Large-Scale Porous Hematite Nanorod Arrays: Direct Growth on

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Large-Scale Porous Hematite Nanorod Arrays: Direct Growth on Titanium Foil and Reversible Lithium Storage Yuqing Song, Shanshan Qin, Yangwei Zhang, Wanqin Gao, and Jinping Liu* Institute of Nanoscience and Nanotechnology, Department of Physics, Huazhong Normal UniVersity, Wuhan 430079, P. R. China ReceiVed: August 11, 2010

Porous single-crystalline hematite (R-Fe2O3) nanorod array has been synthesized on large-area Ti foil via a facile hydrothermal method followed by a simple annealing treatment in Ar gas at 450 °C. The nanorods attained from 6 h hydrothermal reaction are average 30 nm in diameter and 450 nm in length. When used directly as additive-free anode for lithium ion batteries (LIBs), the R-Fe2O3 nanorod array demonstrates excellent cycling performance up to 50 times (∼562 mAh g-1 retained at C/5) and good rate capability, in distinct contrast to R-Fe2O3 nanorod powder-based electrode. The improved electrochemical performance could be ascribed to the enhanced electron transport and Li+ diffusion that result from the well-defined array architecture and the porous nature of the single-crystalline nanorods. Fe3O4 and C/R-Fe2O3 nanorod arrays are further prepared to improve the lithium storage property. Our work represents a successful example of fabricating iron oxide 1D nanostructure arrays directly on nonreactive current collector. Once optimized, the array electrode may hold great promise in thin-film LIBs and other microelectronic systems. Introduction During the past few decades, considerable attention has been focused on the development of new energy resources and materials because of the increasingly serious energy shortage and environment pollution.1 With its relatively high energy density, long cycle life, and environmental friendliness, the lithium-ion battery (LIB) has become one of the most attractive energy storage devices and has come into commercialization for various portable electronic devices including cellular phones, lap-top computers, and so forth. Some future industrial needs such as plug-in hybrid electric vehicles (PHEVs),1-3 however, generally require both high-energy and high-power densities. In addition, the storage of more energy at high power on a small footprint area is also highly desirable for microsystems and flexible thin-film devices.2b These requirements have triggered a great deal of interest in the search for alternative anode materials for commercial graphitic carbon that has a theoretical capacity of only 372 mAh g-1.4-6 Transition metal oxides, discovered with capacities much higher (two or three times) than that of graphite and the volume change during discharge-charge processes not so great as for alloy anodes, hold great promise as electrode materials for reversible lithium storage.7,8 Among them, hematite (R-Fe2O3) has been extensively investigated as a potential anode material candidate due to its low cost and high resistance to corrosion.9-12 The Fe2O3 crystal lattice can react with 6 Li ions per formula unit, exhibiting theoretical capacity as high as 1007 mAh g-1. The overall electrochemical reaction of Li+ intercalation/ deintercalation in Fe2O3 materials can be described by the following equation: Fe2O3 + 6Li+ rf 3Li2O + 2Fe3+.10,11 The quite reversible reaction ensures excellent columbic efficiency and high reversible capacity of R-Fe2O3. Bulk oxide electrode materials generally suffer from poor kinetics and serious capacity fade upon cycling, even at low * To whom correspondence should be addressed: Fax +86-02767861185. E-mail: [email protected].

rates.13 By contrast, nanosized materials can possess distinct electrochemical properties and exhibit improved lithium storage performance arising from the large surface-to-volume ratio and small dimensions.6,14 Nevertheless, further improvement of the cycle life and rate capability is still very challenging. Previous reports provide clear evidence that 1D nanostructured materials can be widely applied in optical, electrical, and optoelectronic devices due to their unique geometry.15-18 Compared to conventional bulk materials and nanoparticles, 1D nanostructures have the advantage of high aspect ratio,19-22 which can facilitate the alleviation of the mechanical stress induced by volume change during repeated charge-discharge cycles.14 Moreover, the 1D structures, especially when aligned directly on current collectors, can provide numerous conduction channels that direct the transport of electrons. In our previous work, some 1D metal oxide nanostructure arrays including ZnO nanowire,23 SnO2 nanorod,24 CoO nanowire,25 and Fe2O3 nanotube26 arrays have already been investigated and demonstrated to be a promising architecture for the electrochemical electrodes. Obviously, using 1D R-Fe2O3 nanostructure arrays as an anode is a feasible approach to improve the battery performance. Substantial efforts have been previously devoted to synthesize 1D R-Fe2O3 nanostructures.6,19,26-28 Once fabricated, however, R-Fe2O3 nanostructures need to be mixed with a polymer binder and acetylene black to form a slurry that is subsequently pasted on an electrode substrate; in the meantime, many undesirable interfaces are developed in the active material.29,30 Thus, not only does this electrode fabrication process negate the benefits of electrochemistry using 1D nanostructures but also it degrades the battery rate performance as well. To our best knowledge, there have been few reports on the direct growth of R-Fe2O3 1D nanostructure arrays so far,26,31-36 and no reports of growing R-Fe2O3 nanorod/wire array on conductive metal foils for use in LIBs have been demonstrated. It remains a huge challenge to develop a simple approach to synthesize array electrodes of R-Fe2O3 nanorods/wires with neither organic binder nor additive. Addressing this challenge will facilitate our in-depth under-

