Reversible Lithium Storage at Highly Populated Vacant Sites in an

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Reversible Lithium Storage at Highly Populated Vacant Sites in an Amorphous Vanadium Pentoxide Electrode Oh B. Chae,† Jisun Kim,† Inchul Park,‡ Hyejeong Jeong,† Jun H. Ku,† Ji Heon Ryu,§ Kisuk Kang,‡ and Seung M. Oh*,† †

Department of Chemical and Biological Engineering and ‡Department of Materials Science and Engineering and Center for Nanoparticle Research at the Institute for Basic Science (IBS), Seoul National University, Seoul 151-742, Korea § Graduate School of Knowledge-based Technology and Energy, Korea Polytechnic University, Siheung-si 429-793, Korea S Supporting Information *

ABSTRACT: A vanadium pentoxide electrode is prepared in the amorphous form (a-V2O5), and its electrode performances are compared to those for its crystalline counterpart (c-V2O5). The aV2O5 electrode outperforms c-V2O5 in several ways. First, it is free from irreversible phase transitions and Li trapping, which evolve in cV2O5, probably due to the lack of interactions between the inserted Li+ ions/electrons and V2O5 matrix. Second, the absence of Li trapping allows a reversible capacity amounting to >600 mA h g−1, which is larger than that given by c-V2O5. Third, it shows an excellent rate property. The notably high reversible capacity and rate capability seem to be due to Li storage at vacant sites that are ill-defined but numerous in a-V2O5, which Li+ ions can easily access. However, irreversible capacity of a-V2O5 is appreciable in the first cycle due to a parasitic Li reaction with surface hydroxyl groups. Treatment with nbutyllithium can suppress the irreversible capacity by removing the surface hydroxyl groups.

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INTRODUCTION A variety of metal oxides in their crystalline form have been exploited as the negative electrode for lithium secondary batteries, wherein the lithiation reactions proceed by two major routes, conversion and insertion reactions. In the former, the metal−oxygen bonds are broken and the metal ions are reduced to their elemental states. In the insertion-type lithiation, however, Li+ ions/electrons are inserted into the oxide lattice without bond cleavage.1−6 In general, the oxides of latetransition-metal elements (for instance, CoO7 and Fe2O38,9) and main-group elements (for instance, SnO210,11) are lithiated by the conversion reaction, whereas the oxides of earlytransition-metal elements (TiO2,1,5 V2O5,2,6 and MoO24,12−16) are lithiated by the insertion reaction. The difference between the two groups is, among others, the metal−oxygen bond strength. The former group has a relatively weaker metal− oxygen bond (the bond dissociation energy is 368 kJ mol−1 for Co−O in CoO vs 644 kJ mol−1 for V−O in V2O5),17 and thus is lithiated with bond cleavage (conversion reaction). The latter group, which possesses a relatively stronger metal−oxygen bond, is lithiated without bond cleavage (insertion reaction). One characteristic feature of the conversion-type oxides is the relatively high lithiation capacity that is governed by the oxidation state of the metal component. For instance, Fe2O3 is lithiated by taking six Li+ ions/electrons per formula unit since the valence of Fe is 3+: Fe2O3 + 6Li+ + 6e− → 2Fe + 3Li2O. The advantage of a high lithiation capacity is, however, offset by © XXXX American Chemical Society

