New Insight into the Reaction Mechanism for Exceptional Capacity of

Oct 24, 2014 - Advanced Photon Source, Argonne National Laboratory, Argonne, Illinois 60439, United States. §. School of Advanced Materials Engineeri...
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New Insight into the Reaction Mechanism for Exceptional Capacity of Ordered Mesoporous SnO2 Electrodes via Synchrotron-Based X‑ray Analysis Hyunchul Kim,† Gwi Ok Park,† Yunok Kim,† Shoaib Muhammad,† Jaeseung Yoo,† Mahalingam Balasubramanian,‡ Yong-Hun Cho,§ Min-Gyu Kim,∥ Byungju Lee,⊥ Kisuk Kang,⊥ Hansu Kim,*,# Ji Man Kim,*,∇ and Won-Sub Yoon*,† †

Department of Energy Science and ∇Department of Chemistry, Sungkyunkwan University, Suwon 440-746, South Korea Advanced Photon Source, Argonne National Laboratory, Argonne, Illinois 60439, United States § School of Advanced Materials Engineering, Kookmin University, Seoul 136-702, South Korea ∥ Pohang Accelerator Beamline Research Division, Pohang 790-834, South Korea ⊥ Department of Materials Science and Engineering, Seoul National University, Seoul, 151-747, South Korea # Department of Energy Engineering, Hanyang University, Seoul 133-791, South Korea ‡

ABSTRACT: Tin oxide-based materials, operating via irreversible conversion and reversible alloying reaction, are promising lithium storage materials due to their higher capacity. Recent studies reported that nanostructured SnO2 anode provides higher capacity beyond theoretical capacity based on the alloying reaction mechanism; however, their exact mechanism remains still unclear. Here, we report the detailed lithium storage mechanism of an ordered mesoporous SnO2 electrode material. Synchrotron X-ray diffraction and absorption spectroscopy reveal that some portion of Li2O decomposes upon delithiation and the resulting oxygen reacts with Sn to form the SnOx phase along with dealloying of LixSn, which are the main reasons for unexpected high capacity of an ordered mesoporous SnO2 material. This finding will not only be helpful in a more complete understanding of the reaction mechanism of Sn-based oxide anode materials but also will offer valuable guidance for developing new anode materials with abnormal high capacity for next generation rechargeable batteries.

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SnO2 + 4Li+ + 4e− → Sn + 2Li 2O

INTRODUCTION

Lithium-ion batteries have been recognized as one of the most promising power source for various applications including portable electronics, electric vehicles, and power storage systems of renewable energy.1 Major challenges of lithium ion batteries for these applications include high energy density, excellent capacity retention, safety, and low cost.1−3 In order to achieve higher energy density of lithium ion battery than that of currently commercialized lithium ion battery,4−7 metal oxides are being investigated as alternative anode materials due to their high energy density achieved by conversion and alloying reactions.8,9 Especially, tin oxide-based materials, including SnO and SnO2, are being considered as one of the best anode materials due to their higher specific lithium storage capacities.10 Previous studies show that SnO2 goes through an irreversible conversion reaction during the initial cycle, which leads to formation of Sn metal and Li2O matrix, followed by a reversible alloying/dealloying reaction of Sn with lithium.11−17 © 2014 American Chemical Society

Sn + x Li+ + x e− → LixSn

(711 mAh/g)

(1)

(783 mAh/g) (0 ≪ x ≪ 4.4)

(2)

Reaction 1 is the main reason for the large initial irreversible capacity due to the formation of Li2O. Reaction 2 is responsible for reversible capacity of these electrode materials after the first discharge. The capacity fading in subsequent cycles is caused by large volume changes of Sn (200%) associated with its alloying/ dealloying reaction with lithium during cycling.18−21 A variety of tactics has been tried on tin oxide-based anode materials to solve the problems related with huge volume changes of Sn.22,23 Use of nanostructured tin oxide-based materials is regarded as one of the effective ways to achieve excellent capacity retention with high reversible capacity by absorbing volume-induced Received: July 14, 2014 Revised: October 16, 2014 Published: October 24, 2014 6361

