology of China in 2000, he worked at the Osaka University in Japan from 2002 to 2003 and at the Lille University of Technology in France in 2004. Now he works at the Shanghai University in China as a full professor. His research interests are focused on the functional materials and applications in environmental and biochemistry science.

nanostructured metal oxides, bulk metallic glasses, and conventional engineering alloys.

C. M. Lawrence Wu received his Ph.D. from the University of Bristol, U.K. in 1986 and has been working in City University of Hong Kong Minghong Wu earned her Ph.D. from the Chinese Academy of Science in 1999. After some years of postdoctoral research in Japan, she became a Professor in 2002 at the School of Environmental and Chemical Engineering of Shanghai University in China. She was the Distinguished Young Scientist supported by the National Natural Science Foundation of China. Her main research interests include the

since 1987. He is currently a Professor in the Department of Physics and Materials Science. Apart from working on the synthesis and characterization of nanocrystals, he also works on plasmonic effects in the application of biosensing and solar energy. 7479


Chemical Reviews

Review

RTJN

random tunnelling junction network

REFERENCES (1) Batzill, M.; Diebold, U. Prog. Surf. Sci. 2005, 79, 47. (2) Chen, Z. W.; Wang, X. P.; Tan, S.; Zhang, S. Y.; Hou, J. G.; Wu, Z. Q. Phys. Rev. B 2001, 63, 165413. (3) Chen, Z. W.; Shek, C. H.; Lai, J. K. L. Mater. Sci. Eng., A 2004, 385, 455. (4) Wang, L. J.; Chen, X. J.; Chen, C.; Liu, Y. Y.; Chen, Z. W.; Shek, C. H.; Wu, C. M. L.; Lai, J. K. L. J. Phys. Chem. C 2012, 116, 21012. (5) Chen, Z. W.; Li, Q. B.; Pan, D. Y.; Zhang, H. J.; Jiao, Z.; Wu, M. H.; Shek, C. H.; Wu, C. M. L.; Lai, J. K. L. Mater. Today 2011, 14, 106. (6) Persson, M. P.; Lherbier, A.; Niquet, Y. M.; Trizon, F.; Roche, S. Nano Lett. 2008, 8, 4146. (7) Nguyen, P.; Ng, H. T.; Meyyappan, M. Adv. Mater. 2005, 17, 1773. (8) Chen, Z. W.; Lai, J. K. L.; Shek, C. H. Phys. Lett. A 2005, 345, 218. (9) Wang, L. J.; Chen, C.; Liu, Y. Y.; Chen, Z. W. Surf. Coat. Technol. 2012, 206, 5091. (10) Chen, Z. W.; Jiao, Z.; Wu, M. H.; Shek, C. H.; Wu, C. M. L.; Lai, J. K. L. J. Nanosci. Nanotechnol. 2012, 12, 26. (11) Chen, Z. W.; Li, Q. B.; Pan, D. Y.; Li, Z.; Jiao, Z.; Wu, M. H.; Shek, C. H.; Wu, C. M. L.; Lai, J. K. L. J. Phys. Chem. C 2011, 115, 9871. (12) Chen, Z. W.; Lai, J. K. L.; Shek, C. H. Chem. Phys. Lett. 2006, 422, 1. (13) Qin, L. P.; Xu, J. Q.; Dong, X. W.; Pan, Q. Y.; Cheng, Z. X.; Xiang, Q.; Li, F. Nanotechnology 2008, 19, 185705. (14) Chen, C.; Wang, L. J.; Liu, Y. Y.; Chen, Z. W.; Pan, D. Y.; Li, Z.; Jiao, Z.; Hu, P. F.; Shek, C. H.; Wu, C. M. L.; Lai, J. K. L.; Wu, M. H. Langmuir 2013, 29, 4111. (15) Chen, Z. W.; Lai, J. K. L.; Shek, C. H. J. Non-Cryst. Solids 2005, 351, 3619. (16) Chen, Z. W.; Lai, J. K. L.; Shek, C. H. Phys. Lett. A 2005, 345, 391. (17) Du, J.; Wang, J.; Jiao, Z.; Wu, M. H.; Shek, C. H.; Wu, C. M. L.; Lai, J. K. L.; Chen, Z. W. J. Nanosci. Nanotechnol. 2011, 11, 9709. (18) Chen, Z. W.; Jiao, Z.; Pan, D. Y.; Li, Z.; Wu, M. H.; Shek, C. H.; Wu, C. M. L.; Lai, J. K. L. Chem. Rev. 2012, 112, 3833. (19) Chen, Z. W.; Zhang, S. Y.; Tan, S.; Li, F. Q.; Wang, J.; Jin, S. Z.; Zhang, Y. H. J. Cryst. Growth 1997, 180, 280. (20) Tamal, T.; Ichinose, N.; Kawanishi, S.; Nishii, M.; Sasuga, T.; Hashida, I.; Mizuno, K. Chem. Mater. 1997, 9, 2674. (21) Mazeina, L.; Picard, Y. N.; Caldwell, J. D.; Glaser, E. R.; Prokes, S. M. J. Cryst. Growth 2009, 311, 3158. (22) Wang, J.; Du, J.; Chen, C.; Li, Z.; Jiao, Z.; Wu, M. H.; Shek, C. H.; Wu, C. M. L.; Lai, J. K. L.; Chen, Z. W. J. Phys. Chem. C 2011, 115, 20523. (23) Ding, X. H.; Zeng, D. W.; Xie, C. S. Sens. Actuators B 2010, 149, 336. (24) Xue, X. Y.; Chen, Z. H.; Ma, C. H.; Xing, L. L.; Chen, Y. J.; Wang, Y. G.; Wang, T. H. J. Phys. Chem. C 2010, 114, 3968. (25) Wang, Y.; Lee, J. Y.; Deivaraj, T. C. J. Phys. Chem. B 2004, 108, 13589. (26) Wijeratne, K.; Akilavasan, J.; Thelakkat, M.; Bandara, J. Electrochim. Acta 2012, 72, 192. (27) Huang, X. W.; Liu, Z. J.; Zheng, Y. F.; Nie, Q. L. Chin. Chem. Lett. 2010, 21, 999. (28) Zhao, H. Y.; Li, Y. H.; Yang, L. F.; Wu, X. H. Mater. Chem. Phys. 2008, 112, 244. (29) Alizadeh, R.; Najafi, N. M.; Poursani, E. M. A. J. Pharm. Biomed. Anal. 2012, 70, 492. (30) Chen, Z. W.; Zhang, S. Y.; Tan, S.; Wang, J.; Jin, S. Z. Mater. Res. Bull. 1999, 34, 1583. (31) Chen, Z. W.; Tan, S.; Zhang, S. Y.; Wang, J.; Jin, S. Z.; Zhang, Y. H.; Sekine, H. Japn. J. Appl. Phys. 2000, 39, 6293. (32) Chen, Z. W.; Zhang, S. Y.; Tan, S.; Wang, J.; Jin, S. Z.; Hou, J. G. J. Mater. Sci. Lett. 2002, 21, 411.

Joseph K. L. Lai won an open scholarship to study Physics at Keble College, Oxford University in 1971 and graduated with first class honours in 1974. He joined the Central Electricity Research Laboratories (CERL) in U.K. after graduation and was appointed Project Leader of the Remaining Life Study Group in 1984. While working at CERL he studied for a Ph.D. at the City University, London, U.K. and was awarded the degree in 1982. He returned to Hong Kong in 1985 and has worked at CityU ever since. He is now Chair Professor of Materials Science in the Department of Physics and Materials Science. Professor Lai is a Fellow of the Hong Kong Institution of Engineers, a Chartered Engineer, and a Fellow of the following U.K. professional institutions: Institute of Materials, Minerals and Mining, Institute of Physics, and Institution of Mechanical Engineers.