the large voltage hysteresis and poor cycling stability of these oxides. The opposite is true for the insertion-type oxides: the Li insertion/removal is reversible without significant voltage hysteresis, but their Li storage capacity is relatively small (0.5 Li+ ion/electron for TiO2 and 1.0 Li+ ion/electron for MoO2) when they are prepared as a highly crystalline form, since Li+ ions are inserted only into the crystallographic Li+ storage sites that must be limited in number in highly crystalline materials. The crystalline vanadium pentoxide (V2O5), which is in the orthorhombic phase (α phase, α-V2O5) in common preparations, is another insertion-type oxide. However, it shows somewhat different lithiation behavior compared to the other insertion-type oxides (TiO2 and MoO2): its lithiation capacity is unexpectedly high (5 Li+ ions/electrons per V2O5) in the first lithiation, but only half of the Li+ ions/electrons are released in the forthcoming delithiation. This implies that, unlike those of TiO2 and MoO2, the lattice of α-V2O5 provides a large number of crystallographic sites that are accessed by Li+ ions, but the lattice is easily transformed into more stable phases after taking Li+ ions/electrons. That is, Li+ ions/electrons are irreversibly trapped inside the lattice. This work was motivated by the premise that V2O5 can be a high-capacity negative electrode if the Li trapping and Received: June 24, 2014 Revised: September 28, 2014

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instrument. The mass signal was recorded as a function of the cell voltage on an HPR-20 gas analyzer (Hiden Analytical Ltd.). Electrochemical Characterizations. For the preparation of working electrodes, the V2O5 powder, Super P, and poly(vinylidene fluoride) (PVdF) (8:1:1, wt %) were dispersed in N-methylpyrrolidone, and the resulting slurry was coated on a piece of copper foil. The electrode plate was dried in a vacuum oven at 120 °C for 12 h and pressed. Coin-type cells (2032-type) were fabricated with lithium foils as the counter and reference electrodes. The used electrolyte was 1.0 M LiPF6 dissolved in a mixture of ethylene carbonate (EC) and dimethyl carbonate (DMC) (1:2, volume ratio). Computational Details. First-principles calculations were performed using the Perdew−Burke−Ernzerhof exchange-correlation parametrization to density functional theory (DFT) with the spinpolarized generalized gradient approximation (GGA).22 A plane-wave basis set and the projector-augmented wave (PAW) method, as implemented in the Vienna Ab-initio Simulation Package (VASP), were used.23 To study the structure and the energy of ω-Li3V2O5, all possible orderings of Li−V with a face-centered cubic (fcc) arrangement of the oxygen framework were considered.24 All plausible configurations of Li−V cation site exchange within the 20 sites were considered for ω-Li3V2O5. For ω-LixV2O5 (x = 0−2), all plausible configurations of Li vacancies were considered. The structure of aV2O5 was calculated from ab initio molecular dynamics (MD) simulation. The model system was first thermally equilibrated in a microcanonical ensemble (NVE) for 2 ps followed by an MD production run of 10 ps where the temperature was controlled by a Nose−Hoover thermostat. The position of lithium ions in the amorphous sample was expected on the basis of an experimental random mixture.

irreversible phase transition are suppressed. Note that V2O5 belongs to the insertion-type oxides that are free from serious voltage hysteresis and massive volume change. Hence, if the critical drawback of Li trapping is solved, V2O5 can be a prospective electrode material for lithium ion batteries. The approach set in this work is the disturbance of crystallographic Li+ storage sites that are responsible for Li+ trapping and irreversible phase transitions by incorporating disorders in the lattice. The expectation here is that the interactions between the inserted Li+ ions/electrons and V2O5 matrix are disturbed by the presence of disorders to discourage the phase transitions. To implement this idea, vanadium pentoxide was prepared in amorphous form (a-V2O5) and its lithiation behavior was examined. In the literature,18−21 amorphous V2O5 was prepared as xerogels or aerogels that have a large surface area (400−1000 m2 g−1). These gel-type oxides show an unusually high Li+ storage capacity. Even if the detailed Li+ storage mechanism was not characterized for these highly porous materials, it is very likely that a considerable amount of Li+ ions/electrons are stored at the surface. Hence, in this work, an amorphous V2O5 having a relatively small surface area (15 m2 g−1) was prepared to eliminate the surface reactions. The following objectives have been identified in this work: (i) to see if the irreversible phase transitions are prevented by preparing V2O5 in the amorphous form, (ii) to identify the nature of Li+ storage sites in a-V2O5, and (iii) to examine the lithiation/delithiation behavior of a-V2O5, including cycle performance, rate capability, and the nature of irreversible reactions.