10.1021/jp1091009  2010 American Chemical Society Published on Web 11/15/2010

Large-Scale Porous Hematite Nanorod Arrays standing of the influence of the ordered architecture of transition metal oxide nanorods/wires on the electrochemical Li storage properties. Herein, we report a facile hydrothermal method utilizing lowcost precursor of FeCl3 followed by a simple annealing, to directly synthesize R-Fe2O3 nanorod array on titanium (Ti) foil and investigate its application as anode for LIBs. The largescale realization of FeOOH nanorod array is first presented. After being annealed in Ar gas, the as-grown FeOOH nanorods can be readily transferred into porous single-crystalline R-Fe2O3 nanorods while maintaining the array architecture unchanged. When used directly as an LIB anode, the R-Fe2O3 array shows significantly improved lithium storage performance as compared to the commercial graphite and nanostructured R-Fe2O3 powders. Our work provides a successful example of growing 1D R-Fe2O3 nanostructure arrays directly on conductive metal substrate, which have potential applications in other thin-film electrochemical devices. Experimental Section In a typical experiment, two aqueous solutions containing 0.946 g of FeCl3 · 6H2O and 0.479 g of Na2SO4 respectively were mixed with vigorous stirring. Distilled water was then added to obtain a final volume of 70 mL. After thorough magnetic stirring for 10 min, with a clean Ti substrate placed against the wall the mixture was transferred into a Teflon-lined stainless-steel autoclave and hydrothermally treated at 120 °C for 6 h. It should be pointed out that our synthesis avoided the additional HCl introduction,34 making it more environmentally friendly and easier to be manipulated. Na2SO4 was employed as a structure-directing agent to facilitate the relatively uniform growth of 1D nanorods. After cooling down to room temperature, the array-covered substrate was rinsed several times with distilled water and then dried at 60 °C. The array has a strong adhesion to the substrate even after ultrasonication for 15 min. The precipitate in solution was also collected and washed sequentially by distilled water and ethanol and then dried in a vacuum oven. Finally, annealing in Ar gas at 450 °C was carried out for 3 h to attain both R-Fe2O3 nanorod array and powder. Fe3O4 nanorod array was obtained by annealing the hydrothermally grown array product in H2 at 350 °C for 3 h.26 C/R-Fe2O3 nanorod array was synthesized based on the protocol reported in our previous work.23a Briefly, the as-grown R-Fe2O3 array was soaked in 0.02 M glucose solution for 15 h and further annealed in Ar gas at 450 °C for 5 h. The microstructure and morphology of the products were directly subjected to powder X-ray diffraction (XRD) (Bruker D-8 Avance) measurement, transmission electron microscopy(TEM) (JEM-2010FEF, 200 kV) observation, and scanning electron microscopy (SEM) (JSM-6700F, 5.0 kV) characterization equipped with an energy-dispersive X-ray spectrometer (EDS). Thermogravimetric analysis (TGA) was carried out on an SDT600 apparatus with a heating rate of 10 °C min-1 in N2. X-ray photoelectron spectroscopy (XPS) measurement were performed on a PerkinElmer PHI 5000C ECSA system with monochromatic Al Ka (1486.6 eV) irradiation. The mass of electrode materials was measured on BS 124 S Balance (Max: 120 g, d ) 0.1 mg). The substrate supporting R-Fe2O3 nanorod arrays was cut into a disk with a surface area of 1.539 cm2 and directly used as the working anode (14 mm diameter, R-Fe2O3 arrays on one side of the foil was removed for electrical contacting beforehand) in Swagelok-type battery which was assembled in an Ar-filled glovebox (Mbraun, Unilab, Germany). A Li-metal circular foil (0.59 mm thick, 14 mm diameter) as