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Figure 1. (a) Small-angle XRD patterns and TEM image (inset) of KIT-6 template; (b) N2 adsorption−desorption isotherm and pore size distribution calculated by the BJH method (inset) of KIT-6 template; (c) low-angle and wide-angle XRD patterns (inset) of mesoporous SnO2 and (d) N2 adsorption−desorption isotherms and pore size distribution (inset) of mesoporous SnO2.

stress.24 Recently, nanostructured SnO2-based anode materials have showed higher capacity beyond the above-mentioned theoretical capacity (783 mAh/g), but their lithium storage reaction mechanism still remains unclear. Demir-Cakan et al. reported that the mesoporous SnO2 electrode showed a reversible capacity of 960 mAh/g, which indicates that 5.6 mol of lithium ions are inserted into the SnO2 phase.25 They attributed the extra capacity (about 1.2 mol of lithium ions) to the partial reversible decomposition of the Li2O phase by using only cyclic voltammograms of the SnO2 electrode. Cyclic voltammetry is the fundamental tool for studying electrochemical reactions, but it cannot provide direct evidence of phase formation. Their explanation was contrary to the widespread understanding of the reaction mechanism of Snbased oxide anode materials, because the conversion reaction of Sn-based oxide materials is well-known to be completely irreversible and, thus, regarded as a main reason for irreversible capacity at the first cycle. Given that high irreversible capacity of Sn-based oxide anode materials is regarded as one of the key reasons that hinders the commercialization of Sn-based oxide anode materials, it is of utmost importance to completely understand the reaction mechanism of nanostructured Sn-based oxide materials that shows extra high capacity beyond theoretical capacity based on the alloying reaction mechanism. In this work, we report the detailed lithium storage and release mechanism of the nanostructured SnO2 anode material using an ordered mesoporous SnO2, which exhibits high surface area, high crystallinity, well-defined mesopore, and regular framework thickness, as a model material. Phase transition and the changes in the electronic and local structures around Sn atoms were investigated by synchrotron-based X-ray diffraction

(XRD) and X-ray absorption spectroscopy (XAS) analysis together with electrochemical tests.26−30

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EXPERIMENTAL SECTION

Mesoporous silica template with three-dimensional mesostructure (KIT-6, cubic Ia3d symmetry) was synthesized by the following previously reported method.31 Triblock copolymer (Pluronic 123, EO20PO70EO20, Aldrich, Mav = 5800) was used as a structure directing agent for KIT-6. A total of 9.0 g of P123 and 9.0 g of n-butanol was dissolved in 325.5 g of doubly distilled water at room temperature. A total of 17.7 g of concentrated HCl (Samchun, 35 wt %) solution was added in this mixture with continuous stirring. After stirring for 1 h, 19.4 g of tetraethyl orthosilicate (TEOS, 98%, Samchun) was added into the solution under vigorous stirring at 308 K for 24 h, and then the mixture was transferred in an oven at 373 K for another 24 h under static conditions. White precipitated product was filtered, washed with ethanol, dried at 353 K, and finally calcined in air atmosphere at 823 K for 3 h. Mesoporous SnO2 was synthesized via nanoreplication method by the following steps. Sn(II) chloride dihydrate (SnCl2 2H2O, 97%, Junsei) was used as precursors and melted at 373 K to the liquid state. Typically, 5.0 g of the KIT-6 was heated at 373 K for 1 h. The preheated KIT-6 was poured into a polypropylene bottle containing Sn precursors (5.14 g of SnCl2-2H2O). The bottle was closed and shaken vigorously for 1 h in order to mix the precursor solution homogeneously with the KIT-6 template. Subsequently, the bottle was placed in an oven at 353 K overnight for spontaneous infiltration of Sn precursor within the mesopores of the silica template. The composite materials then were heated to 973 K under ambient atmosphere for 3 h. The silica template was removed by treating the composite material with an aqueous solution of HF (10 wt %) several times. Finally, the obtained mesoporous SnO2 material was washed with distilled water and acetone several times and dried at 353 K for 24 h. SnO2 electrodes were prepared by slurring the active material powder (70 wt % mesoporous SnO2, 15 wt % Super-P, and 15 wt % 6362

