Comparison of Storage Mechanisms in RuO2, SnO2, and SnS2 for

Article pubs.acs.org/JPCC

Comparison of Storage Mechanisms in RuO2, SnO2, and SnS2 for Lithium-Ion Battery Anode Materials Ayorinde S. Hassan,† Kathleen Moyer,†,‡ B. Ramu Ramachandran,*,† and Collin D. Wick*,† †

Chemistry, College of Engineering & Science, Louisiana Tech University, Ruston, Louisiana 71272, United States Stevens Institute of Technology, Hoboken, New Jersey 07030, United States

‡

ABSTRACT: A comparison of the molecular mechanisms of lithium sorption in RuO2, SnO2, and SnS2 was conducted with first-principles calculations. The calculated discharge curves for lithium absorption qualitatively agreed with experiment for the three materials. Our computations show that high-capacity lithium sorption in RuO2 beyond the stoichiometric conversion limit is due to “interfacial” storage, i.e., lithium absorption in the interface between the metallic ruthenium islands and the lithium oxide produced by conversion. Lithium sorption beyond the conversion limit in SnO2 and SnS2 is due to alloying, but interfacial storage is also found to contribute beyond the alloying limit. Therefore, interfacial storage appears to be a universal high-capacity mechanism for metal oxide and metal sulfide materials, accounting at least partially for the observed capacity beyond stoichiometric limits. Among the three materials examined, SnS2 is shown to expand the least with lithium sorption, furthering its promise as an anode material for Li-ion batteries. with respect to Li/Li+.22−28 Metallic tin is particularly attractive because of its low cost and wide availability.29 However, pure tin suffers from very high volume fluctuations which lead to fracturing and pulverization of the electrode materials and loss of electrical contact, resulting in rapid capacity fading.30 SnO2, on the other hand, has reduced volume fluctuations,30−32 along with a theoretical specific capacity of 782 mAh/g that is more than twice the reversible gravimetric capacity of graphite (372 mAh/g).33−36 SnO2 is believed to absorb lithium during its first discharge cycle in a two-step process:31

I. INTRODUCTION The future of energy sustainability is dependent on the discovery and design of new materials for energy storage and generation, one important facet of which is electrochemical storage. Among such storage devices, rechargeable lithium-ion batteries (LiBs) are very attractive because of their high gravimetric energy density1−3 resulting in their pervasive usage in portable consumer electronic devices, and potential wide use in electric vehicles.4,5 Cathodes and anodes represent important components of LiBs because they control the capacity of the cell, are critical for their cyclability, and the cell potential is dependent on them. In order to meet future energy demands, the design and development of new safe low-cost electrode materials with higher energy density, cyclability, and mechanical stability is required.3,5−7 However, the capacity of the present generation of LiBs anode is limited to 372 mAh/g, which is the theoretical capacity of the widely used layered/intercalating graphite anode.8−11 Among the current research trends in this area is, therefore, the search for materials with low cost, higher capacity, good Coulombic efficiency, and cyclability which can replace standard graphite anodes. In this effort to increase the energy storage capacity of LiBs, a wide range of electrode materials have been investigated including metal oxides, metal fluorides, silicon, germanium, tin and their alloys, and nanocomposites.12,13 Materials such as silicon,14−16 tin,17,18 antimony,19,20 and aluminum21 can form alloys with lithium metal at low potentials. These materials are considered promising candidates for anodes because of their high theoretical capacities (silicon ∼4200 mAh/g, germanium 1385 mAh/g, tin 993 mAh/g, antimony 660 mAh/g), and they have suitable equilibrium potential in the range of 1.0 and 0.3 V © 2016 American Chemical Society

SnO2 + 4Li+ + 4e− → Sn + 2Li 2O

(1)

Sn + x Li+ + x e− → LixSn

(2)

(0 ≤ x ≤ 4.4)