ACKNOWLEDGMENTS The work described in this article was financially supported by the National Natural Science Foundation of China (Project Nos. 11375111, 11074161, 11025526, 41373098, and 41173120), the Research Fund for the Doctoral Program of Higher Education of China (Project No. 20133108110021), the Key Innovation Fund of Shanghai Municipal Education Commission (Project Nos. 14ZZ098 and 10ZZ64), the Science and Technology Commission of Shanghai Municipality (Project Nos. 14JC1402000 and 10JC1405400), the Shanghai Pujiang Program (Project No. 10PJ1404100), and the Program for Innovative Research Team in University (Project No. IRT13078). This work was also supported by a General Research Fund from the Research Grants Council, Hong Kong (Project No. CityU 119212). LIST OF SnO2 PLD XRD SEM TEM HRTEM AFM SAED EDS MFM UV IR FET RMS FIPS D

ABBREVIATIONS tin dioxide pulsed laser deposition X-ray diffraction scanning electron microscopy transmission electron microscopy high resolution transmission electron microscopy atomic force microscopy selected area electron diffraction energy dispersive X-ray spectroscopy magnetic force microscopy ultraviolet infrared field-effect transistor root-mean-square fractal images process software dimension 7480


Chemical Reviews

Review

(70) Wang, Z. L.; Kang, Z. C. Functional and Smart MaterialsStructural Evolution and Structure Analysis; Plenum Press: New York, 1998. (71) Pan, Z. W.; Dai, Z. R.; Wang, Z. L. Science 2001, 291, 1947. (72) Chen, Z. W.; Lai, J. K. L.; Shek, C. H. Phys. Rev. B 2004, 70, 165314. (73) Chen, Z. W.; Lai, J. K. L.; Shek, C. H. Appl. Phys. Lett. 2006, 88, 033115. (74) Dai, Z. R.; Gole, J. L.; Stout, J. D.; Wang, Z. L. J. Phys. Chem. B 2002, 106, 1274. (75) Wang, Z. L. Adv. Mater. 2003, 15, 432. (76) Ng, H. T.; Li, J.; Smith, M. K.; Nguyen, P.; Han, A.; Cassell, J.; Meyyappan, M. Science 2003, 300, 1249. (77) Cheng, B.; Russell, J. M.; Shi, W. S.; Zhang, L.; Samulski, E. T. J. Am. Chem. Soc. 2004, 126, 5972. (78) Wu, Q. H.; Li, J.; Sun, S. G. Curr. Nanosci. 2010, 6, 525. (79) Mikko, U.; Hanna, L.; Heli, V.; Lauri, N.; Resch, R.; Gernot, F. Mikrochim. Acta 2000, 133, 119. (80) Hidalgo-Falla, P.; Peres, H. E.; Gouvêa, M. D.; RamirezFernandez, F. J. Mater. Sci. Forum 2005, 636, 498. (81) Oyabu, T. J. Appl. Phys. 1982, 53, 2785. (82) Chen, Z. W.; Lai, J. K. L.; Shek, C. H. Appl. Phys. Lett. 2006, 89, 231902. (83) Demarne, V.; Grisel, A. Sens. Actuators 1988, 13, 301. (84) Santos, O. D.; Weiller, M. L.; Junior, D. Q.; Medina, A. N. Sens. Actutaors B 2001, 75, 83. (85) Shukla, S.; Seal, S.; Ludwing, L.; Parish, C. Sens. Actutaors B 2004, 97, 256. (86) Harbeck, S.; Szatvanyi, A.; Barsan, N.; Weimar, U.; Hoffmann, V. Thin Solid Films 2003, 436, 76. (87) Windischmann, H.; Mark, P. J. Electrochem. Soc. 1979, 126, 627. (88) Lantto, V.; Romppainen, P. Surf. Sci. 1987, 192, 243. (89) Romppainen, P.; Lantto, V. J. Appl. Phys. 1988, 63, 5159. (90) Shek, C. H.; Lai, J. K. L.; Lin, G. M. J. Phys. Chem. Solids 1999, 60, 189. (91) Chen, Z. W.; Zhang, S. Y.; Tan, S.; Hou, J. G.; Zhang, Y. H. Thin Solid Films 1998, 322, 194. (92) Feder, J. Fractal; Plenum Press: New York, 1988. (93) Chen, Z. W.; Zhang, S. Y.; Tan, S.; Hou, J. G.; Wu, Z. Q. Appl. Phys. A: Mater. Sci. Process. 2003, 76, 33. (94) Chen, Z. W.; Zhang, S. Y.; Tan, S.; Hou, J. G. Mater. Res. Bull. 2002, 37, 825. (95) Chen, Z. W.; Tan, S.; Zhang, S. Y.; Hou, J. G. Int. J. Mod. Phys. B 2002, 16, 159. (96) Chen, Z. W.; Zhang, S. Y.; Tan, S.; Hou, J. G.; Zhang, Y. H.; Sekine, H. J. Appl. Phys. 2001, 89, 783. (97) Fallah, H. R.; Ghasemi, M.; Hassanzadeh, A. Phys. E 2007, 39, 23. (98) Chen, Z. W.; Lai, J. K. L.; Shek, C. H.; Chen, H. D. Appl. Phys. A: Mater. Sci. Process. 2005, 81, 959. (99) Kim, Y. J.; Kim, Y. T.; Yang, H. K.; Park, J. C.; Han, J. I.; Lee, Y. E.; Kim, H. J. J. Vac. Sci. Technol. A 1997, 15, 1103. (100) Chen, Z. W.; Shek, C. H.; Lai, J. K. L. Phys. B 2005, 358, 56. (101) Chen, Z. W.; Lai, J. K. L.; Shek, C. H.; Chen, H. D. Appl. Surf. Sci. 2005, 250, 3. (102) Li, Q. B.; Chen, C.; Chen, Z. W.; Jiao, Z.; Wu, M. H.; Shek, C. H.; Wu, C. M. L.; Lai, J. K. L. Inorg. Chem. 2012, 51, 8473. (103) Chen, Z. W.; Zhang, S. Y.; Tan, S.; Hou, J. G.; Wu, Z. Q. Appl. Phys. A: Mater. Sci. Process. 2004, 78, 603. (104) Chen, Z. W.; Lai, J. K. L.; Shek, C. H.; Chen, H. D. J. Phys. D: Appl. Phys. 2004, 37, 2726. (105) Zhu, S.; Lu, Y. F.; Hong, M. H.; Chen, X. Y. J. Appl. Phys. 2001, 89, 2400. (106) Kim, D.; Lee, H. J. Appl. Phys. 2001, 89, 5703. (107) Zhu, S.; Lu, Y. F.; Hong, M. H. Appl. Phys. Lett. 2001, 79, 1396. (108) Berthe, L.; Fabbro, R.; Peyre, P.; Tollier, L.; Bartnicki, E. J. Appl. Phys. 1997, 82, 2826. (109) Penn, R. L.; Banfield, J. F. Geochim. Cosmochim. Acta 1999, 63, 1549.