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polyamide−imide (PAI)) and then coating the mixture on Cu foil. The anodes were incorporated into coin cells (2032 type) with a metallic Li foil as counter electrode and a Celgard separator. The electrolyte used was commercially available 1.3 M LiPF6 in a 3:7 ethylene carbonate (EC):diethyl carbonate (DEC) solvent, and the mass loading of active material per electrode was ∼3.25 mg/cm2. The discharge/charge cycling tests were conducted at constant current rate of C/10 between 0.001 and 2.0 V on the battery test station (WBCS3000, WonATech). The C rate was calculated by the weight of active material and its theoretical capacity. The low and wide-angle X-ray diffraction (XRD) patterns were obtained from a Rigaku D/MAX-III instrument equipped with Cu Kα radiation operating at 30 kV and 40 mA. N2 sorption isotherms were obtained using a Micromeritics ASAP 2000 at liquid N2 temperature. Before the measurements, the material was degassed for 12 h at 325 K. The Brunauer−Emmett−Teller (BET) and Barrett−Joyner−Halenda (BJH) methods were utilized to estimate the surface areas and pore size distributions. Scanning electron microscopy (SEM) images were collected using LEO SUPRA 55, GENESIS 2000 instrument at an accelerating voltage of 15 kV. High-resolution SEM images were obtained using a Hitachi UHR S-5500 FE-SEM operating at 30 kV. Transmission electron microscope (TEM) images were obtained using a G2 FE-TEM at operating voltage of 200 kV. High resolution synchrotron X-ray powder diffraction (HRPD) measurements of the samples were carried out on 9B HRPD beamline at Pohang Light Source-II (PLS-II). Pristine mesoporous SnO2 and bulk SnO2 powders were scanned from 10 to 130.5° at a scan rate of 0.1°/min. The incident X-rays were vertically collimated using a mirror and monochromatized to a wavelength of 1.5475 Å using a doublecrystal Si(111) monochromator. The detector comprises a set of six analyzer crystals and six scintillation detectors. XRD patterns for mesoporous SnO2 were collected on the 5A XRD beamline at PLS-II. Samples were prepared by subdividing points in the first and second discharge/charge cycles. Electrodes were collected by disassembling the coin cell after achieving the set discharge or charge capacity. Collected electrodes were washed by using diethyl carbonate (DEC) and dried in the dry room. Samples were sealed by using Kapton tape and packed in an Al pouch in order to prevent contamination. The wavelength of the X-ray beam was 0.7653 Å, and XRD patterns were recorded as a set of circles on a Mar 345-image plate detector in the transmission mode for about 1 min of exposure time. The total recording time was about 2.6 min because of the scanning time of the image plate and transferring time of spectral information. The two theta angles of all the XRD patterns presented in this article have been recalculated to corresponding angles for λ = 1.54 Å, which is the wavelength of the conventional X-ray tube source with Cu Kα radiations, for easy comparison with other published results. The XAS experiment was conducted using the 10C wide XAFS beamline with the Si(111) double crystal monochromator (PLS-II). The Sn K-edge spectra were recorded in fluorescence mode at room temperature. The storage ring was operated at 2.5 GeV with a ring current of 100−150 mA. Reference spectra of the Sn metal were simultaneously collected using Sn foil.

respectively. The well-arranged mesostructure of KIT-6 can also be observed in the HR-TEM image (inset of Figure 1a). Low- and wide-angle XRD patterns of mesoporous SnO2 are shown in Figure 1c. The low-angle XRD pattern has a wellresolved peak at 1.07° in 2θ, which can be indexed to the (211) reflection of the Ia3d space group expected as the exact negative replica structure of KIT-6. The other intense (110) reflection at 0.55° in 2θ indicates that the mesostructure of mesoporous SnO2 has transformed from the cubic Ia3d to tetragonal I41/a (or lower) during the replication. The wideangle XRD pattern shown in Figure 1c (inset) implies that mesoporous SnO2 has high crystallinity and the peaks can be indexed to tetragonal SnO2 (JCPDS 88−0287). To further investigate the mesoporous nature of prepared mesoporous SnO2, N2 adsorption−desorption measurements were carried out. The isotherms of mesoporous SnO2 exhibit a typical type IV in Figure 1d, indicating uniform mesoporosity. According to the corresponding pore size distribution curve in Figure 1d (inset), mesoporous SnO2 involves dual size pores of ∼3 and 20 nm. Narrow pore size distribution of ∼3 nm is originated from the exact replica of silica frameworks in the KIT-6 template, and broad pore size distribution of ∼20 nm arises from the exclusive pore formation in one of the two chiral pore channel system as expected for other mesoporous metal oxides.33 The BET surface area and pore volume of mesoporous SnO2 are 99 m2/g and 0.32 cm3/g, respectively. Figure 2 represents electron microscopy images of mesoporous SnO2. Mesoporous SnO2 consists of several