The first step (eq 1) is an irreversible reduction step, which forms a Li2O matrix surrounding tin metal. The formation of the Li2O matrix is considered to be an important step in reducing volume fluctuations in subsequent charge/discharge cycles.31 The second step (eq 2) represents an alloying mechanism, which is known to induce large volume fluctuations but is somewhat mitigated by the formation of the Li2O matrix. However, the practical application of bulk crystalline SnO2 is still hampered by volume changes that occur during the charge−discharge process.3,37,38 Tin disulfide (SnS2), an n-type semiconductor “layered compound” with a hexagonal cadmium iodide (CdI2 ) structure,39−41 could also serve as a potential high-capacity Received: September 17, 2015 Revised: January 13, 2016 Published: January 15, 2016 2036

DOI: 10.1021/acs.jpcc.5b09078 J. Phys. Chem. C 2016, 120, 2036−2046

Article

The Journal of Physical Chemistry C

coefficient. Section IV provides a summary of this work and concluding remarks.

anode material for Li-ion batteries. Due to its layered structure, SnS2 can host molecular guest species at the empty sites between its adjacent sulfur layers,42 in a fashion possibly similar to lithium intercalation in graphite. Therefore, it has been examined as an electrode material for Li-ion batteries.43−48 The advantage of this layered morphology in accommodating lithium has also been demonstrated in other metal dichalcogenides such as molybdenum sulfide.49 Other studies explored several redox metal oxides with capacities of ∼1000 mAh/g including MnO2, CoO, FeO, Fe2O3, NiO, CuO, Cu2O, Fe3O4, and Co3O4.50 Among most transition metal oxides, RuO2 is of particular interest because of its metallic conductivity, high chemical and thermal stability, catalytic activities, electrochemical redox properties, and field emitting behavior.51,52 It has been studied extensively as an electrode material for supercapacitor and exhibits remarkably high specific capacitance and a long cycle life.51−56 Several previous studies of lithium intercalation in RuO2 show high reversible capacities.57−59 The interest in RuO2 for LiB anodes increased tremendously after Balaya et al. reported on its electrochemical properties for LiBs in 2002.60 They demonstrated an exceptional reversibility in the first cycle over a wide voltage range (0.05−4.3 V), exhibiting a unique combination of high capacity (1130 mAh/g, corresponding to the storage of 5.6 mol of Li ions per mole of RuO2) and high initial Coulombic efficiency (98%). The mechanism of RuO2 lithiation has been explored by various experimental methods, including X-ray diffraction,60−62 selected area electron diffraction,60,61 high-resolution transmission electron microscopy (HR-TEM),60,61 magic-angle spinning 6Li NMR,62 Raman spectroscopy,60 and in situ transmission electron microscopy.63 The experimental observations are consistent with a two-step mechanism for lithium sorption, the first of which is simple intercalation of lithium ions in the RuO2 crystal lattice followed by conversion at higher capacities. As in the case of SnO2 (eq 1), the conversion reaction is of the form RuO2 + 4Li+ + 4e− → Ru + 2Li 2O