(33) Liu, J. P.; Li, Y. Y.; Huang, X. T.; Ding, R. M.; Hu, Y. Y.; Jiang, J.; Liao, L. J. Mater. Chem. 2009, 19, 1859. (34) Wang, W. F.; Chen, Z. W.; Hou, L. G.; Hu, P. F.; Shek, C. H.; Wu, C. M. L.; Lai, J. K. L. J. Phys. Chem. C 2013, 117, 8903. (35) Chen, Z. W.; Li, Q. B.; Wang, J.; Pan, D. Y.; Jiao, Z.; Wu, M. H.; Shek, C. H.; Wu, C. M. L.; Lai, J. K. L. Inorg. Chem. 2011, 50, 6756. (36) Chen, C.; Ding, G. J.; Zhang, D.; Jiao, Z.; Wu, M. H.; Shek, C. H.; Wu, C. M. L.; Lai, J. K. L.; Chen, Z. W. Nanoscale 2012, 4, 2590. (37) Chen, Z. W.; Lai, J. K. L.; Shek, C. H. J. Chem. Phys. 2006, 124, 184707. (38) Chen, Z. W.; Tan, S.; Zhang, S. Y.; Hou, J. G.; Wu, Z. Q.; Sekine, H. Japn. J. Appl. Phys. 2001, 40, 3960. (39) Li, Q. B.; Wang, J.; Jiao, Z.; Wu, M. H.; Shek, C. H.; Wu, C. M. L.; Lai, J. K. L.; Chen, Z. W. Chaos Solitons Fractals 2011, 44, 640. (40) Chen, Z. W.; Lai, J. K. L.; Shek, C. H. J. Phys. D: Appl. Phys. 2006, 39, 4544. (41) Velásquez, C.; Rojas, F.; Esparza, J. M.; Ortiz, A.; Campero, A. J. Phys. Chem. B 2006, 110, 11832. (42) Yadav, B. C.; Verma, N.; Singh, S. Opt. Laser Technol. 2012, 44, 1681. (43) Thangaraju, B. Thin Solid Films 2002, 402, 71. (44) Ohodnicki, P. R., Jr.; Natesakhawat, S.; Baltrus, J. P.; Howard, B.; Brown, T. D. Thin Solid Films 2012, 520, 6243. (45) Acharya, R.; Zhang, Y. Q.; Cao, X. A. Thin Solid Films 2012, 520, 6130. (46) Liu, H. D.; Huang, J. M.; Li, X. L.; Liu, J.; Zhang, Y. X. Ceram. Int. 2012, 38, 5145. (47) Khuc, Q. T.; Vu, X. H.; Dang, D. V.; Nguyen, D. C. Adv. Nat. Sci.: Nanosci. Nanotechnol. 2010, 1, 025010. (48) He, H., Jr.; Wu, T. H.; Hsin, C. L.; Li, K. M.; Chen, L. J.; Chueh, Y. L.; Chou, L. J.; Wang, Z. L. Small 2006, 1, 116. (49) Niu, M.; Cheng, Y.; Wang, Y.; Cui, L.; Bao, F.; Zhou, L. Cryst. Growth Des. 2008, 8, 1727. (50) Liu, Y.; Dong, Y.; Wang, G. Appl. Phys. Lett. 2003, 82, 260. (51) Zhang, D. F.; Sun, L. D.; Jia, C. J.; Yan, Z. G.; You, L. P.; Yan, C. H. J. Am. Chem. Soc. 2005, 127, 13492. (52) Budak, S.; Miao, G. X.; Ozdemir, M.; Chetry, K. B.; Gupta, A. J. Cryst. Growth 2006, 291, 405. (53) Mathur, S.; Barth, S. Small 2007, 3, 2070. (54) Dou, X. C.; Sabba, D.; Mathews, N.; Wong, L. H.; Lam, Y. M.; Mhaisalkar, S. Chem. Mater. 2011, 23, 3938. (55) Li, H. X.; Ma, H. Q.; Zeng, Y. P.; Pan, A. L.; Zhang, Q. L.; Yu, H. C.; Wang, T. H.; Wang, Y. G.; Zou, B. S. J. Phys. Chem. C 2010, 114, 1844. (56) Chen, Z. W.; Jiao, Z.; Wu, M. H.; Shek, C. H.; Wu, C. M. L.; Lai, J. K. L. Prog. Mater. Sci. 2011, 56, 901. (57) Varghese, O. K.; Malhotra, L. K. Sens. Actuators B 1998, 53, 19. (58) He, Y. S.; Campbell, J. C.; Murphy, R. C.; Arendt, M. F.; Swinnea, J. S. J. Mater. Res. 1993, 8, 3131. (59) Wang, D. Z.; Wen, S. L.; Chen, J.; Zhang, S. Y.; Li, F. Q. Phys. Rev. B 1994, 49, 14282. (60) Chen, Z. W.; Pan, D. Y.; Zhao, B.; Ding, G. J.; Jiao, Z.; Wu, M. H.; Shek, C. H.; Wu, C. M. L.; Lai, J. K. L. ACS Nano 2010, 4, 1202. (61) Chen, Z. W.; Zhang, H. J.; Li, Z.; Jiao, Z.; Wu, M. H.; Shek, C. H.; Wu, C. M. L.; Lai, J. K. L. Acta Mater. 2009, 57, 5078. (62) Wu, H. B.; Chen, J. S.; Lou, X. W. D.; Hng, H. H. J. Phys. Chem. C 2011, 115, 24605. (63) Liu, J. P.; Li, Y. Y.; Huang, X. T.; Ding, R. M.; Hu, Y. Y.; Jiang, J.; Liao, L. J. Mater. Chem. 2009, 19, 1895. (64) Kar, A.; Kundu, S.; Patra, A. J. Phys. Chem. C 2011, 115, 118. (65) Jang, H. S.; Kang, S. O.; Kim, Y. I. Solid State Commun. 2006, 140, 495. (66) Ng, M. F.; Shen, L.; Zhou, L.; Yang, S. W.; Tan, V. B. C. Nano Lett. 2008, 8, 3662. (67) Ye, J. F.; Qi, L. M. J. Mater. Sci. Technol. 2008, 24, 529. (68) Ji, X. X.; Huang, X. T. Nanoscale Res. Lett. 2010, 5, 649. (69) Chen, Z. W.; Liu, G.; Zhang, H. J.; Ding, G. J.; Jiao, Z.; Wu, M. H.; Shek, C. H.; Wu, C. M. L.; Lai, J. K. L. J. Non-Cryst. Solids 2009, 355, 2647. 7481


Chemical Reviews

Review

(110) Penn, R. L.; Banfield, J. F. Science 1998, 281, 969. (111) Hou, J. G.; Wu, Z. Q. Phys. Rev. B 1989, 40, 1008. (112) Chen, Z. W.; Zhang, S. Y.; Tan, S.; Hou, J. G.; Zhang, Y. H. J. Vac. Sci. Technol. A 1998, 16, 2292. (113) Cooper, R. B.; Advani, G. N.; Jordan, A. G. J. Electron. Mater. 1981, 10, 455. (114) Cicera, A.; Dieguez, A.; Diaz, R.; Cornet, A.; Morante, J. R. Sens. Actuators B 1999, 58, 360. (115) Chopra, K. L.; Major, S.; Pandya, D. K. Thin Solid Films 1983, 102, 1. (116) Abello, L.; Bochu, B.; Gaskov, A.; Koudryavtseva, S.; Lucazeau, G.; Roumyantseva, M. J. Solid State Chem. 1998, 135, 78. (117) Ansari, S. G.; Boroojerdian, P.; Sainkar, S. R.; Karekar, R. N.; Aiyer, R. C.; Kulkarni, S. K. Thin Solid Films 1997, 295, 271. (118) Chen, Z. W.; Wu, C. M. L.; Shek, C. H.; Lai, J. K. L.; Jiao, Z.; Wu, M. H. Crit. Rev. Solid State Mater. Sci. 2008, 33, 197. (119) Lee, S. W.; Kim, Y. W.; Chen, H. D. Appl. Phys. Lett. 2001, 78, 350. (120) Mason, M. G.; Hung, L. S.; Tang, C. W.; Lee, S. T.; Wong, K. W.; Wang, M. J. Appl. Phys. 1999, 86, 1688. (121) Göpel, W.; Schierbaum, K. D. Sens. Actuators B 1995, 26, 1. (122) Madou, M. J.; Morrison, S. R. Chemical sensing with solid state devices; Academic Press: New York, 1989. (123) Sung, J. H.; Lee, Y. S.; Lim, J. W.; Hong, Y. H.; Lee, D. D. Sens. Actuators B 2000, 66, 149. (124) Kim, C. K.; Choi, S. M.; Noh, I. H.; Lee, J. H.; Hong, C.; Chae, H. B.; Jang, G. E.; Park, H. D. Sens. Actuators B 2001, 77, 463. (125) El Khakani, M. A.; Dolbec, R.; Serventi, A. M.; Horrillo, M. C.; Trudeau, M.; Saint-Jacques, R. G.; Rickerby, D. G.; Sayago, I. Sens. Actuators B 2001, 77, 383. (126) Dolbec, R.; El Khakani, M. A.; Serventi, A. M.; Trudeau, M.; Saint-Jacques, R. G. Thin Solid Films 2002, 419, 230. (127) Chen, Z. W.; Lai, J. K. L.; Shek, C. H.; Chen, H. D. J. Mater. Res. 2003, 18, 1289. (128) Kim, T. W.; Lee, D. U.; Lee, J. H.; Choo, D. C.; Jung, M.; Yoon, Y. S. J. Appl. Phys. 2001, 90, 175. (129) Mandayo, G. G.; Castano, E.; Gracia, F. J.; Cirera, A.; Cornet, A.; Morante, J. R. Sens. Actuators B 2003, 95, 90. (130) Ogawa, H.; Nishikawa, M.; Abe, A. J. Appl. Phys. 1982, 53, 4448. (131) Katti, V. R.; Debnath, A. K.; Muthe, K. P.; Kaur, M.; Dua, A. K.; Gadkari, S. C.; Gupta, S. K.; Sahni, V. C. Sens. Actuators B 2003, 96, 245. (132) Min, B. K.; Choi, S. D. Sens. Actuators B 2004, 98, 239. (133) Korotcenkov, G.; Brinzari, V.; Schwank, J.; Cerneavschi, A. Mater. Sci. Eng., C 2002, 19, 73. (134) Alfonso, C.; Charai, A.; Armigliato, A.; Narducci, D. Appl. Phys. Lett. 1996, 68, 1207. (135) Shukla, S.; Patil, S.; Kuiry, S. C.; Rahman, Z.; Du, T.; Ludwig, L.; Parish, C.; Seal, S. Sens. Actuators B 2003, 96, 343. (136) Kotsikau, D.; Ivanovskaya, M.; Orlik, D.; Falasconi, M. Sens. Actuators B 2004, 101, 199. (137) Zhang, G.; Liu, M. L. Sens. Actuators B 2000, 69, 144. (138) Korotcenkov, G.; Brinzari, V.; Golovanov, V.; Blinov, Y. Sens. Actuators B 2004, 98, 41. (139) Korotcenkov, G.; Macsanov, V.; Tolstoy, V.; Brinzari, V.; Schwank, J.; Faglia, D. Sens. Actuators B 2003, 96, 602. (140) Xu, C. N.; Tamaki, J.; Miura, N.; Yamazoe, N. Sens. Actuators B 1991, 3, 147. (141) Sberveglieri, G. Sens. Actuators B 1992, 6, 239. (142) Chrisey, D. B.; Hubler, G. K. Pulsed laser deposition of thin films; Wiley: New York, 1994. (143) Willmott, P. R.; Huber, J. R. Rev. Mod. Phys. 2000, 72, 315. (144) Auciello, O.; Engemann, J. Multicomponent and multilayered thin films for advanced microtechnologies: techniques, fundamentals and devices; Kluwer: Netherlands, 1993. (145) Bäuerle, D. Laser processing and chemistry; Springer: New York, 1996.