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Figure 1. XRD pattern of the tetragonal FeOOH nanorod array. The standard bar pattern of pure Ti substrate was also provided.

the counter and reference electrodes, a microporous polypropylene membrane as the separator, and 1 M solution of LiPF6 dissolved in a mixed solution of ethylene carbonate (EC) and diethyl carbonate (DEC) (1:1 by volume) as the electrolyte. For the comparative study of the electrochemical performance, R-Fe2O3 disordered nanorod-based electrode was also prepared. The random nanorods were derived from the precipitation in the solution, which consists of R-Fe2O3 nanorods with similar morphology to that in the array. The electrode was made in the conventional way: a slurry was first obtained by thoroughly mixing 75 wt % R-Fe2O3 disordered nanorods, 15 wt % carbon black, and 10 wt % polyvinylidene fluoride (PVDF) in N-methyl-2-pyrrolidene (NMP) solvent. The slurry was then spread on the Ti surface at room temperature and further heated under vacuum overnight at 100 °C to remove the water or organic solvent. Before used, the film electrode was cut into circular disks of 14 mm in diameter and pressed between two stainless-steel plates at 20 MPa. The cells were aged for 12 h before measurement. The discharge-charge cycling was performed at room temperature by using a multichannel battery tester (model SCN, USA). Results and Discussion To identify the structure and chemical composition of the as-grown array, XRD and XPS measurements were conducted. Figure 1 shows XRD pattern of the product on Ti foil obtained after the hydrothermal treatment for 6 h. With the exception of the reflection from the substrate (labeled with star symbols), the characteristic peaks in XRD can be well indexed as tetragonal FeOOH (JCPDS Card No.75-1594). In addition, the intensity ratio of (002)/(200) peaks increases significantly with respect to that in the standard pattern, revealing a substantial texture effect in accordance with the crystal shape anisotropy and 〈001〉 orientation. For XPS, the binding energy scale was referenced to the characteristic carbon 1s binding energy of 285.0 eV (part a of Figure 2). Part b of Figure 2 shows the Fe 2p XPS spectrum of the array after subtraction of the spectrum background. The 2p photoelectron peaks appear around 712.0 and 725.0 eV with a shakeup satellite line at 719.6 eV, characteristic of Fe3+ in FeOOH. The line shape and binding energies of Fe 2p agree well with the literature values for FeOOH.37,38 Part c of Figure 2 presents the deconvoluted O 1s spectrum, where three peaks located at 532.2, 529.2, and 533.0 eV can be observed, which correspond to O2- in hydroxide (the second O in FeOOH) and oxide (the first O in FeOOH) lattices and to adsorbed oxygen, respectively.39 Top-view and tilted SEM images of the product are further shown in parts a and b of Figure 3, which clearly demonstrate

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Figure 2. XPS spectra of tetragonal FeOOH array: (a) wide range spectrum, (b) Fe 2p spectrum, (c) O 1s spectrum.

Figure 3. SEM images of FeOOH nanorod arrays obtained after 6 h (a,b) and 11 h reaction (c,d). Insert in part a of Figure 3 shows an optical image of the as-grown product. Insert in part d of Figure 3 is a cross-sectional image.

the growth of vertically aligned FeOOH nanorod array on Ti foil. The nanorods typically have sharp tips with average diameter and length of ∼30 and 450 nm, respectively. As shown in the insert of part a of Figure 3, the array can grow uniformly on a ∼9 cm2 alloy substrate. The array formation was initiated by heterogeneous nucleation. In this regard, nuclei were formed directly on the substrate under the hydrothermal conditions because of the lower interfacial nucleation energy on the substrate.23,24 As long as the precursor concentration is appropriate, external crystal growth would take place from these nuclei along the easy direction of crystallization, thus the singlecrystalline nanorods perpendicular to the substrate will be generated. For tetragonal FeOOH, the anisotropic growth along [001] direction may be attributed to the relatively high surface energy of (002) facets. The FeOOH nanorods, however, were partially agglomerated into bundles, probably owing to their high surface energy resulting from the small size. Thus, adjacent nanorods running in the same direction tend to combine with each other to reduce the surface energy with the maximum contact between the surfaces. More obvious bundlelike nanostructures can be achieved via tuning the solution concentration or reaction time. It can be seen from parts c and d of Figure 3 that more nanorod bundles dominate in the array when the reaction time was prolonged to 11 h. In addition, the length of the nanorods (the array thickness) increases to nearly one micrometer, as shown in the insert of part d of Figure 3. However, increasing the array thickness while keeping the rod diameter unchanged is difficult at the present time. We hope to achieve this by introducing surfactants into the reaction solution in future. Considering that small diameter of the nanorod is highly required to reduce Li+ diffusion pathway, which will be