II. COMPUTATIONAL DETAILS We employed first-principles DFT calculations with Blöchl’s projector augmented wave (PAW) method65,66 and the generalized gradient approximation of Perdew−Burke−Ernzerhof (PBE)67 for the exchange-correlation interaction potentials. All calculations were carried out in a periodic working cell using a 7 × 7 × 7 Monkhorst−Pack k-point grid68 and an energy cutoff of 500 eV. The Vienna Ab Initio Simulations Package (VASP) was utilized for the DFT calculations.69 The initial working cells were taken from unit cells replicated into a 2 × 2 × 1 supercell (24 atoms8 tin or ruthenium atoms, 16 oxygen or sulfur atoms). RuO2 has the tetragonal rutile structure with P42/mnm symmetry, space group number 136, with a = b = 4.492 Å and c = 3.107 Å lattice parameters (JCPDS no. 21-1172) under normal conditions.70,71 SnO2 possesses the same tetragonal rutile-type unit cell structure as RuO2, with lattice constants of a = b = 4.7374 Å and c = 3.186 Å (JCPDS no. 41-1445).72 Due to layer alternations, SnS2 has the 2H, 4H, and 18R polytypes, which can be distinguished by their different layer stackings.73 The calculations were carried out for the 2H polytype of SnS2 which is reported to have lattice parameters a = b = 3.638 Å, and c = 5.880 Å (JCPDS no. 23-0677).74 Starting with these experimental lattice parameters, our unit cell optimizations allowed atomic positions and initial cell lattice constants to relax in order to determine the minimum energy structure. The optimized lattice parameters were a = b = 4.517 Å and c = 3.127 Å for RuO2 and a = b = 4.768 Å and c = 3.220 Å for SnO2, while SnS2 gave a = b = 3.696 Å and c = 5.918 Å, all of which compare favorably with experiment. Two models were used for computational studies of lithiation of these anode materials. The first started with supercells constructed from the optimized unit cells which corresponded to the formula (MX2)8, where M = Ru or Sn, and X = O or S, and successively added lithium to the cell, followed by structural optimization by full relaxation of atom positions as well as cell shape and size. The second type started with a (Li4MX2)8 cell in which the metal atoms were arranged as an island in a sea of Li2X, representing the stoichiometric limit of full conversion. Lithium atoms were then added or removed from this structure. The purpose of lithium removal was not to simulate a charge cycle, but to assess the thermodynamic favorability of such structures at capacities lower than 4:1 Li/M conversion limit. In other words, this was done to estimate the capacity at which the conversion structures become less favorable thermodynamically than the structures resulting from lithiation of the original pristine metal oxides or sulfide. The delithiation was carried out by removing lithium atoms in multiples of eight that were within 5 Å of metal atoms. Furthermore, the removed lithium atoms were chosen to maximize the distance between them. Once again, full relaxation of atom positions as well as cell shape and volume were carried out. The rationale for these models has been described in our previous work,64 and will be clarified below. The results from first model is referred to as “curve I” in the following section, while those from the second model is labeled as “curve II.” In both cases, lithium atoms were added in the cells in multiples of eight atoms via a hit and miss Monte Carlo procedure described in our previous work.64 Briefly, the insertion of each lithium atoms was carried out by choosing

(3)

Computational studies of Hassan et al. lend support to this two-step mechanism for low to moderate capacities.64 In spite of their high capacities, redox metal oxides such as RuO2 have low cyclability, and are often coated with carbon to improve their cyclability.12,13,60 Improving the cycle performance of electrode materials requires a detailed understanding of the mechanism by which the electrode materials operate and factors that contribute to their deficiencies. In this paper, we report the computational study of the mechanism of lithium sorption and the electrochemical properties of RuO2, SnO2, and SnS2 using plane wave density functional theory (DFT). Computational approaches, when validated by experiment, can provide valuable atomic level insights into the molecular mechanism of different material properties, which can provide avenues to explore or employ in improving capacity and cyclability. A detailed understanding of the thermodynamics and chemistry of electrode materials is essential in maximizing the energy density and capacity of battery materials. The remainder of this paper is organized as follows. In section II, we describe the computational approach taken. In section III, we discuss the computational discharge curves for the three materials in reference to a Li|Li+ electrode and the molecular level insights provided by the computational results including bond and free volume analysis and expansion 2037


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Figure 1. Calculated voltage profiles for the lithium sorption in (a) RuO2, (b) SnO2, and (c) SnS2. Lithium sorption in crystalline (MO2)8 and (MS2)8 are denoted as calculated curve I, while lithium sorption in M8 island structures is calculated curve II. Experimental first discharge curves are given for comparison for RuO2 from Kim et al. (ref 77), for SnO2 from Wang et al. (ref 78), and for SnS2 from Kim et al. (ref 44).