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RESULTS AND DISCUSSION Figure 1a shows the XRD patterns of two V2O5 samples. The sample prepared by drying the precipitate in a vacuum oven at

EXPERIMENTAL SECTION

Materials Synthesis. The amorphous vanadium pentoxide (aV2O5) powder was prepared by a simple precipitation method. In detail, 1.63 g of VOSO4 (Sigma-Aldrich, 97%) was dissolved in 100 mL of deionized water, and the resulting solution was slowly added to an air-purged NH4OH solution. A brown precipitate formed immediately. The precipitate was isolated as a wet cake by centrifugation and washed with ethanol. The a-V2O5 powder was obtained by drying the precipitate in a vacuum oven at 100 °C for 24 h. A crystalline vanadium pentoxide (c-V2O5) powder was also prepared for comparison purposes by heat treatment of the a-V2O5 powder at 600 °C for 3 h in air. Acidic surface hydroxyl groups on aV2O5 were removed by treatment with n-butyllithium (n-BuLi), which is known as a strong base. For this treatment, 0.5 g of a-V2O5 powder was dispersed in 0.05 M n-butyllithium/hexane solution (SigmaAldrich) (50 mL) and the resulting dispersion stirred overnight. The treated powder was collected by centrifugation and dried in a vacuum oven at 100 °C for 24 h. Material Characterizations. X-ray diffraction (XRD) patterns were recorded with a D-MAX2500-PC (Rigaku Co.) using Cu Kα radiation (1.54056 Å). The particle morphology was examined using field-emission scanning electron microscopy (FE-SEM; JEOL JSM6700F). The surface area was measured from the nitrogen adsorption isotherm (Micromeritics, ASAP 2010). The CHNS contents in the samples were measured by elemental analysis (EA). The adsorbed water and hydroxyl groups were determined from thermogravimetric analysis (TGA) data that were obtained from room temperature to 600 °C at 10 °C min−1 in an argon atmosphere. X-ray absorption spectroscopy (XAS) data were collected near the K-edge of vanadium (E0 = 5465.1 eV) in the transmission mode at the Pohang Light Source (PLS) with a ring current of 120−170 mA at 2.5 GeV. Fouriertransform infrared (FT-IR) spectra were obtained by using a Nicolet 6700 (Thermo Scientific) in the transmission mode. The samples were prepared as a KBr pellet. H2 evolution was analyzed by using a homemade differential electrochemical mass spectroscopy (DEMS)

Figure 1. (a) XRD patterns of c-V2O5 and a-V2O5, (b) FE-SEM image of c-V2O5, and (c) FE-SEM image of a-V2O5.

100 °C turns out to be amorphous (a-V2O5), whereas the other one that was prepared by heat treatment of the as-prepared amorphous sample at 600 °C is crystalline (c-V2O5). The diffraction peaks for the latter are well-matched with those for the orthorhombic phase (α phase, α-V2O5, JCPDS no. 86B

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Figure 2. Galvanostatic lithiation/delithiation voltage profiles obtained from (a) Li/c-V2O5 and (b) Li/a-V2O5 cells. The solid and dashed lines represent the first and second voltage profiles, respectively. Corresponding differential capacity (dQ/dV) plots for (c) Li/c-V2O5 and (d) Li/a-V2O5 cells. Current density 100 mA g−1. Voltage cutoff 0.01−3.0 V (vs Li/Li+).