(146) Von Allmen, M.; Blatter, A. Laser-beam interactions with materials; Springer: New York, 1995. (147) Dolbec, R.; El Khakani, M. A.; Serventi, A. M.; Saint-Jacques, R. G. Sens. Actuators B 2003, 93, 566. (148) Petrik, P.; Biró, L. P.; Fried, M.; Lohner, T.; Berger, R.; Schneider, C.; Gyulai, J.; Ryssel, H. Thin Solid Films 1998, 315, 186. (149) McCarthy, G.; Welton, J. J. Powder Diffraction 1989, 4, 156. (150) Chen, Z. W.; Lai, J. K. L.; Shek, C. H.; Chen, H. D. Appl. Phys. A: Mater. Sci. Process. 2005, 81, 1073. (151) Duparre, A. Handbook of optical properties; CRC Press: Boca Raton, FL, 1995; Vol.1. (152) Senthilkumar, M.; Sahoo, N. K.; Thakur, S.; Tokas, R. B. Appl. Sur. Sci. 2005, 245, 114. (153) Lindström, T.; Isidorsson, J.; Niklasson, G. A. Thin Solid Films 2001, 401, 165. (154) Barsan, N.; Schweizer-Berberich, M.; Göpel, W. Fresenius’ J. Anal. Chem. 1999, 365, 287. (155) Gong, J.; Chen, Q.; Fei, W.; Seal, S. Sens. Actuators B 2004, 102, 117. (156) Amalric-Popescu, D.; Bozon-Verduraz, F. Catal. Today 2001, 70, 139. (157) Zhang, Y.; Kolmakov, A.; Lilach, Y.; Moskovits, M. J. Phys. Chem. B 2005, 109, 1923. (158) Nicholas, C. P.; Marks, T. J. Nano Lett. 2004, 4, 1557. (159) Stjerna, B.; Olsson, E.; Granqvist, C. G. J. Appl. Phys. 1994, 76, 3797. (160) Kay, A.; Grätzel, M. Chem. Mater. 2002, 14, 2930. (161) Arias, A. C.; De Lima, J. R.; Hümmelgen, I. A. Adv. Mater. 1998, 10, 392. (162) Andersson, A.; Johannsson, N.; Bröms, P.; Yu, N.; Lupo, D.; Salaneck, W. R. Adv. Mater. 1998, 10, 859. (163) Yang, P. D.; Zhao, D. Y.; Margolese, D. I.; Chmelka, B. F.; Stucky, G. D. Chem. Mater. 1999, 11, 2813. (164) Severin, K. G.; Abdel-Fattah, T. M.; Pinnavaia, T. J. Chem. Commun. 1998, 1471. (165) Hyodo, T.; Abe, S.; Shimizu, Y.; Egashira, M. Sens. Actuators B 2003, 93, 590. (166) Fujihara, S.; Maeda, T.; Ohgi, H.; Hosono, E.; Imai, H.; Kim, S. H. Langmuir 2004, 20, 6476. (167) Toupance, T.; Babot, O.; Jousseaume, B.; Vilacüa, G. Chem. Mater. 2003, 15, 4691. (168) Takenaka, S.; Takahashi, R.; Sato, S.; Sodesawa, T.; Matsumoto, F.; Yoshida, S. Microporous Mesoporous Mater. 2003, 59, 123. (169) Pan, J. H.; Chai, S. Y.; Lee, C.; Park, S. E.; Lee, W. I. J. Phys. Chem. C 2007, 111, 5582. (170) Liu, Y. K.; Zheng, C. L.; Wang, W. Z.; Yin, C. R.; Wang, G. H. Adv. Mater. 2001, 13, 1883. (171) Xu, C. K.; Xu, G. D.; Liu, Y. K.; Zhao, X. L.; Wang, G. H. Scripta Mater. 2002, 46, 789. (172) Dai, Z. R.; Pan, Z. W.; Wang, Z. L. Solid State Commun. 2001, 118, 351. (173) Hu, J. Q.; Ma, X. L.; Shang, N. G.; Xie, Z. Y.; Wong, N. B.; Lee, C. S.; Lee, S. T. J. Phys. Chem. B 2002, 106, 3823. (174) Maddalena, A.; Maschio, R. D.; Dire, S.; Raccanelli, A. J. J. NonCryst. Solids 1990, 121, 365. (175) Shek, C. H.; Lai, J. K. L.; Lin, G. M. NanoStuct. Mater. 1999, 11, 887. (176) Ghostagore, R. N. J. Electrochem. Soc. 1978, 125, 110. (177) Tarey, R. D.; Raju, T. A. Thin Solid Films 1995, 128, 181. (178) Minami, T.; Nanto, H.; Takata, S. J. J. Appl. Phys. 1988, 27, L287. (179) Zhu, J. J.; Lu, Z. H.; Aruna, S. T.; Aurbach, D.; Gedanken, A. Chem. Mater. 2000, 12, 2557. (180) Schosser, V.; Wind, G. Proceedings of the 8th EC Photovoltaic Solar Energy Conference; Florence: Italy, 1998. (181) Serventi, A. M.; Dolbec, R.; El Khakani, M. A.; Saint-Jacques, R. G.; Rickerby, D. G. J. Phys. Chem. Solids 2003, 64, 2097. 7482


Chemical Reviews

Review

(215) Reddy, S. A.; Figueiredo, N. M.; Cavaleiro, A. Appl. Surf. Sci. 2012, 258, 8902. (216) Li, H.; Wang, J.; Liu, H.; Zhang, H.; Li, X. J. Cryst. Growth 2005, 275, e943. (217) Bouhssira, N.; Abed, S.; Tomasella, E.; Cellier, J.; Mosbah, A.; Aida, M. S.; Jacquet, M. Appl. Surf. Sci. 2006, 252, 5594. (218) Hambergend, I.; Granquist, C. G. J. Appl. Phys. 1986, 60, R123. (219) He, G.; Bhat, G.; Chen, Z. W. J. Alloys Compd. 2011, 509, 9513. (220) Bhatti, M. T.; Rana, A. M.; Khan, A. F. Mater. Chem. Phys. 2004, 84, 126. (221) Olguín, D.; Cardona, M.; Cantarero, A. Solid State Commun. 2002, 122, 575. (222) Das, S.; Kar, S.; Chaudhuri, S. J. Appl. Phys. 2006, 99, 114303. (223) Peaker, A. R.; Horsley, B. Rev. Sci. Instrument 1971, 42, 1825. (224) Von Ortenberg, M.; Link, J.; Helbig, R. J. Opt. Soc. Am. 1977, 67, 968. (225) Bucher, E. Appl. Phys. 1978, 17, 1. (226) Ghosh, A. K.; Fishman, C.; Feng, T. J. Appl. Phys. 1978, 49, 3490. (227) Watson, J. Sens. Actuators 1984, 5, 29. (228) Feng, X.; Ma, J.; Yang, F.; Ji, F.; Zong, F.; Luan, C.; Ma, H. Mater. Lett. 2008, 62, 1779. (229) Mukashev, B. N.; Tokmoldin, S. Z.; Beisenkhanov, N. B.; Kikkarin, S. M.; Valitova, I. V.; Glazman, V. B.; Aimagambetov, A. B.; Dmitrieva, E. A.; Veremenithev, B. M. Mater. Sci. Eng., B 2005, 118, 164. (230) Wu, Q. H.; Song, J.; Kang, J.; Dong, Q. F.; Wu, S. T.; Sun, S. G. Mater. Lett. 2007, 61, 3679. (231) Du, J.; Zhang, H. J.; Jiao, Z.; Wu, M. H.; Shek, C. H.; Wu, C. M. L.; Lai, J. K. L.; Chen, Z. W. J. Nanosci. Nanotechnol. 2011, 11, 10659. (232) Lee, J. Thin Solid Films 2008, 516, 1386. (233) Chen, Z. W.; Zhang, S. Y.; Tan, S.; Hou, J. G.; Zhang, Y. H.; Sekine, H. J. Appl. Phys. 2001, 89, 783. (234) Ku, D. Y.; Kim, I. H.; Lee, I.; Lee, K. S.; Lee, T. S.; Jeong, J. H.; Cheong, B.; Baik, Y. J.; Kim, W. M. Thin Solid Films 2006, 515, 1364. (235) Chen, D.; Xu, J.; Xie, Z.; Shen, G. Z. ACS Appl. Mater. Interfaces 2011, 3, 2112. (236) Tsokkou, D.; Othonos, A.; Zervos, M. Appl. Phys. Lett. 2012, 100, 133101. (237) Liu, J.; Chen, X. L.; Wang, W. J.; Song, B.; Huang, Q. S. Cryst. Growth Des. 2009, 9, 1757. (238) Kuang, Q.; Lao, C. S.; Wang, Z. L.; Xie, Z. X.; Zheng, L. S. J. Am. Chem. Soc. 2007, 129, 6070. (239) Meduri, P.; Pendyala, C.; Kumar, V.; Sumanasekera, G. U.; Sunkara, M. K. Nano Lett. 2009, 9, 612. (240) Liu, J.; Hu, Y. J.; Gu, F.; Ma, J.; Li, C. Z. Ind. Eng. Chem. Res. 2011, 50, 5584. (241) Wang, B.; Zhu, L. F.; Yang, Y. H.; Xu, N. S.; Yang, G. W. J. Phys. Chem. C 2008, 112, 6643. (242) Park, J. Y.; Choi, S. W.; Kim, S. S. J. Phys. Chem. C 2011, 115, 12774. (243) Cheng, C. W.; Liu, B.; Yang, H. Y.; Zhou, W. W.; Sun, L.; Chen, R.; Yu, S. F.; Zhang, J. X.; Gong, H.; Sun, H. D.; Fan, H. J. ACS Nano 2009, 3, 3069. (244) Pan, J.; Song, X. F.; Zhang, J.; Shen, H.; Xiong, Q. H. J. Phys. Chem. C 2011, 115, 22225. (245) Woo, H. S.; Hwang, I. S.; Na, C. W.; Kim, S. J.; Choi, J. K.; Lee, J. S.; Choi, J.; Kim, G. T.; Lee, J. H. Mater. Lett. 2012, 68, 60. (246) Lee, S. Y.; Shin, Y. H.; Kim, Y.; Kim, S.; Ju, S. J. Lumin. 2011, 131, 2565. (247) Thanasanvorakun, S.; Mangkorntong, P.; Choopun, S.; Mangkorntong, N. Ceram. Int. 2008, 34, 1127. (248) Duraia, E. S. M. A.; Mansorov, Z. A.; Tokmolden, S. Phys. B 2009, 404, 3952. (249) Braun, P. V.; Osenar, P.; Stupp, S. I. Nature 1996, 380, 325. (250) Zhang, Y.; Suenaga, K.; Colliex, C.; Lijima, S. Science 1998, 281, 973.