and tin metals, accompanied by the formation of Li2O.31,60,79 However, these mechanisms do not account for all of the observed capacity for these materials. For RuO2, the stochiometric capacity based on the conversion limit (eq 3) is 806 mAh/g, but capacities of 932,79 1027,77 and 1130 mAh/g,60,61 have been reported experimentally. Similarly, SnO2 has a theoretical first discharge capacity based on conversion and alloying (eqs 1 and 2) of 1494 mAh/g, while capacities over 2000 mAh/g have been recorded.78,80 The origin of these extra capacities will be discussed later. SnS2 is thought to behave in essentially the same manner as SnO2, with metal displacement and formation of Li2S.46 When comparing computed and experimental voltages, the role of activation polarization,81 which is associated with kinetic factors and the activation barriers that need to be overcome, must be taken into consideration. Because of this, in general, experimental voltages can be expected to be lower than computed values. However, methods such as galvanostatic intermittent titration technique (GITT)82,83 can remove some of the influence of activation polarization. In general, the GITT method increases the voltage with its degree depending on the relaxation time used. For instance, in the conversion regime in RuO2, where activation barrier is important, employing GITT with a relaxation time of 3 h increases the voltage by approximately 0.2−0.3 V, improving the agreement with our computed results.77 With these taken into account, Figure 1 shows qualitative agreement between the calculated and experimental discharge curves, and characteristic regions observed in experiments are also present in the calculations. For instance, the discharge curve for RuO2 has two voltage plateaus followed by a sloping region, similar to the three characteristic regions observed in the experimental discharge curve. SnO2 has the same sloping curve in experiment and computations, and the small plateau at 200−400 mAh/g found experimentally can be accounted for by considering the highest computed voltage at each capacity. The SnS2 experimental results have two plateaus before 600 mAh/g followed by a sloping curve, which are replicated to a degree by the calculated results. The higher voltage of curve II for 600 mAh/g and higher capacities indicate that computations predict all three materials to undergo conversion reactions into metal and Li2O or Li2S domains, which is consistent with available experimental data. The simulated annealing calculations described in the previous section resulted in minor changes to the voltages for the SnO2 system. For instance, annealing the 16 Li system used for curve I resulted in its voltage increasing by 0.08 V, and

100 000 random positions in the working cell. The position that was the most “open”, determined by the average distance to neighboring atoms including periodic images, was chosen for lithium insertion. This process was repeated eight times, each time resulting in the addition of one lithium to the cell. The resulting cell was used as the input structure for structural and energy optimization using the described DFT methodologies. We also carried out calculations for RuO2 and SnO2 with double the cell size for low to moderate lithiation to determine the effect of system size on voltages and cell volume as described later. Finally, to gauge the effectiveness of our algorithm to place lithium atoms in the SnO2 working cell, we carried out simulated annealing calculations on three different structures: 16 Li for curve I, 32 Li for curve II, and 72 Li for curve II. These included 15 ps molecular dynamics simulations at 500 K in the NVT ensemble, followed by full structural and cell relaxation. The voltage potential relative to a Li|Li+ electrode was calculated via the Nernst equation, in which the free energy differences of the system were approximated to be equivalent to the energy differences as has been done elsewhere:75,76 Φ(V) =

−[E Lix(MX )n − E(MX 2)n − xE Li] 2

xF

(4)

where M = Ru or Sn, X = O or S, and the E are the optimized energies of the associated systems. Each lithium is associated with one electron; therefore, the number of electrons involved in the process, x, is also the number of lithium atoms added to the solid, and F is Faraday’s constant. The energy of the lithium atom was taken as the energy per atom from a DFT unit cell optimization of bcc lithium metal.

III. RESULTS AND DISCUSSIONS III.A. Voltage Profiles. Voltage profiles for the first discharge of RuO2, SnO2, and SnS2 were obtained using the methodology outlined in the previous section and are shown in Figure 1. As described in the previous section, “curve I” denotes calculations starting from the pristine crystal structures and “curve II” denotes those started from the “island” structures resulting from conversion reactions. Experimental first discharge curves for these materials are also shown, with data taken from the literature.44,77,78 For all three systems, except for the low-capacity region, lithium absorption into the working cells containing the island structures become more favorable as indicated by the higher voltage of curve II around 600 mAh/g and higher capacities. This is consistent with the mechanisms proposed on the basis of experimental observations: the lithiation of RuO2 and SnO2 reduces them into ruthenium 2038


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Table 1. Comparison of the Voltage/Electromotive Force (EMF, in Volts) and Volume Change (in Percent) of M8O16 and M16O32 Systems Lix(RuO2)8

Lix(RuO2)16

Lix(SnO2)8

Lix(SnO2)16

Li/M ratio

EMF (V)

% vol change

EMF (V)

% vol change

EMF (V)