Figure 3. (a) XRD patterns obtained from the crystalline V2O5 electrode (α-V2O5, JCPDS no. 86-2248) after lithiation (1.5 and 0.01 V) and delithiation (3.0 V). (b) Calculated voltage and galvanostatic lithiation/delithiation voltage profiles in the first cycle obtained from the Li/c-V2O5 cell.

can be mitigated by treatment with n-butyllithium. The details will be advanced in a later section. The galvanostatic discharge (lithiation) and charge (delithiation) voltage profiles and the differential capacity (dQ/dV) plots obtained from the Li/c-V2O5 and Li/a-V2O5 cells are presented in Figure 2. In the first lithiation (solid line), the crystalline electrode (c-V2O5) gives several voltage plateaus (Figure 2a) that appear as peaks in the dQ/dV plot (Figure 2c). The voltage plateaus at >1.5 V are known to be associated with the consecutive phase transitions from the initial α phase to the ε, δ, γ, and eventually ω phases, which proceed with lithium insertion.24,28,29 The formation of the final ω phase is confirmed in this work by XRD analysis. As shown in Figure 3a, the pristine α phase is converted into the ω phase (ωLixV2O5) upon lithiation down to 1.5 V. Interestingly, the asformed ω phase is retained during the forthcoming lithiation down to 0.01 V and delithiation up to 3.0 V. As a result, the

2248). The particle size of the two powder samples was less than 1 μm as seen in the FE-SEM images (Figure 1b). The Brunaur−Emmett−Teller surface area for the amorphous and crystalline samples was 15 and 10 m2 g−1, respectively. The oxidation state of vanadium in the two samples was estimated to be 5+ from the presence of a strong pre-edge peak at 5469 eV in the X-ray absorption near-edge structure (XANES) data (Figure S1, Supporting Information).25,26 The absence of residual sulfate ions in a-V2O5 was confirmed by the elemental analysis data, in which the sulfur content was less than the detection limit. TGA was performed to assess the amount of adsorbed water and hydroxyl groups in a-V2O5 (Figure S2, Supporting Information). In the TGA profile, the mass loss at 1.5 V in the dQ/dV plot (Figure 2c) is much larger than that for the delithiation peak. This means that Li+ ions/ electrons, which are consumed in the serial phase transitions at >1.5 V, are not released during the delithiation. In short, the crystalline V2O5 in the α phase is converted into the ω phase (ω-LixV2O5) in the first lithiation, but a substantive amount of Li+ ions/electrons are irreversibly trapped inside the lattice. The origin of Li+ trapping is further studied by first-principles calculations. Figure 3b shows a theoretical prediction of the phase transitions in the first cycle. In lithiation, it is predicted that the α-V2O5 phase as a layered structure initially takes 3 Li+ ions/electrons per V2O5 at 2.41 V on average and transforms into the ω-LixV2O5 phase as a rock-salt structure. This prediction is well-matched with the experimental data up to x = 3 in LixV2O5. Further lithiation below 1.5 V is not considered in our calculations and will be discussed in detail later. In delithiation (>1.5 V), however, the rock-salt ω-LixV2O5 needs a higher voltage to extract Li+. The average extraction voltage is 2.80 V, which is higher than that for lithiation. The change in charging potential is the result of irreversible phase transitions to ω-LixV2O5 and the corresponding delithiation potential from the new framework. While the predicted charging potential in the initial delithiation (>1.5) agrees well with the experiments in general, the experimental delithiation of the higher potential region could not be observed and, thus, could not be compared with the calculations. We believe this is partly due to a large overpotential occurring during the delithiation from the rocksalt ω-LixV2O5, which is likely to exhibit very sluggish Li+ ion diffusivity compared to that of the layered α-V2O5. The discharge/charge voltage profiles and differential capacity (dQ/dV) plots obtained from the Li/a-V2O5 cell are presented in parts b and d, respectively, of Figure 2. The first lithiation voltage profile shows a monotonic decrease in the whole potential range (3.0−0.01 V) without discernible voltage plateaus. Hence, the peaks at >1.5 V in the dQ/dV plot (Figure 2c), which are associated with the irreversible phase transitions in c-V2O5, almost disappear (Figure 2d). Certainly, the approach set in this work to suppress the phase transitions works, which is the incorporation of disorders in the lattice. It is also very likely that Li trapping is not serious since this electrode is free from the irreversible phase transitions. Evidently, the first Coulombic efficiency for this electrode (aV2O5) is much larger (72%; 944 mA h g−1 for lithiation and 684 mA h g−1 for delithiation) than that for c-V2O5 (52%; 731 mA h g−1 for lithiation and 381 mA h g−1 for delithiation). Note, however, that this electrode (a-V2O5) still shows an appreciable irreversible capacity in the first cycle (Figure 2b,d). The details on this will be discussed later.