(182) Serrini, P.; Briois, V.; Horrillo, M. C.; Traverse, A.; Manes, L. Thin Solid Films 1997, 304, 113. (183) Xu, C.; Tamaki, J.; Miura, N.; Yamazoe, N. J. Mater. Sci. Lett. 1989, 8, 1092. (184) Chen, Z. W.; Lai, J. K. L.; Shek, C. H. Chem. Phys. Lett. 2006, 422, 1. (185) Chen, Z. W.; Lai, J. K. L.; Shek, C. H. Solid State Chem. 2005, 178, 892. (186) Nolsson, G.; Nelin, G. Phys. Rev. B 1972, 6, 3777. (187) Zhang, S. L.; Zhu, B. F.; Huang, F. M.; Yan, Y.; Shang, E. Y.; Fan, S. S.; Han, W. G. Solid State Commun. 1999, 111, 647. (188) Bouzerar, G.; Ziman, T. Phys. Rev. Lett. 2006, 96, 207602. (189) Rahman, G.; Garcia-Suarez, V. M.; Hong, S. C. Phys. Rev. B 2008, 78, 184404. (190) Wang, H.; Yan, Y.; Li, K.; Du, X.; Lan, Z.; Jin, H. Phys. Status Solidi B 2010, 247, 444. (191) Ogale, S. B.; Choudhary, R. J.; Buban, J. P.; Lofland, S. E.; Shinde, S. R.; Kale, S. N.; Kulkarni, V. N.; Higgins, J.; Lanci, C.; Simpson, J. R.; Browning, N. D.; Das Sarma, S.; Drew, H. D.; Greene, R. L.; Venkatesan, T. Phys. Rev. Lett. 2003, 91, 077205. (192) Kiliç, Ç .; Zunger, A. Phys. Rev. Lett. 2002, 88, 095501. (193) Mohan Kant, K.; Chandrasekaran, K.; Ogale, S. B.; Venkatesan, T.; Rao, M. S. R. J. Appl. Phys. 2005, 97, 10A925. (194) Ghosh, S.; Khan, G. G.; Mandal, K. ACS Appl. Mater. Interfaces 2012, 4, 2048. (195) Herng, T. S.; Wong, M. F.; Qi, D.; Yi, J.; Kumar, A.; Huang, A. F.; Kartawidjaja, C.; Smadici, S.; Abbamonte, P.; Sánchez-Hanke, C.; Shannigrahi, S.; Xue, J. M.; Wang, J.; Feng, Y. P.; Rusydi, A.; Zeng, K.; Ding, J. Adv. Mater. 2011, 23, 1635. (196) Philip, J.; Punnoose, A.; Kim, B. I.; Reddy, K. M.; Layne, S.; Holmes, J. O.; Satpati, B.; Leclair, P. R.; Santos, T. S.; Moodera, J. S. Nat. Mater. 2006, 5, 298. (197) Jeon, H. C.; Jeong, Y. S.; Kang, T. W.; Kim, T. W.; Chung, K. J.; Chung, K. J.; Jhe, W.; Song, S. A. Adv. Mater. 2002, 14, 1725. (198) Venkatesan, M.; Fitzgerald, C. B.; Coey, J. M. D. Nature 2004, 430, 630. (199) Sundaresan, A.; Bhargavi, R.; Rangarajan, N.; Siddesh, U.; Rao, C. N. R. Phys. Rev. B 2006, 74, 161306. (200) Ghosh, S.; Khan, G. G.; Das, B.; Mandal, K. J. Appl. Phys. 2011, 109, 123927. (201) Ney, A.; Ollefs, K.; Ye, S.; Kammermeier, T.; Ney, V.; Kaspar, T. C.; Chambers, S. A.; Wilhelm, F.; Rogalev, A. Phys. Rev. Lett. 2008, 100, 157201. (202) Xu, Q.; Zhou, S.; Marko, D.; Potzger, K.; Fassbender, J.; Vinnichenko, M.; Helm, M.; Hochmuth, H.; Lorenz, M.; Grundmann, M.; Schmidt, H. J. Phys. D: Appl. Phys. 2009, 42, 085001. (203) Barla, A.; Schmerber, G.; Beaurepaire, E.; Dinia, A.; Bieber, H.; Colis, S.; Scheurer, F.; Kappler, J. P.; Imperia, P.; Nolting, F.; Wilhelm, F.; Rogalev, A.; Müller, D.; Grob, J. J. Phys. Rev. B 2007, 76, 125201. (204) Ghosh, S.; De Munshi, D.; Mandal, K. J. Appl. Phys. 2010, 107, 123919. (205) Fernandes, V.; Schio, P.; de Oliveira, A. J. A.; Schreiner, W. H.; Varalda, J.; Mosca, D. H. J. Appl. Phys. 2011, 110, 113902. (206) Ghosh, S.; Mandal, K. J. Magn. Magn. Mater. 2010, 322, 1979. (207) Dakhel, A. A.; El-Hilo, M. J. Appl. Phys. 2010, 107, 123905. (208) Sakamoto, N. J. Phys. Soc. Jpn. 1962, 17, 99. (209) Kumagai, H.; Oka, Y.; Kawata, S.; Ohba, M.; Inoue, K.; Kurmoo, M.; Okawa, H. Polyhedron 2003, 22, 1917. (210) Presley, R. E.; Munsee, C. L.; Park, C. H.; Hong, D.; Wager, J. F.; Keszler, D. A. J. Phys. D: Appl. Phys. 2004, 37, 2810. (211) Law, M.; Kind, H.; Messer, B.; Kim, F.; Yang, P. D. Angew. Chem., Int. Ed. 2002, 41, 2405. (212) Nayral, C.; Viala, E.; Fau, P.; Senocq, F.; Jumas, J. C.; Maisonnat, A.; Chaudret, B. Chem.Eur. J. 2000, 6, 4082. (213) De Monredon, S.; Cellot, A.; Ribot, F.; Sanchez, C.; Armelao, L.; Gueneau, L.; Delattre, L. J. Mater. Chem. 2002, 12, 2396. (214) Cao, Y. A.; Zhang, X. T.; Yang, W. S.; Du, H.; Bai, Y. B.; Li, T. J.; Yao, J. N. Chem. Mater. 2000, 12, 3445. 7483