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RESULTS AND DISCUSSION Mesoporous SnO2 was synthesized via nanoreplication method using KIT-6 (space group of Ia3d) as hard template according to the literature reports.32 The small-angle XRD pattern of the KIT-6 template in Figure 1a shows a narrow, well-resolved (211) reflection peak at 0.98° in 2θ and several order reflections between 1° and 3°, indicating bicontinuous cubic gyroidal mesostructure (Ia3d). The lattice parameter (a0) calculated based on (211) reflection is 22.0 nm. The N2 adsorption−desorption isotherms and pore size distribution of the KIT-6 template, determined by BJH method, are presented in Figure 1b. The isotherms exhibit a typical type IV with a H1 hysteresis loop. The BET surface area, pore volume, and pore size of KIT-6 are 738 m2/g, 0.89 cm3/g, and 7.2 nm,

Figure 2. (a) Low magnification FE-SEM image; (b) high magnification FE-SEM image; (c) STEM image; and (d) HR-TEM image of mesoporous SnO2.

micrometers of particle with no specific shaped morphologies as shown in the low magnification SEM image (Figure 2a). High magnification SEM image (Figure 2b) and STEM image (Figure 2c) reveal that the mesostructure regularity is successfully preserved with uniform repeating units in longrange order even after nanoreplication. No silica and chloride are detected from energy-dispersive X-ray spectrum analysis, 6363

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which indicates silica template and a part of the precursor are all removed. The high resolution TEM image (Figure 2d) shows that mesoporous SnO2 is well crystalline, and lattice fringes in the 5 nm of framework are also observed, in agreement with the wide-angle XRD results. Rietveld refinement was carried out on SnO2 pristine powder diffraction data in order to collect in-depth structural and atomic information. Figure 3 shows the refined XRD pattern of

Figure 4. Voltage profile at C/10 rate between 0.001 and 2 V for the first two cycles and cycle performance during 10 cycles (inset) of mesoporous SnO2.

bulk SnO2, nano-SnO2, C/SnO2, and even some reports of GO/SnO2.36,43,44 This result implies that 6.25 mol of lithium ions are released from the mesoporous SnO2 electrode, which is much higher than the theoretical capacity of SnO2 based on the alloying mechanism of Sn with lithium (4.4 mol of lithium ions). In order to obtain the direct evidence on the origin of additional capacity of the mesoporous SnO2 electrode, we carried out synchrotron X-ray analysis on the mesoporous SnO2 electrode during cycling. Combination of XRD and XAS was used for probing the bulk and local structure of the mesoporous SnO2 structure. Figure 5 shows XRD patterns of mesoporous SnO2 at different states of charge. The XRD pattern of the pristine sample was completely indexed as the tetragonal space group P42/mnm, which confirms that synthesized material has high crystalline rutile SnO2 phase. The intensity of the XRD peaks drops significantly during discharge, and a closer look at the XRD patterns reveals that the

Figure 3. Observed, calculated, and difference plots obtained from Rietveld refinement of pristine mesoporous SnO2 high resolution powder diffraction pattern.

the SnO2 powder in which all of the peaks were indexed to single-phase tetragonal P42/mnm phase. Refined parameters and reliability factors are presented in Table 1. The unit-cell parameters are a = 4.74297 Å, c = 3.18918 Å, and V = 71.733 Å3, which are consistent with previous reports.34−36 Table 1. Atomic Sites (number of positions and Wyckoff notation) and Coordinates x, y, and z (in units of lattice constants) of Mesoporous Pristine SnO2a,b atom

site

Sn O

2a 4f

Wyckoff positions 0.00 0.3045(5) Rp Rwp Rexp S

0.00 0.00 0.3045(5) 0.00 reliability factors

B, Å2

OCCc

0.42(1) 0.01(7)

0.5 1.0

4.57 6.34 5.03 1.26

a

The weighted factor Rωp = 6.34%. bMesoporous SnO2 space group: P42/mnm, a (Å) = 4.74297(6), c (Å) = 3.18918(5). cFixed parameter.