% vol change

EMF (V)

% vol change

1 2 3 4

1.68 1.39 1.30 1.24

16.3 68.1 91.5 127.8

1.70 1.39 1.30 1.26

16.9 68.6 91.8 133.3

1.79 1.35 1.23 1.16

46.2 82.5 115.6 149.9

1.69 1.36 1.24 1.15

42.4 86.4 115.8 154.1

annealing the 32 Li and 72 Li systems used for curve II resulted in a small voltage decrease of 0.02 and 0.03 V, respectively. This shows that the curve I voltages may be slightly underestimated, but the difference in voltages should not have a significant impact on the qualitative nature of the results. Furthermore, the annealing calculations show that the Monte Carlo lithium placement methodology appears to do a very good job of finding a low-energy configuration. In an effort to assess if the M8O16 working cells is sufficient to capture the inherent physicochemical properties of the systems, we explored the effect of the size of the system on voltage and volume change for low to moderate lithium content. The discharge curve for RuO2 and SnO2 were explored using a M16O32 working cell (twice the size of the original M8O16 working cell). Presented in Table 1 is the comparison of voltages and percentage volume change for the two different cell sizes. As can be observed, doubling the cell size does not significantly change the voltage or volume of the structures, producing similar results. Reaching high lithium content, one of the aims of this paper, with the M16O32 cell would require far more lithium atoms and drastically increase required computational resources. In light of the smaller M8O16 and larger M16O32 resulting in the same qualitative results and trends for moderate lithiation, we therefore conclude that the former was sufficient to computationally characterize the lithiation of the considered systems. III.B. Analysis of Structures. A few representative structures used to calculate curve I are given in Figure 2. It is well-accepted from experimental as well as computational evidence,64,84 that RuO2 acts as an intercalation material up to ∼1:1 ratio of Li/Ru, accommodating Li into the RuO2 lattice with minimal disruption to the crystalline structure. This is evident in the top row of Figure 2. However, the rutile crystalline structure gets distorted beyond the 1:1 ratio. This disruption of the crystalline structure continues as more lithium atoms are added, resulting in a substantial distortion by the time the stoichiometric conversion limit (see eq 3) of 4:1 Li/ Ru ratio (32 Li) is reached. The tendency of the ruthenium atoms to aggregate into metallic ruthenium is evident at a 4:1 Li/Ru ratio (32 Li) and beyond, while the lithium and oxygen atoms appear to cluster into relatively ruthenium-free domains, which is an expected feature of the conversion mechanism. Previous charge analysis based on Bader charges85,86 has also shown that the ruthenium charges decrease with higher lithium content, approaching zero, consistent with the conversion mechanism.64 However, complete segregation of the ruthenium and Li2O phases does not occur in these energy minimization computations, presumably because of the energy barriers which need to be surmounted to reach the phase-separated final state observed in in situ HR-TEM experiments.63 The SnO2 crystal lattice suffers some structural disruption even at low lithiation, and the distortion gets progressively greater as the Li/Sn ratio increases.

Figure 2. Selected representative optimized structures for (MX2)8Lin where M = Ru or Sn and X = O or S. Ruthenium atoms are dark green (teal), tin atoms are blue, oxygens are red, sulfurs are yellow, and lithium atoms are purple.

However, in marked contrast to RuO2, SnO2 does not show significant aggregation of tin atoms with higher lithium content. In the case of SnS2, the added Li migrates away from tin and appears to form Li2S even at low capacities. This allows the tin atoms initially to cluster together at intermediate capacities, but as one goes beyond the conversion limit of 4:1 (32 Li), the clusters disintegrate and become intermingled with added Li. As already mentioned above, the computational methodology used to calculate curve I does not lead to structures that represent full conversion of the electrode material, presumably because of the energy barriers between the various minima, which would likely require substantial annealing to overcome. Even experimentally, full conversion is not always reached in the first discharge; a combination of Li2O and LixRuO2 domains are found with in situ TEM images of RuO2 after the first discharge.63 The model underlying curve II of Figure 1, in which the input structure itself contains a metal island embedded in a sea of Li2X (X = O or S), provides a starting 2039