necessary to achieve robust electrochemical kinetics, we chose the array attained after six hour growth for the following discussion. The formation of FeOOH was further confirmed by TGA and differential thermal analysis (DTA), as illustrated in part a of Figure 4. The FeOOH mainly undergoes the weight loss in two steps during the temperature increase process. Associated with this loss, two characteristic temperatures (T1, 110 °C; T2, 250 °C) can be seen in its TGA curve. The first weak loss of weight at 110 °C may be ascribed to the emission of physically absorbed water molecules; and the second striking weight decrease centered at 250 °C corresponds to the decomposition of FeOOH, which is a typical endothermic reaction. The weight loss ceases at ∼550 °C, afterward the mass remains the same. The stable residue can reasonably be attributed to the thermodynamically stable phase of iron oxide - R-Fe2O3, in accordance with previous results.40 Accordingly, the observed weight loss of physically absorbed water and the weight loss with regard to decomposition are ca. 1.06% and 9.85% (expected weight loss ) 10.11%), respectively. Both the XRD pattern (part c of Figure 4) of the product obtained after being annealed in Ar at 450 °C for 3 h and the color change from yellow green to red brown confirm the formation of R-Fe2O3. Part b of Figure 4 reveals that the 1D nanostructure is stable and the nanorod arrays are well preserved during the thermal treatment, even though the composition has been converted. The EDS result shown in part d of Figure 4 further indicates the annealed array contains only Fe and O elements with the molar ratio of ∼2:3. To provide further insights into the morphology of the generated arrays, TEM investigations were performed. Part a of Figure 5 shows the TEM image of initially formed FeOOH nanorods. After scraped from the substrate and ultrasonicated in ethanol, the sample dispersed on the copper grid still shows a rodlike morphology. The FeOOH nanorods have diameters of ∼30 nm and lengths of several hundred nanometers, in good agreement with SEM observation. The high-resolution TEM (HRTEM) image (part b of Figure 5) of an individual FeOOH nanorod (denoted by a circle in part a of Figure 5) demonstrates the single-crystal nature of the nanorod with an interplanar spacing of ca. 0.26 nm, consistent with the standard value for (021) plane of tetragonal FeOOH. The insert image in part b of Figure 5 is the corresponding fast Fourier transform (FFT) pattern. It can be indexed to single-crystal FeOOH with the growth orientation along [002]. Part c of Figure 5 displays the TEM image of R-Fe2O3 nanorods. The insert is the selectedarea electron diffraction (SAED) pattern of a single R-Fe2O3 nanorod, revealing that R-Fe2O3 nanorod is also single crystal and grows along the [10-10] direction. It is worth mentioning that some pores exist in the nanorods (highlighted by arrow in part d of Figure 5), which is reasonably attributed to dehydration and the lattice contraction occurred in the thermal decomposition

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Figure 4. (a) TGA and DTA curves of the FeOOH nanorod array. SEM image (b), XRD (c) and EDS analysis, (d) of R-Fe2O3 nanorod array.

Figure 6. Cyclic voltammograms of the electrode of R-Fe2O3 nanorod array at a scan rate of 1 mV/s.

Figure 5. (a) TEM image of several FeOOH nanorods. (b) HRTEM image of an individual nanorod. The insert is the corresponding FFT pattern. (c) TEM image of R-Fe2O3 nanorods. The insert is the SAED pattern taken from one nanorod denoted by a circle. (d) Enlarged TEM image showing the porous structure. Insert is the pore size distribution curve of R-Fe2O3 nanorods.