Figure 4 shows the voltage profiles of mesoporous SnO2 at the first two cycles at a constant current rate of C/10 along with cycling performance of the first 10 cycles in the inset. The initial discharge/charge capacities were 2060.49 mAh/g and 1043.04 mAh/g, and the second discharge/charge capacities were 1109.68 mAh/g and 1021.07 mAh/g, respectively. The discharge capacity in the first cycle is higher than the collective discharge capacity of reactions 1 and 2 mentioned above. It is well-known that exceptionally high capacity during initial discharge in metal oxide anodes is attributed to irreversible electrolyte decomposition and its reaction with the electrode surface to form a solid electrolyte interphase layer at low voltage.37−42 The initial charge capacity achieved by mesoporous SnO2 structure, i.e., 1048.43 mAh/g, is also higher than

Figure 5. X-ray diffraction patterns of mesoporous SnO2 electrode during first discharge, first charge, and second discharge. 6364

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Figure 6. Sn K-edge XANES and EXAFS spectra with corresponding voltage profile taken in (a) the first discharge region, (b) the middle discharge region, and (c) the last discharge region of first cycle.

(110) reflection moves toward lower angles, indicating the expansion of lattice parameters in the a and b directions and gradual conversion of SnO2 crystallites into the amorphous nano-LiSnO2 phase.45 The XRD peaks almost disappeared after discharging below ∼0.7 V. However, by discharging below ∼0.2 V, two broad peaks appeared at about 22° and 38°, and their

intensity increased steadily with further depth of discharge. These newly observed peaks at the end of discharge can be associated with the crystal structure of LixSn alloy. As the quantity of Li increases in Sn host structure, the crystal structure of the LixSn alloys goes through a series of phase transitions.46 In the start of charge, the peaks damp down again 6365

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Figure 7. Sn K-edge XANES and EXAFS spectra with corresponding voltage profile taken in (a) the first charge region and (b) the last charge region of first cycle.

and disappear above ∼0.9 V, which comes from dealloying the reaction of the LixSn alloy. In the end of the second discharge, Bragg peaks corresponding to the LixSn phase appear reversibly, and the alloy composition in the fully discharged state can be assigned as Li4.4Sn.13,47,48 These XRD patterns clearly confirm the reversibility of the alloying reaction, which should yield a maximum capacity of 783 mAh/g. Figure 6a represents selected Sn K-edge XANES and EXAFS patterns in the first discharge region of the first cycle. The oxidation state of Sn in the ordered mesoporous SnO2 is found as Sn4+. The reduction of Sn ions takes places in the start of discharge; the Sn K-edge XANES spectra show prominent shift toward lower energy values representing the decrease in the average oxidation of Sn. This reduction of Sn through conversion reaction affects the local environment around the Sn atom and influences the EXAFS spectra. The first prominent peak at ∼1.5 Å in Sn K-edge EXAFS spectra corresponds to the Sn−O interaction in the first coordination sphere, and the broad peaks in the region of 2.2−3.9 Å are due to the Sn−Sn, Sn−O, and Sn−Sn interactions in the second, third, and fourth coordination spheres, respectively. The amplitudes of these peaks decrease significantly with the increase of depth of the discharge due to displacement of reacting species during conversion reaction. Figure 6b shows XAS data obtained in the middle discharge region of the first cycle. In this discharge

region, Sn K-edge XANES spectra show negligible shift toward lower energy values. However, white line intensities keep decreasing in this region, suggesting the formation of metallic Sn phase involved in the alloying reaction. XRD patterns show evolution of a totally amorphous phase in this discharge region, but the changes in the EXAFS spectrum show that atomic rearrangements are still taking place. In the EXAFS spectra after 600 mAh/g, the amplitude of the Sn−O peak steadily decreases, and a new peak at around 2.6 Å appears, which corresponds to the Sn−Sn(Li) pair in the LixSn alloy, and amplitude of this new peak grows with the increase of the Li/Sn ratio.49 Note that both the conversion reaction of SnO2 and the alloying reaction of Sn formed in the SnO2 electrode can simultaneously occur in this region where the potential of the SnO2 electrode ranges between 1.0 and 0.3 V( vs Li/Li+).50 However, previous GITT studies showed that the reaction kinetics of the conversion reaction are significantly slower than those of the alloying reaction;51,52 thus, conversion and alloying reactions might take place simultaneously in a sequential way, and Sn metal formed during the conversion reaction reacts with lithium to form LixSn alloy during the process of the conversion reaction. The intensity of the Sn−O peak decreases, and that of the Sn−Sn(Li) peak increases during this discharge region. The Sn−O peak disappears completely when the discharge capacity approaches 1500 mAh/g, which indicates the completion of the 6366