Chemical Reviews

Review

(251) Zhang, Y.; Ichihashi, T.; Landree, E.; Nihey, F.; Lijima, S. Science 1999, 285, 1719. (252) Peng, X.; Manna, L.; Yang, W.; Wickham, J.; Scher, E.; Kadavanich, A.; Alivisatos, A. P. Nature 2000, 404, 59. (253) Duan, X. F.; Huang, Y.; Cui, Y.; Wang, J. F.; Lieber, C. M. Nature 2001, 409, 66. (254) Peng, H. Y.; Pan, Z. W.; Xu, L.; Fan, X. H.; Wang, N. B.; Lee, C. S.; Lee, S. T. Adv. Mater. 2001, 13, 317. (255) Murray, C. B.; Cagan, C. R.; Bawendi, M. G. Science 1995, 270, 1335. (256) Leite, E. R.; Weber, I. T.; Longo, E.; Varela, J. A. Adv. Mater. 2000, 12, 965. (257) Postma, H. W.; Teepen, T.; Yao, Z.; Grifoni, M.; Deckker, C. Science 2001, 293, 76. (258) Gordon, R. G. MRS Bull. 2000, 25, 52. (259) Li, C.; Zhang, D. H.; Han, S.; Liu, X. L.; Tang, T.; Zhou, C. W. Adv. Mater. 2003, 15, 143. (260) Idota, Y.; Kubota, T.; Matsufuji, A.; Maekawa, Y.; Miyasaka, T. Science 1997, 276, 1395. (261) Liu, Z. Q.; Zhang, D. H.; Han, S.; Li, C.; Tang, T.; Jin, W.; Liu, X. L.; Lei, B.; Zhou, C. W. Adv. Mater. 2003, 15, 1754. (262) Zhang, D. F.; Sun, L. D.; Yin, J. L.; Yan, C. H. Adv. Mater. 2003, 15, 1022. (263) Chen, Z. W.; Lai, J. K. L.; Shek, C. H. J. Non-Cryst. Solid 2006, 352, 3285. (264) Chen, Z. W.; Zhang, S. Y.; Tan, S.; Wang, J.; Jin, S. Z. Appl. Phys. A: Mater. Sci. Process. 2004, 78, 581. (265) Chen, Z. W.; Lai, J. K. L.; Shek, C. H. Scripta Mater. 2006, 55, 735. (266) Chen, Z. W.; Shek, C. H.; Lai, J. K. L. Appl. Phys. A: Mater. Sci. Process. 2005, 80, 703. (267) Chen, Z. W.; Lai, J. K. L.; Shek, C. H. Appl. Phys. Lett. 2005, 86, 181911. (268) Huang, M. H.; Wu, Y.; Feick, H.; Tran, N.; Weber, E.; Yang, P. Adv. Mater. 2001, 13, 113. (269) Leite, E. R.; Gomes, J. W.; Oliveira, M. M.; Lee, E. J. H.; Longo, E.; Varela, J. A.; Paskocimas, C. A.; Boschi, T. M.; Lanciotti, F., Jr.; Pizani, P. S.; Soares, P. C., Jr. Appl. Sci. Res. 2002, 2, 125. (270) Law, M.; Kind, H.; Messer, B.; Kim, F.; Yang, P. Angew. Chem. 2002, 114, 2511. (271) Liu, H. I.; Maluf, N. I.; Pease, R. F. W. J. Vac. Sci. Technol. B 1992, 10, 2846. (272) Ono, T.; Saitoh, H.; Esashi, M. Appl. Phys. Lett. 1997, 70, 1852. (273) Frank, F. C. Discovery Faraday 1949, 5, 48. (274) Wagner, R. S.; Ellis, W. C. Appl. Phys. Lett. 1964, 4, 89. (275) Chen, Z. W.; Jiao, Z.; Wu, M. H.; Shek, C. H.; Wu, C. M. L.; Lai, J. K. L. Mater. Chem. Phys. 2009, 115, 660. (276) Zheng, J. G.; Pan, X. Q.; Schweizer, M.; Weimar, U.; Göpel, W.; Rühle, M. Philos. Mag. Lett. 1996, 73, 93. (277) Cheng, Y.; Xiong, P.; Fields, L.; Zheng, J. P.; Yang, R. S.; Wang, Z. L. Appl. Phys. Lett. 2006, 89, 093114. (278) McDowell, M. G.; Sanderson, R. J.; Hill, I. G. Appl. Phys. Lett. 2008, 92, 013502. (279) Ng, H. T.; Han, J.; Yamada, T.; Nguyen, P.; Chen, Y. P.; Meyyappan, M. Nano Lett. 2004, 4, 1247. (280) Lai, E.; Kim, W.; Yang, P. Nano Res. 2008, 1, 123. (281) Law, M.; Greene, L. E.; Johnson, J. C.; Saykally, R.; Yang, P. Nat. Mater. 2005, 4, 455. (282) Feng, X.; Shankar, K.; Varghese, O. K.; Paulose, M.; Latempa, T. J.; Grimes, C. A. Nano Lett. 2008, 8, 3781. (283) Wang, Z. L.; Song, J. Science 2006, 312, 242. (284) Kim, W. S.; Kim, D. H.; Choi, K. J.; Park, J. G.; Hong, S. H. Cryst. Growth Des. 2010, 10, 4746. (285) Beltran, A.; Andres, J.; Longo, E.; Leite, E. R. Appl. Phys. Lett. 2003, 83, 635. (286) Li, Q.; Creighton, J. R.; Wang, G. T. J. Cryst. Growth 2008, 310, 3706. (287) Cui, Y.; Lauhon, L. J.; Gudiksen, M. S.; Wang, J.; Lieber, C. M. Appl. Phys. Lett. 2001, 78, 2214.