Figure 4. (a) Cyclic voltammograms obtained from the amorphous V2O5 electrode as a function of the scan rate. (b) Current versus square root of the scan rate. The current values at 0.01 V (lithiation) and 3.0 V (delithiation) were plotted according to the square root of the scan rate.

potentials (0.01 and 3.0 V) are plotted as a function of the square root of the scan rate (v1/2) in Figure 4b. Both the lithiation and delithiation currents exhibit a linear relation with v1/2. Clearly, the capacitor behavior (current ∞ v) is not observed, eloquently demonstrating that Li+ ions/electrons are stored in the bulk of a-V2O5 rather than at the surface. As presented above, the a-V2O5 electrode shows reversible Li storage without phase transitions. This feature is further evidenced by the vanadium K-edge extended X-ray absorption fine structure (EXAFS) data shown in Figure 5. The c-V2O5 electrode in its pristine state gives two distinct peaks in the FT profile (Figure 5a). The one at 1.0−2.5 Å comes from the V−O bonds of the first coordination sphere, which forms a square pyramidal structure.26,30,31 The peak at 2.5−3.4 Å represents the distance between two neighboring vanadium ions (V− V).26,30,31 When c-V2O5 is lithiated down to 0.01 V, the peaks D

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more symmetric than that (square pyramidal) for the initial αV2O5. The peak sharpening and intensifying is caused by the merging of four shorter equatorial V−O bonds and one longer apical vanadyl bond (in square pyramidal structure) into one V−O bond in octahedral geometry. Also noted in Figure 5a is the peak shift to the shorter radial distance for V−O and V−V. This illustrates that the first coordination sphere around V ions is contracted upon lithiation. Meanwhile, the FT profile in Figure 5a reveals that the peak intensity for V−O and V−V decreases upon delithiation (3.0 V) to be comparable with that for the pristine sample, indicating that the local structure becomes less symmetric with lithium removal. However, the peak position is not the same as that for the pristine sample. The V−O bonds are longer and the V−V distance is shorter than those observed in the pristine sample. This illustrates that the local structure near V ions is not perfectly restored back to the initial one even if lithium ions are removed from the lattice. Surely, this is due to the irreversible phase transitions in this crystalline sample. The EXAFS data obtained from the amorphous electrode are displayed in Figure 5b. An immediately apparent feature is the absence of FT peaks at >3 Å, indicative of the absence of longrange ordering. Short-range ordering is, however, clearly seen in Figure 5b: V−O bonds at 1−2 Å and V−V bonds at 2−3 Å for the pristine sample. The peak growing and sharpening for V−O and V−V is also observed when this electrode is lithiated down to 0.01 V, indicating that this amorphous electrode also experiences a similar symmetry change upon lithium insertion. Nonetheless, there appears a feature distinctively different from that of the crystalline electrode. As seen in Figure 5b, the FT peaks for V−O and V−V for the delithiated (up to 3.0 V) sample are very close in intensity and position to those observed with the pristine one. This explains that the local structure near V ions is restored back to the initial one when Li ions are removed from the amorphous electrode. Surely, this resulted from the reversible Li insertion/removal without irreversible phase transitions. Figure 6 illustrates the schematic structural change of c-V2O5 and a-V2O5 during the lithiation down to 0.01 V/delithiation up to 3.0 V in the first cycle. In c-V2O5, the layered structure