of FeOOH. BET pore size distribution measurement in part d of Figure 5 further demonstrates that R-Fe2O3 nanorods have the dominant pore sizes of 3.4 and 15 nm. The presence of such porous structures is highly desirable for electrochemical lithium storage.14b,26 Next, we study the electrochemical lithium storage of the porous single-crystalline R-Fe2O3 nanorod array. The electro-

chemical property of the array anode was first investigated by cyclic voltammetry (CV), as shown in Figure 6. When the electrode is scanned cathodically from 3.0 to 0.005 V in the first cycle at a scan rate of 1 mV/s at room temperature, two peaks, a tiny peak and an intensive one, are observed at ∼1.4 V (P1) and 0.45 V (P2), which can be attributed to the lithium insertion into the crystal structure of R-Fe2O3 (the formation of cubic Li2Fe2O3) and the reduction from Fe2+ to Fe0 accompanied with irreversible decomposition reaction of electrolyte, respectively.14c,e Meanwhile, in the first-cycle anodic process, a broad peak is present at ∼2.14 V (P3), corresponding to the reversible oxidation from Fe0 to Fe3+.6 During the subsequent cycles, the reduction peak P2 potential shifts to ∼0.65 V while the oxidation peak position is nearly unchanged. Importantly, the shape of CV curves in the following cycles remains similar to that of the second cycle with only a very small decrease in the integrated area, revealing the good reversibility of lithium storage of the nanorod array electrode.

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Figure 7. Charge-discharge curves of the R-Fe2O3 nanorod array.

Figure 8. Cycling performance of R-Fe2O3 nanorod array and R-Fe2O3 powder at C/5.

Charge-discharge tests were carried out to investigate the battery performance of the array as LIB anode. Figure 7 shows the charge-discharge curves of the R-Fe2O3 anode at various cycles, which were tested at a current rate of C/5 (about 134.2 mA g-1, 1C was defined as 4 Li+/h) within a voltage range of 0.005-3.0 V. The first discharge capacity is ∼1236 mAh g-1, whereas the first charge capacity is approximate 840 mAh g-1, giving a Coulombic efficiency of as high as ∼67.9%. After the first charge-discharge cycle, there is very little irreversible capacity loss. The relatively large irreversible capacity in the first cycle should be caused by the decomposition reaction of electrolyte and formation of the SEI (solid-electrolyte interphase) film onto the surface of R-Fe2O3. In the second, fifth and tenth cycles, although the discharge capacities decrease gradually to about 859, 771, and 747 mAh g-1 respectively, the Coulombic efficiencies are all above 98%. To demonstrate the advantages of R-Fe2O3 nanorod array electrode, anode fabricated from R-Fe2O3 nanorod powder was also tested for comparison. The cyclability of both R-Fe2O3 nanorod array and powder-based electrodes was tested via constant current (C/5) charge-discharge cycling up to 50 cycles. It is evident that the R-Fe2O3 nanorod array electrode exhibits much better cyclability, though the capacity decreases sharply during the first five cycles (Figure 8). After 50 cycles, the array electrode still delivers a discharge capacity of 562 mAh g-1, which is much higher than that of graphite (372 mAh g-1) and comparable to the results for previous purpose-designed R-Fe2O3 nanostructures.14b,19 By contrast, the discharge capacity of the R-Fe2O3 powder-based electrode decays dramatically during the first 15 cycles (from 1260 mAh g-1 to 402 mAh g-1). After the 50th cycling, the reversible capacity is only 329 mAh g-1,

Figure 9. (a) Rate performance of R-Fe2O3 nanorod array; (b) Cycling stability comparison of R-Fe2O3 nanorod array, Fe3O4 nanorod array, and C/R-Fe2O3 nanorod array at C/2. Charge, hollow symbol; Discharge, solid symbol.

even lower than that of graphite. It has been reported that, in addition to the intrinsic crystal structure, the lithium storage performance is greatly related to extrinsic morphology and assembly fashion of active materials, which have an obvious effect on the reactivity not only in the initial charge-discharge cycle but also in the subsequent cycles.2b,c One main reason for the superior electrochemical performance of R-Fe2O3 array is the array architecture consisting of individual nanorods growing directly on Ti foil, without additives and polymer binders. The direct alignment of single-crystalline nanorods on the current collector ensures many fast electron transport pathways.23-26,41-43 In addition, the interspaces between neighboring nanorods allow for better accommodation of volume changes while providing enough spaces for electrolyte penetration into the inter part of the array, which will significantly reduce the interfacial resistance.44,45 Another key aspect is the porosity of the R-Fe2O3 nanorods. The presence of porous structure enhances the contact area between materials and electrolyte, rendering an increased number of active sites for electrochemical lithium storage;46,47 the diffusion length of the Li+ is also greatly shortened in porous nanorods. We emphasize that this beneficial effect is more pronounced for ordered array because all of the pores can be exposed directly to electrolyte, in distinct contrast to R-Fe2O3 nanorod powder-based electrode, in which the contact of pores to electrolyte may be greatly hindered by additive and polymer binder. In addition, it should be pointed out that the array has a much smaller thickness as compared to traditional electrode film, which facilitates the transportation/diffusion of both electrons and Li+ and thus leads to better performance. However, to gain more energy per unit area, our future effort will be made to design an array with optimized thickness that can deliver both high energy and power density.