(288) Kuykendall, T.; Pauzauskie, P. J.; Zhang, Y.; Goldberger, J.; Sirbuly, D.; Denlinger, J.; Yang, P. D. Nat. Mater. 2004, 3, 524. (289) Nguyen, P.; Ng, H. T.; Kong, J.; Cassell, A. M.; Quinn, R.; Li, J.; Han, J.; McNeil, M.; Meyyappan, M. Nano Lett. 2003, 3, 925. (290) Leonardy, A.; Hung, W. Z.; Tsai, D. S.; Chou, C. C.; Huang, Y. S. Cryst. Grow. Des. 2009, 9, 3958. (291) Pan, J.; Shen, H.; Werner, U.; Prades, J. D.; HernandezRamirez, F.; Soldera, F.; Mücklich, F.; Mathur, S. J. Phys. Chem. C 2011, 115, 15191. (292) Mathur, S.; Barth, S.; Shen, H.; Pyun, J. C.; Werner, U. Small 2005, 1, 713. (293) Pan, J.; Xiao, L.; Shen, H.; Mathur, S. Ceram. Eng. Sci. Proc. 2010, 30, 9. (294) Uchic, M. D.; Holzer, L.; Inkson, B. J.; Principe, E. L.; Munroe, P. MRS Bull. 2007, 32, 408. (295) Munroe, P. R. Mater. Charact. 2009, 60, 2. (296) Ghosh, R.; Basak, D.; Fujihara, S. J. Appl. Phys. 2004, 96, 2689. (297) Schwenzer, B.; Gomm, J. R.; Morse, D. E. Langmuir 2006, 22, 9829. (298) Lee, H. Y.; Ko, H. J.; Yao, T. Appl. Phys. Lett. 2003, 82, 523. (299) Shi, W. S.; Agyeman, O.; Xu, C. N. J. Appl. Phys. 2002, 91, 5640. (300) Zhou, X. T.; Heigl, F.; Murphy, M. W.; Regier, T.; Coulthard, I.; Blyth, R. I. R.; Sham, T. K. Appl. Phys. Lett. 2006, 89, 213109. (301) Zhou, X. T.; Zhou, J. G.; Murphy, M. W.; Ko, J. Y. P.; Heigl, F.; Regier, T.; Blyth, R. I. R.; Sham, T. K. J. Chem. Phys. 2008, 128, 144703. (302) Wang, D. N.; Yang, J. L.; Li, X. F.; Wang, J. J.; Li, R. Y.; Cai, M.; Sham, T. K.; Sun, X. L. Cryst. Growth Des. 2012, 12, 397. (303) Sharma, S.; Sunkara, M. K.; Miranda, R.; Lian, G.; Dickey, E. C. Mater. Res. Soc. Symp. Proc. 2001, 676, Y1.6.1. (304) Dobrokhotov, V.; Mcllroy, D. N.; Norton, M. G.; Abuzir, A. J. Appl. Phys. 2006, 99, 104302. (305) Cabot, A.; Dieguez, A.; Romano-Rodriguez, A.; Morante, J. R.; Barsan, N. Sens. Actuators B 2001, 79, 98. (306) Mathur, S.; Ganesan, R.; Grobelsek, I.; Shen, H.; Ruegamer, T.; Barth, S. Adv. Eng. Mater. 2007, 9, 658. (307) Srivastava, R.; Dwivedi, R.; Srivastava, S. K. Microelectron. J. 1998, 29, 833. (308) Chaturvedi, A.; Mishra, V. N.; Dwivedi, R.; Srivastava, S. K. Microelectron. J. 2000, 31, 283. (309) Forleo, A.; Francioso, L.; Capone, S.; Casino, F.; Siciliano, P.; Tan, O. K.; Hui, H. Procedia Chem. 2009, 1, 196. (310) Hui, H.; Tan, O. K.; Lee, Y. C.; Tran, T. D.; Tse, M. S. Appl. Phys. Lett. 2005, 87, 163123. (311) Pan, J.; Ganesan, R.; Shen, H.; Mathur, S. J. Phys. Chem. C 2010, 114, 8245. (312) Mai, Y. J.; Wang, X. L.; Xiang, J. Y.; Qiao, Y. Q.; Zhang, D.; Gu, C. D.; Tu, J. P. Electrochim. Acta 2011, 56, 2306. (313) Zhu, J. W.; Zeng, G. Y.; Nie, F. D.; Xu, X. M.; Chen, S.; Han, Q. F.; Wang, X. Nanoscale 2010, 2, 988. (314) Li, N.; Wang, Z. Y.; Zhao, K. K.; Shi, Z. J.; Xu, S. K.; Gu, Z. N. J. Nanosci. Nanotechnol. 2010, 10, 6690. (315) Kim, H.; Seo, D. H.; Kim, S. W.; Kim, J.; Kang, K. Carbon 2011, 49, 326. (316) Lu, T.; Zhang, Y. P.; Li, H. B.; Pan, L. K.; Li, Y. L.; Sun, Z. Electrochim. Acta 2010, 55, 4170. (317) Kim, Y. J.; Lee, J. H.; Yi, G. C. Appl. Phys. Lett. 2009, 95, 213101. (318) Wu, J.; Shen, X. P.; Jiang, L.; Wang, K.; Chen, K. M. Appl. Surf. Sci. 2010, 256, 2826. (319) Wang, D. H.; Choi, D. W.; Li, J.; Yang, Z. G.; Nie, Z. M.; Kou, R.; Hu, D. H.; Wang, C. M.; Saraf, L. V.; Zhang, J. G.; Aksay, I. A.; Liu, J. ACS Nano 2009, 3, 907. (320) Lambert, T. N.; Chavez, C. A.; Sanchez, B. H.; Lu, P.; Bell, N. S.; Ambrosini, A.; Friedman, T.; Boyle, T. J.; Wheeler, D. R.; Huber, D. L. J. Phys. Chem. C 2009, 113, 19812. (321) Williams, G.; Seger, B.; Kamat, P. V. ACS Nano 2008, 2, 1487. 7484


Chemical Reviews

Review

(322) Zhang, X. Y.; Li, H. P.; Cui, X. L.; Lin, Y. H. J. Mater. Chem. 2010, 20, 2801. (323) Kim, S. R.; Parvez, M. K.; Chhowalla, M. Chem. Phys. Lett. 2009, 483, 124. (324) Yao, J.; Shen, X. P.; Wang, B.; Liu, H. K.; Wang, G. X. Electrochem. Commun. 2009, 11, 1849. (325) Lian, P.C.; Zhu, X. F.; Liang, S. Z.; Li, Z.; Yang, W. S.; Wang, H. H. Electrochim. Acta 2011, 56, 4532. (326) Paek, S. M.; Yoo, E. J.; Honma, I. Nano Lett. 2009, 9, 72. (327) Wang, X. Y.; Zhou, X. F.; Yao, K.; Zhang, J. G.; Liu, Z. P. Carbon 2011, 49, 133. (328) Zhang, L. S.; Jiang, L. Y.; Yan, H. J.; Wang, W. D.; Wang, W.; Song, W. G.; Guo, Y. G.; Wan, L. J. J. Mater. Chem. 2010, 20, 5462. (329) Du, Z. F.; Yin, X. M.; Zhang, M.; Hao, Q. Y.; Wang, Y. G.; Wang, T. H. Mater. Lett. 2010, 64, 2076. (330) Li, Y. M.; Lv, X. J.; Lu, J.; Li, J. H. J. Phys. Chem. C 2010, 114, 21770. (331) Wang, Z. Y.; Zhang, H.; Li, N.; Shi, Z. J.; Gu, Z. N.; Cao, G. P. Nano Res. 2010, 3, 748. (332) Li, F.; Song, J.; Yang, H.; Gan, S.; Zhang, Q.; Han, D.; Ivaska, A.; Niu, L. Nanotechnology 2009, 20, 455602. (333) Gaidi, M.; Hajjaji, A.; Smirani, R.; Bessais, B.; El Khakani, M. A. J. Appl. Phys. 2010, 108, 063537. (334) Jean, S. T.; Her, Y. C. J. Appl. Phys. 2009, 105, 024310. (335) Ding, J. J.; Yan, X. B.; Li, J.; Shen, B. S.; Yang, J.; Chen, J. T.; Xue, Q. J. ACS Appl. Mater. Interfaces 2011, 3, 4299. (336) Hwang, J. O.; Lee, D. H.; Kim, J. Y.; Han, T. H.; Kim, B. H.; Park, M.; No, K.; Kim, S. O. J. Mater. Chem. 2011, 21, 3432. (337) Kim, T. W.; Lee, D. U.; Yoon, Y. S. J. Appl. Phys. 2000, 88, 3759. (338) Wang, B.; Yang, Y. H.; Wang, C. X.; Xu, N. S.; Yang, G. W. J. Appl. Phys. 2005, 98, 124303. (339) Chen, Y. J.; Xue, X. Y.; Wang, Y. G.; Wang, T. H. Appl. Phys. Lett. 2006, 88, 083105. (340) Chen, S.; Wang, M.; Ye, J.; Cai, J.; Ma, Y.; Zhou, H.; Qi, L. Nano Res. 2013, 6, 243. (341) Bass, J. D.; Schaper, C. D.; Rettner, C. T.; Arellano, N.; Alharbi, F. H.; Miller, R. D.; Kim, H. C. ACS Nano 2011, 5, 4065. (342) Xi, G. C.; Ye, J. H. Inorg. Chem. 2010, 49, 2302. (343) Cheng, G.; Chen, J. Y.; Ke, H. Z.; Shang, J.; Chu, R. Mater. Lett. 2011, 65, 3327. (344) Chen, D. L.; Gao, L. Chem. Phys. Lett. 2004, 398, 201. (345) Zhang, H. L.; Hu, C. G. Catal. Commun. 2011, 14, 32. (346) Hodes, G. Adv. Mater. 2007, 19, 639. (347) Noda, S.; Chutinan, A.; Imada, M. Nature 2000, 407, 608. (348) Noda, S.; Yamamoto, N.; Imada, M.; Kobayashi, H.; Okano, M. IEEE J. Lightwave Technol. 1999, 17, 1948. (349) Alivisatos, P. Science 1996, 271, 933. (350) Lauhon, L. H.; Gudiksen, M. S.; Wang, D.; Lieber, C. M. Nature 2002, 420, 57. (351) Li, J.; Zhao, D.; Meng, X.; Zhang, Z.; Zhang, J.; Shen, D.; Lu, Y.; Fan, X. J. Phys. Chem. B 2006, 110, 14685. (352) Shi, L.; Xu, Y.; Hark, S.; Liu, Y.; Wang, S.; Peng, L. M.; Wong, K.; Li, Q. Nano Lett. 2007, 7, 3559. (353) Shimpi, P.; Gao, P. X.; Goberman, D. G.; Ding, Y. Nanotechnology 2009, 20, 125608. (354) Richter, J. P.; Voss, T.; Kim, D. S.; Scholz, R.; Zacharias, M. Nanotechnology 2008, 19, 305202. (355) Jin, C.; Kim, H.; Lee, W. I.; Lee, C. Adv. Mater. 2011, 23, 1982. (356) Kordas, G. J. Mater. Chem. 2000, 10, 115. (357) Dekker, C. Phys. Today 1999, 52, 22. (358) Morales, A. M.; Lieber, C. M. Science 1998, 279, 208. (359) Duan, X.; Lieber, C. M. J. Am. Chem. Soc. 2000, 122, 188. (360) Wu, Y.; Yang, P. Chem. Mater. 2000, 12, 605. (361) Ginley, D. S.; Bright, C. Mater. Res. Soc. Bull. 2000, 25, 15. (362) Yamazoe, N. Sens. Actuators B 1991, 5, 7. (363) Yoon, K. H.; Cho, Y. S.; Kang, D. H. J. Mater. Sci. 1998, 33, 2977.