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Figure 8. Sn K-edge XANES and EXAFS spectra with corresponding voltage profile taken in (a) the first discharge region and (b) the last discharge region of the second cycle.

alloying metals.54 Charge redistribution takes place to minimize the electrostatic energy and shifts the Sn K-edge during the alloying reaction. In the EXAFS spectra, the amplitude of the Sn−Sn peak continuously decreases with the increase of the depth of discharge in this region. Since the increase of the Li/ Sn ratio increases the amount of Li around Sn, Li has a much smaller electron scattering cross section as compared to Sn, and the intensity of the Sn−Sn(Li) peak starts to decrease when the Li/Sn molar ratio exceeds ∼3.55 This trend of XANES and EXAFS data suggests that the capacity in this deep discharge region is solely attained by further lithium alloying of the LixSn phase until it achieves its nominal terminal composition of Li4.4Sn. Figure 7a shows XAS data taken from the ordered mesoporous SnO2 electrode in the first charge region of the first cycle. The Sn K-edge XANES spectra shift reversibly toward lower energy values, and EXAFS spectra show the rise of Sn−Sn(Li) related peaks, implying that only the dealloying reaction occurs in this potential region. After the cell was charged up to 500 mAh/g, the amplitude of the Sn−Sn(Li) peaks starts to decrease and the Sn−O peak starts to rise as shown in Figure 7b. Appearance of the Sn−O peak and damping of the Sn−Sn peak show the formation of the SnOx phase in this charge region, indicating that partial reversibility of converted SnO2 in the first discharge at the cost of the Li2O decomposition. Formation of the SnOx phase was also suggested by Shiva et al. by using cyclic voltammetry during

conversion reaction, and the remaining discharge capacity can be attributed to the alloying reaction only. Figure 6c shows the XAS data taken from the ordered mesoporous SnO2 electrode in the last discharge region of the first cycle. Considering that lithium is more electropositive than Sn, it is expected that the edge energy of Sn in this region would be decreased with the increased of amount of lithium inserted. However, we found that XANES spectra in this region shift toward higher energy values upon further lithiation. A similar XANES result had been already reported by Mansour et al.53 They reported that XANES spectra of Sn-based composite oxide glass slightly shift to a higher energy without significant change in the overall shape of the data; thus, they assigned the data as Sn−Sn type of bonding. However, our XANES results obtained from the ordered mesoporous SnO2 electrode clearly shows the shift toward high energy of the Sn K-edge. This shift can be explained by charge transfer between different metals during alloy formation. A new covalent type bond is superimposed over the metallic bond between dissimilar metals, i.e., Li and Sn, during the alloying process, which causes the redistribution of electrons between these metals. Due to the intrinsic difference between the metallic bond and its superimposed covalent bond, delocalized electrons have the freedom to slide past each other. In alloys, atomic sites tend to remain neutral, but a small amount of net charge transfer takes place depending on the electronic structure, atomic radii, and electronegativity of the 6367

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Figure 9. Sn K-edge XANES and EXAFS spectra with corresponding voltage profile taken in (a) the first charge region and (b) the last charge region of the second cycle.

the electrochemical cycling of the SnO2−rGO composite.56 These results suggest that the reversible charge capacity at the end of the charge can be attributed to not only dealloying of LixSn phase but also the conversion reaction of Sn into the SnOx phase. After achieving charge capacity of 900 mAh/g, the Sn−O peak shows significant rise with negligible change of the Sn−Sn peak, implying that reversible charge capacity is mostly achieved by conversion reaction in this region. XANES spectra do not show substantial shift in this charge region. The dealloying reaction of the LixSn alloy phase tends to move the Sn K-edge toward lower energy level whereas Sn oxidation due to the formation of SnOx should shift the K-edge toward higher energy positions. These two processes take place simultaneously during the deepest charge region; thus Sn K-edge maintains the stagnant position. The EXAFS and diffraction patterns in the first cycle show that active material in SnO2 electrode does not come back to its starting composition after one complete discharge/charge cycle; instead the SnO2 electrode turned into metallic Sn with a small quantity of amorphous SnOx as well as LixSn phase. To further investigate electrochemical reactions of mesoporous SnO2 electrode material, XAS studies were performed during the second cycle at an interval of 100 mAh/g. The first discharge region in the second cycle is shown in Figure 8a. Considering that the electrode potential drastically drops below 1 V after achieving discharge capacity of 100 mAh/g only, the