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Figure 10. (a) XRD pattern of the Fe3O4 nanorod array. (b) Raman spectrum of C/R-Fe2O3 nanorod array.

High rate performance is generally of great importance for various applications of batteries such as electric vehicles and portable powers.1,2 Benefiting from the novel porous singlecrystalline array architecture, good rate capability of the array electrode has been further demonstrated. Part a of Figure 9 displays rate performance of the R-Fe2O3 nanorod array electrode at 0.5, 1, and 2C rates for the first cycle. The charge capacity reaches 707 mAh g-1 at 0.5C current rate, with the Coulombic efficiency of 63.2%. At 1C rate, discharge capacity of 966 mAh g-1 and 605 mAh g-1 charge capacity can be maintained. More importantly, the reversible capacity at 2C rate is ∼459 mAh g-1, still much higher than that of graphite. As shown in part b of Figure 9, under 0.5C continuous cycling, the R-Fe2O3 nanorod array electrode still delivers discharge and charge capacities of 444 and 438 mAh g-1 after 50 cycles, respectively. The charge capacity retention after 50 cycles at 0.5C is ∼62.0%. To further explore the application potential of the directly grown array electrode, we have prepared Fe3O4 and C/R-Fe2O3 nanorod arrays from the precursor FeOOH array and tested their lithium storage properties under the same electrochemical conditions. These two arrays are well-known to have better electronic conductivity, thus this feature will facilitate the electron transport and the maintenance of electrical continuity in the active material. XRD result shown in part a of Figure 10 confirms the Fe3O4 array growth. In part b of Figure 10, in addition to the Raman peaks in the range of 200-700 cm-1 from R-Fe2O3, the G band and D band peaks of carbon can be clearly observed, indicating the successful carbon modification of pristine R-Fe2O3 array. The assembled cells from Fe3O4 and C/R-Fe2O3 nanorod array electrodes were cycled in 0.005-3.0 V range at room temperature with a current rate of 0.5C. Their cycling performances are illustrated in part b of Figure 9. In general, these two cells demonstrate obvious improved cycleability as compared with the cell assembled using pristine R-Fe2O3 array. Discharge capacities of 533 and 595 mAh g-1 can be retained after 50 times cycling for Fe3O4 and C/R-Fe2O3 nanorod array electrodes, respectively. The above results specifically indicate that the electronic conductivity plays an important role in the electrochemical property of electrode materials. Conclusions In summary, the simple synthesis and battery application of R-Fe2O3 nanorod array on Ti foil have been investigated. Not only the array architecture of 1D porous single-crystalline nanostructure but also the direct-growth method has remarkable effects on the electrochemical performance toward lithium storage. It is found that R-Fe2O3 nanorod array exhibits good