Figure 5. Fourier transforms of vanadium K-edge EXAFS profiles: (a) crystalline V2O5 and (b) amorphous V2O5.

for V−O and V−V are intensified and sharpened, reflecting that the local structure near V becomes more symmetric with lithium insertion.26 This feature can be explained on the basis of previous reports20,32 which illustrate that vanadium ions in αV2O5 migrate from their original square pyramidal sites into the neighboring empty octahedral sites with lithiation. This means that the local symmetry of V ions in ω-LixV2O5 (octahedral) is

Figure 6. Schematic representation of structural changes upon lithiation down to 0.01 V and delithiation up to 3.0 V for c-V2O5 and a-V2O5. E

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transforms into the rock-salt structure upon lithiation. This structure is, however, not restored back to the initial one, even if Li+ ions are removed from the lattice. Meanwhile, a-V2O5 retains its structure during the cycle, and the local structure is only slightly changed upon lithiation. The structural stability of a-V2O5 may contribute to the reversible Li insertion/extraction. It is time to discuss the nature of Li storage sites in a-V2O5. Our previous work reported that an amorphous molybdenum dioxide (a-MoO2) electrode delivers a larger capacity than the crystalline counterpart.33 Such an enlarged capacity is also observed with nanosized TiO234 and Li4Ti5O1235 and less crystalline carbons.36,37 One unique feature for all these electrodes is the sloping voltage profile, which is the case for a-V2O5 in this work. This feature can be explained on the basis of the voltage profiles given by carbon electrodes. It is wellknown that highly crystalline carbons (graphite) are lithiated by a series of two-phase reactions that are known as the stage phenomenon. The lithiation potential for each two-phase reaction is discrete to appear as a series of voltage plateaus. In contrast, the lithiation voltage profiles are sloped for less crystalline carbons (hard carbons and soft carbons), illustrating that the lithiation reactions take place over a wide potential range since the Li+ storage sites, which have been proposed to be vacancies, void spaces, cluster gaps, or interstitial sites, are electrochemically nonequivalent to each other.36−38 In this work, the c-V2O5 electrode exhibits a behavior similar to that of graphite: the lithiation potential is discrete for each two-phase reaction (for instance, α to ε phase transition), which appears as a series of voltage plateaus at >1.5 V (Figure 2a) and as peaks in dQ/dV plot (Figure 2c). The a-V2O5 electrode, however, shows a sloping voltage profile like the less crystalline carbons. Hence, the Li+ storage sites in a-V2O5 can be described on the basis of those proposed for the less crystalline carbons. Namely, amorphous materials have randomly ordered atoms in the structure, such that a high population of vacant sites (cation and anion vacancies, void spaces, cluster gaps, or interstitial sites) can be assumed in a-V2O5. Moreover, Li+ ions seem to easily access those sites since the active surface area for Li+ ions is enlarged due to the tiny particle size of a-V2O5. Cycle performance is compared for two cells in Figure 7a, in which a-V2O5 outperforms c-V2O5 with respect to reversible capacity and cycle performance. The reversible capacity of aV2O5 amounts to 600 mA h g−1, demonstrating that the abovementioned Li+ storage sites are abundant in number and stably accept/release Li+ ions/electrons. The rate capability is also excellent for a-V2O5 (Figure 7b). The capacity delivered at 5000 mA g−1 amounts to 85% of the value obtained at 100 mA g−1. Note that the current density of 5000 mA g−1 corresponds to 13 C for the graphite electrode. This excellent rate property for a-V2O5 is likely due to the rather opened Li+ diffusion pathways provided by the vacant sites.33 As pointed out above, the a-V2O5 electrode still shows an appreciable irreversible capacity in the first cycle. The absence of lithiation peaks at >1.5 V in the dQ/dV plot (Figure 2d) illustrates that this irreversible capacity does not result from Li trapping. Two possibilities are responsible for the irreversible reactions at