(364) Wada, H.; Sakane, K.; Kitamura, T. J. Mater. Sci. Lett. 1991, 10, 1076. (365) Hashimoto, S.; Yamaguchi, A. J. Eur. Ceram. Soc. 2000, 20, 397. (366) Kajiwara, M. J. Mater. Sci. 1987, 22, 1223. (367) Tas, A. C. J. Am. Ceram. Soc. 2001, 84, 295. (368) Hayashi, Y.; Kimura, T.; Yamaguchi, T. J. Mater. Sci. 1986, 21, 757. (369) Chen, Z. W.; Du, J.; Zhang, H. J.; Jiao, Z.; Wu, M. H.; Shek, C. H.; Wu, C. M. L.; Lai, J. K. L. Acta Mater. 2009, 57, 4632. (370) Ocana, M.; Serna, C. J. Spectrochim. Acta, Part A 1991, 47, 765. (371) Korotcenkov, G. Mater. Sci. Eng., B 2007, 139, 1. (372) Vayssieres, L.; Graetzel, M. Angew. Chew. Int. Ed. 2004, 43, 3666. (373) Wang, Y. L.; Jiang, X. C.; Xia, Y. N. J. Am. Chem. Soc. 2003, 125, 16176. (374) Kolmakov, A.; Zhang, Y.; Cheng, G.; Moskovits, M. Adv. Mater. 2003, 15, 997. (375) Wang, Y.; Lee, J. Y.; Zeng, H. C. Chem. Mater. 2005, 17, 3899. (376) Hu, J.; Bando, Y.; Liu, Q.; Golberg, D. Adv. Funct. Mater. 2003, 13, 493. (377) Wang, C.; Zhou, Y.; Ge, M. Y.; Xu, X. B.; Zhang, Z. L.; Jiang, J. Z. J. Am. Chem. Soc. 2010, 132, 46. (378) Ho, S. Y.; Wong, A. S. W.; Ho, G. W. Cryst. Growth Des. 2009, 9, 732. (379) Korotcenkov, G.; Brinzari, V.; Boris, Y.; Ivanov, M.; Schwank, J.; Morante, J. Thin Solid Films 2003, 436, 119. (380) Joshi, R. K.; Kruis, F. E. Appl. Phys. Lett. 2006, 89, 153116. (381) Epifani, M.; Arbiol, J.; Pellicer, E.; Comini, E.; Siciliano, P.; Faglia, G.; Morante, J. R. Cryst. Growth Des. 2008, 8, 1774. (382) Wang, Y. D.; Djerdj, I.; Antonietti, M.; Smarsly, B. Small 2008, 4, 1656. (383) Wang, J.; Zhang, P.; Qi, J. Q.; Yao, P. J. Sens. Actuators B 2009, 136, 399. (384) Sun, P.; Yu, Y. S.; Xu, J.; Sun, Y. F.; Ma, J.; Lu, G. Y. Sens. Actuators B 2011, 160, 244. (385) Tricoli, A.; Graf, M.; Pratsinis, S. E. Adv. Funct. Mater. 2008, 18, 1969. (386) Liu, J.; Gu, F.; Hu, Y. J.; Li, C. Z. J. Phys. Chem. C 2010, 114, 5867. (387) Shang, G. L.; Wu, J. H.; Tang, S.; Liu, L.; Zhang, X. P. J. Phys. Chem. C 2013, 117, 4345. (388) Gratzel, M. Nature 2001, 414, 338. (389) Kang, S. H.; Choi, S. H.; Kang, M. S.; Kim, J. Y.; Kim, H. S.; Hyeon, T.; Sung, Y. E. Adv. Mater. 2008, 20, 54. (390) Yang, S.; Hou, Y.; Zhang, B.; Yang, X. H.; Fang, W. Q.; Zhao, H. J.; Yang, H. G. J. Mater. Chem. A 2013, 1, 1374. (391) Choi, J. A.; Kim, S. H.; Kim, D. W. J. Power Sources 2010, 195, 6192. (392) Kumai, Y.; Shirai, S.; Sudo, E.; Seki, J.; Okamoto, H.; Sugiyama, Y.; Nakano, H. J. Power Sources 2011, 196, 1503. (393) Li, Y.; Tan, B.; Wu, Y. Nano Lett. 2008, 8, 265. (394) Cai, Y.; Liu, S.; Yin, X. M.; Hao, Q. Y.; Zhang, M.; Wang, T. H. Phys. E 2010, 43, 70. (395) Cui, Z. M.; Hang, L. Y.; Song, W. G.; Guo, Y. G. Chem. Mater. 2009, 21, 1162. (396) Huang, X. H.; Xia, X. H.; Yuan, Y. F.; Zhou, F. Electrochim. Acta 2011, 56, 4960. (397) Li, L. M.; Yin, X. M.; Liu, S. A.; Wang, Y. G.; Chen, L. B.; Wang, T. H. Electrochem. Commun. 2010, 12, 1383. (398) Pfanzelt, M.; Kubiak, P.; Fleischhammer, M.; WohlfahrtMehrens, M. J. Power Sources 2011, 196, 6815. (399) Guo, P.; Song, H.; Chen, X. Electrochem. Commun. 2009, 11, 1320. (400) Wu, Z. S.; Ren, W.; Wen, L.; Gao, L.; Zhao, J.; Chen, Z.; Zhou, G.; Li, F.; Cheng, H. M. ACS Nano 2010, 4, 3187. (401) Xing, L.; Cui, C.; Ma, C.; Xue, X. Mater. Lett. 2011, 65, 2104. (402) Zhou, G.; Wang, D. W.; Li, F.; Zhang, L.; Li, N.; Wu, Z. S.; Wen, L.; Lu, G. Q.; Cheng, H. M. Chem. Mater. 2010, 22, 5306. 7485


Chemical Reviews

Review

(403) Wang, F.; Yao, G.; Xu, M.; Zhao, M.; Sun, Z.; Song, X. J. Alloys Compd. 2011, 509, 5969. (404) Wu, P.; Du, N.; Zhang, H.; Yu, J. X.; Qi, Y.; Yang, D. R. Nanoscale 2011, 3, 746. (405) Kim, J. G.; Nam, S. H.; Lee, S. H.; Choi, S. M.; Kim, W. B. ACS Appl. Mater. Interfaces 2011, 3, 828. (406) Lupan, O.; Chow, L.; Chai, G.; Heinrich, H.; Park, S.; Schulte, A. Phys. E 2009, 41, 533. (407) Lupan, O.; Chow, L.; Chai, G.; Schulte, A.; Park, S.; Heinrich, H. Mater. Sci. Eng., B 2009, 157, 101. (408) Vayssieres, L.; Graetzel, M. Angew. Chem., Int. Ed. 2004, 116, 3752. (409) Park, M. S.; Wang, G. X.; Kang, Y. M.; Wexler, D.; Dou, S. X.; Liu, H. K. Angew. Chem., Int. Ed. 2007, 119, 764. (410) Wang, Y.; Lee, J. Y. J. Phys. Chem. B 2004, 108, 17832. (411) Huggins, R. A. Ionics 1997, 3, 245. (412) Ying, Z.; Wan, Q.; Cao, H.; Song, Z. T.; Feng, S. L. Appl. Phys. Lett. 2005, 87, 113108. (413) Poizot, P.; Laruelle, S.; Grugeon, S.; Dupont, L.; Tarascon, J. M. Nature 2000, 407, 496. (414) Hu, Y. S.; Rezan, D. C.; Titirici, M. M.; Müller, J. O.; Schlögl, R.; Antonietti, M.; Maier, J. Angew. Chem., Int. Ed. 2008, 47, 1645. (415) Reddy, M. V.; Yu, T.; Sow, C. H.; Shen, Z. X.; Lim, C. T.; Rao, G. V. S.; Chowdari, B. V. R. Adv. Funct. Mater. 2007, 17, 2792. (416) Wang, Y.; Lee, J. Y. J. Power Sources 2005, 144, 220. (417) Zhang, W. M.; Hu, J. S.; Guo, Y. G.; Zheng, S. F.; Zhong, L. S.; Song, W. G.; Wan, L. J. Adv. Mater. 2008, 20, 1160. (418) Persaud, K.; Dodd, G. Nature 1982, 299, 352.

7486


Recent Advances in Tin Dioxide Materials: Some ... - ACS Publications

Recommend Documents