cycling performance and a high reversible capacity when used directly as LIB anode, which are much superior to R-Fe2O3 nanorod powder-based electrode. The lithium storage property of the nanorod array could be further improved by converting the active component to Fe3O4 or modification with conductive carbon. Our work indicates that there are good prospects for using directly grown iron oxide nanorod arrays as anode materials for thin-film LIBs. Acknowledgment. Financially supported by self-determined research funds of CCNU from the colleges’ basic research and operation of MOE (No.CCNU09A01019), the National Natural Science Foundation of China (No. 50872039), China Postdoctoral Science Foundation (20090460996), the Open Project Program of Key Laboratory of Quak & Lepton Physics (Huazhong Normal University), Ministry of Education, China(QLPL200902) and National Innovation Experiment Program for University Students (CCNU, 091051111). We are appreciative of valuable suggestions from the reviewers for revision of this paper. References and Notes (1) Armand, M.; Tarascon, J.-M. Nature 2008, 451, 652. (2) (a) Winter, M.; Besenhard, J. O.; Spahr, M. E.; Novak, P. AdV. Mater. 1998, 10, 725. (b) Bruce, P. G. ; Scrosati, B.; Tarascon, J.-M. Angew. Chem., Int. Ed. 2008, 47, 2930. (c) Larcher, D.; Beattie, S.; Morcrette, M.; Edstro¨m, K.; Jumas, J.-C.; Tarascon, J.-M. J. Mater. Chem. 2007, 17, 3759. (d) Xie, Y.; Wu, C. Z. Dalton Trans. 2007, 5235. (3) (a) Arico, A. S.; Bruce, P. G.; Scrosati, B.; Tarascon, J. M.; Van Schalkwijk, W. Nat. Mater. 2005, 4, 366. (b) Wang, Y.; Takahashi, K.; Lee, K. H.; Cao, G. Z. AdV. Funct. Mater. 2006, 16, 1133. (4) Liu, H.; Wang, G. X.; Wang, J. Z.; Wexler, D. Electrochem. Commun. 2008, 10, 1879. (5) NuLi, Y.; Zeng, R.; Zhang, P.; Guo, Z.; Liu, H. J. Power Sources 2008, 184, 456. (6) NuLi, Y.; Zeng, R.; Zhang, P.; Guo, Z.; Liu, H. J. Electrochem. Soc. 2008, 155, A196. (7) Poizot, P.; Laruelle, S.; Grugeon, S.; Dupont, L.; Tarascon, J. M. Nature 2000, 407, 496. (8) Yang, L. C.; Gao, Q. S.; Zhang, Y. H.; Tang, Y.; Wu, Y. P. Electrochem. Commun. 2008, 10, 118. (9) Gong, C.; Chen, D.; Jiao, X.; Wang, Q. J. Mater. Chem. 2002, 12, 1844. (10) Hosono, E.; Fujihara, S.; Honma, I.; Ichihara, M.; Zhou, H. J. Electrochem. Soc. 2006, 153, A1273. (11) Li, J.; Dahn, H. M.; Sanderson, R. J.; Todd, A. D. W.; Dahna, J. R. J. Electrochem. Soc. 2008, 155, A975. (12) Wu, Z. C.; Yu, K.; Zhang, S. D.; Xie, Y. J. Phys. Chem. C 2008, 112, 11307. (13) Li, H.; Huang, X.; Chen, L. Solid State Ionics 1999, 123, 189. (14) (a) Liu, S. L.; Zhang, L. N.; Zhou, J. P.; Xiang, J. F.; Sun, J. T.; Guan, J. G. Chem. Mater. 2008, 20, 3623. (b) Chen, J.; Xu, L.; Li, W.; Gou, X. AdV. Mater. 2005, 17, 582. (c) 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. (d) Zhang, W.-M.; Wu, X.-L.; Hu, J.-S.; Guo, Y.-

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G.; Wan, L.-J. AdV. Funct. Mater. 2008, 18, 3941. (e) Wu, X. L.; Guo, Y. G.; Wan, L. J.; Hu, C. W. J. Phys. Chem. C 2008, 112, 16824. (15) Patel, A. C.; Li, S.; Wang, C.; Zhang, W.; Wei, Y. Chem. Mater. 2007, 19, 1231. (16) Fan, H. J.; Werner, P.; Zacharias, M. Small 2006, 2, 700. (17) Goldberger, J.; He, R. R.; Zhang, Y. F.; Lee, S. W.; Yan, H. Q.; Choi, H. J.; Yang, P. D. Nature 2003, 422, 599. (18) Tenne, R. Nat. Nanotechnol. 2006, 1, 103. (19) (a) Liu, H.; Wexler, D.; Wang, G. X. J. Alloys Compd. 2009, 487, L24. (b) Zeng, S. Y.; Tang, K. B.; Li, T. W.; Liang, Z. H. J. Phys. Chem. C 2010, 114, 274. (20) Son, S. J.; Reihel, J.; He, B.; Schuchman, M.; Lee, S. B. J. Am. Chem. Soc. 2005, 127, 7136. (21) Jia, C.-J.; Sun, L.-D.; Yan, Z.-G.; You, L.-P.; Luo, F.; Han, X.-D.; Pang, Y.-C.; Zhang, Z.; Yan, C.-H. Angew. Chem., Int. Ed. 2005, 44, 4328. (22) Sui, Y. C.; Skomski, R.; Sorge, K. D.; Sellmyer, D. J. Appl. Phys. Lett. 2004, 84, 1525. (23) (a) Liu, J. P.; Li, Y. Y.; Ding, R. M.; Jiang, J.; Hu, Y. Y.; Ji, X. X.; Chi, Q. B.; Zhu, Z. H.; Huang, X. T. J. Phys. Chem. C 2009, 113, 5336. (b) Liu, J. P.; Huang, X. T.; Li, Y. Y.; Ji, X. X.; Li, Z. K.; He, X.; Sun, F. L. J. Phys. Chem. C 2007, 111, 4990. (24) 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. (25) Jiang, J.; Liu, J. P.; Ding, R. M.; Ji, X. X.; Hu, Y. Y.; Li, X.; Hu, A. Z.; Wu, F.; Zhu, Z. H.; Huang, X. T. J. Phys. Chem. C 2010, 114, 929. (26) Liu, J. P.; Li, Y. Y.; Fan, H. J.; Zhu, Z. H.; Jiang, J.; Ding, R. M.; Hu, Y. Y.; Huang, X. T. Chem. Mater. 2010, 22, 212. (27) Han, Q.; Liu, Z. H.; Xu, Y. Y.; Chen, Z. Y.; Wang, T. M.; Zhang, H. J. Phys. Chem. C 2007, 111, 5034. (28) Liu, H.; Wang, G. X.; Park, J.; Wang, J. Z.; Liu, H. K.; Zhang, C. Electrochim. Acta 2009, 54, 1733. (29) Li, J.; Dahn, H. M.; Krause, L. J.; Le, D. B.; Dahna, J. R. J. Electrochem. Soc. 2008, 155, A812. (30) Hanga, B. T.; Okadab, S.; Yamaki, J. J. Power Sources 2008, 178, 402.