amount of crystalline SnO2 in the electrode would be much smaller compared to that of the pristine electrode. At the initial discharge region, XANES spectra did not show any significant shift due to simultaneous occurrence of both the alloying and conversion reactions during the second cycle. In EXAFS data, the peak corresponding to the Sn−O bond disappears, and the peaks corresponding to the Sn−Sn(Li) bond emerges after achieving discharge capacity of 400 mAh/g, which indicates the completion of the conversion reaction and progression of the alloying reaction. Figure 8b shows the last discharge region in the second cycle. The Sn K-edge shifts to higher energy values upon further discharge. The amplitude of the Sn−Sn(Li) peak in the EXAFS spectra increases up to 500 mAh/g and decreases rapidly until the end of discharge, while the Sn−O peak remains absent in the EXAFS spectrum. These results suggest that only the alloying reaction prevails in this deepest discharge region. Figure 9a shows the XAS data obtained in the first charge region of the second cycle. The Sn K-edge XANES spectra shifts to the lower energy values, and intensity of Sn− Sn(Li) peaks grows in the EXAFS spectra as a result of lithium extraction from the Li4.4Sn phase. Figure 9b represents the XAS data obtained in the last charge region of the second cycle. Sn K-edge XANES spectra maintain their position, and the Sn−O peak starts to appear in the EXAFS spectra, which indicates that dealloying along with partial conversion into the SnOx phase occurs in the second cycle as well. XAS results at the second 6368

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Figure 10. Diagram of overall electrochemical reaction mechanism of mesoporous SnO2 during two cycles.

cycle confirm the proposed mechanism derived from the first cycle, which includes complete reversibility of alloying and partial reversibility of the conversion process. These results successfully explain the origin of high capacity of mesoporous SnO2 beyond its reported theoretical capacity. On the basis of this set of results, the overall reaction mechanism is summarized in Figure 10.

is gratefully acknowledged. M.B. is supported by the U.S. DOE, Office of Science (Contract No. DE-AC02-06CH11357).

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CONCLUSIONS Ordered mesoporous SnO2 prepared by hard templating showed a reversible capacity of about 1000 mAh/g, which is higher than the predicted value based on the alloying reaction mechanism of Sn, formed from conversion reaction of SnO2 with lithium. Synchrotron XAS analysis combined with XRD revealed that some portion of the Li2O phase decomposes to form the SnOx phase with Sn upon delithiation, parallel with dealloying of LixSn, which leads to unexpected high capacity of an ordered mesoporous SnO2 material. We firmly believe that these results will not only help further understanding of the complicated reaction mechanism of nanostructured Sn-based anode materials but also provide a valuable insight into designing new electrode materials with abnormal high capacity for next generation rechargeable batteries.

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REFERENCES

AUTHOR INFORMATION

Corresponding Authors

*E-mail: [email protected] (Hansu Kim). *E-mail: [email protected] (Ji Man Kim). *E-mail: [email protected] (Won-Sub Yoon). Notes

The authors declare no competing financial interest.

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ACKNOWLEDGMENTS This work was supported by Samsung Research Funding Center for Future Technology (SRFC-MA1401-03); the Energy Efficiency & Resources Core Technology Program of the Korea Institute of Energy Technology Evaluation and Planning (KETEP) granted financial resources from the Ministry of Trade, Industry & Energy, Republic of Korea (No. 20132020000260). We also thank the National Research Foundation of Korea (NRF-2010-C1AAA001-2010-0029065 and the Mid-Career Researcher Program No. 2012R1A2A2A01010011) for partial support. Preliminary studies at sector-20 BM in Argonne’s Advanced Photon Source 6369

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