Song et al. (31) (a) Yu, T.; Zhu, Y.; Xu, X.; Yeong, K.-S.; Shen, Z.; Chen, P.; Lim, C.-T.; Thong, J. T.-L.; Sow, C.-H. Small 2006, 2, 80–84. (b) Zhang, W. X.; Yang, S. H. Acc. Chem. Res. 2009, 42, 1617. (32) LaTempa, T. J.; Feng, X. J.; Paulose, M.; Grimes, C. A. J. Phys. Chem. C 2009, 113, 16293. (33) Mohapatra, S. K.; John, S. E.; Banerjee, S.; Misra, M. Chem. Mater. 2009, 21, 3048. (34) Vayssieres, L.; Beermann, N.; Lindquist, S.-E.; Hagfeldt, A. Chem. Mater. 2001, 13, 233. (35) Li, L.; Koshizaki, N. J. Mater. Chem. 2010, 20, 2972. (36) Bachmann, J.; Jing, J.; Knez, M.; Barth, S.; Shen, H.; Mathur, S.; Go¨sele, U.; Nielsch, K. J. Am. Chem. Soc. 2007, 129, 9554. (37) Mansour, A. N.; Brizzolara, R. A. Surf. Sci. Spectra 1998, 4, 357. (38) Sun, Z.; Yuan, H.; Liu, Z.; Han, B.; Zhang, X. AdV. Mater. 2005, 17, 2993. (39) Yan, D.; Yan, P. X.; Cheng, S.; Chen, J. T.; Zhuo, R. F.; Feng, J. J.; Zhang, G. A. Cryst. Growth Des. 2009, 9, 218. (40) Tang, B.; Wang, G. L.; Zhuo, L. H.; Ge, J. C.; Cui, L. J. Inorg. Chem. 2006, 45, 5196. (41) Liu, J. P.; Li, Y. Y.; Huang, X. T.; Li, G. Y.; Li, Z. K. AdV. Funct. Mater. 2008, 18, 1448. (42) Taberna, P. L.; Mitra, S.; Poizot, P.; Simon, P.; Tarascon, J. M. Nat. Mater. 2006, 5, 567. (43) Chan, C. K.; Peng, H.; Liu, G.; McIlwrath, K.; Zhang, X. F.; Huggins, R. A.; Cui, Y. Nat. Nanotechnol. 2008, 3, 31. (44) Li, Y.; Tan, B.; Wu, Y. Nano Lett. 2008, 8, 265. (45) Xiao, W.; Chen, J. S.; Li, C. M.; Xu, R.; Lou, X. W. Chem. Mater. 2010, 22, 746. (46) Chen, J. S.; Tan, Y. L.; Li, C. M.; Cheah, Y. L.; Luan, D.; Madhavi, S.; Boey, F. Y.; Archer, L. A.; Lou, X. W. J. Am. Chem. Soc. 2010, 132, 6124. (47) Wang, Y.; Su, F. B.; Lee, J. Y.; Zhao, X. S. Chem. Mater. 2006, 18, 1347.

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