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Li, and 3.5 Å for S−Li. M represents metal atoms, ruthenium, or tin. These cutoff distances were chosen by calculating the radial distribution functions between different atom types, finding the first peak in this function, and placing the cutoff just past this first peak, as has been described previously.64 The number of nearest neighbors serves as an indication of how strongly the constituent atoms in the system interact as the lithium content increases. Figure 4 gives the number of atoms of a certain type surrounding each atom type for RuO2, SnO2, and SnS2. The number of Ru−Ru neighbors is roughly constant, initially increasing from ∼3 to ∼4 and then decreasing to ∼2 at the highest ratios of Li/Ru considered. In comparison, the tin-based systems start out with low numbers of Sn−Sn neighbors and the number rapidly drops further as lithium atoms are accommodated in the cell. The number of Ru−O nearest neighbors decreases upon lithiation. This same reduction can be observed for Sn−O and Sn−S, signifying the separation and breaking of bonding between the metal (ruthenium and tin) and oxygen or sulfur atoms. This is accompanied by a simultaneous increase in the number of Li− O and Li−S nearest neighbors until the 32 Li stoichiometric limit of conversion is reached, consistent with the formation of Li2O and Li2S domains. At high Li concentrations, i.e., beyond a 4:1 Li/Sn ratio (32 Li), the number of Sn−O and Sn−S nearest neighbors decreases while the Li−Sn value increases dramatically, indicating alloying. An alternate diagnostic of the nature of bonding in the cell is provided by the average distance to nearest neighbors of a specified type, which is given in Figure 5. The Ru−Ru distance is roughly constant, independent of lithium content, while the Sn−Sn distances increase for both SnO2 and SnS2. The M−O and M−S distances increase with increasing lithium content, an indication of the separation of the metal and oxygen/sulfur phases, consistent with the decrease in the number of M−O and M−S neighbors with increasing lithium content seen in Figure 4. Of particular interest is the fact that the M−Li distance reaches approximately the same value at high lithium content in all three cases. The accommodation of lithium beyond the conversion limit of a 4:1 ratio of Li/M (32 Li in this case) is of particular interest. In the case of RuO2, an interfacial storage mechanism has already been described for lithium sorption at high capacities,62,87 which was confirmed by computational results in our previous work.64 One indication of interfacial storage in RuO2 is that the O−Li average distance steadily increases with

point for investigating the lithiation behavior of the anode material at and beyond the stoichiometric conversion limit (Figure 3).

Figure 3. Selected representative optimized structures for the absorption of lithium atoms onto Ru8 and Sn8 islands surrounded by Li2O/Li2S matrix. Ruthenium atoms are dark green (teal), tin is blue, oxygen is red, sulfur is yellow, and lithium atoms are purple.

III.C. Bond Analysis. In order to further assess the local environment surrounding each atom, we calculated the average number and type of atoms close to each atom type. These analyses were done for the systems used to generate curve II so that the results reported are for the most stable structures at high lithium content. Atoms were considered to be neighbors if their separation was less than a certain cutoff: 2.8 Å for Ru−Ru, 3.1 Å for Sn−Sn, 2.5 Å for M−O, 3.5 Å for M−Li, 2.5 Å for O−

Figure 4. Number of nearest neighbor of the constituent atoms in RuO2, SnO2, and SnS2 as number of lithium atoms increases. 2040


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Figure 5. Average distance to nearest neighbors of specified type in RuO2, SnO2, and SnS2 as number of lithium atoms increases in the island system.

this work at the same lithium content. In the present work, a steep increase in volume change occurs beyond a 2:1 Li/Sn, whereas this steep volume change manifested beyond a 1:1 ratio in Liu’s work.89 III.E. Free Volume Analysis. To discern the extent to which different phases of the working cell contribute to the volume increases, a free volume analysis was carried out for the various systems. The free volume was determined by finding the number of 1 Å diameter spheres that can be placed in different regions in the working cell without overlapping with existing atoms in the cell or each other. The structures used for this analysis were taken from those used to calculate the voltages using the atomic van der Waals radii from the literature.90 An overlap was defined if the distance between the centers of the inserted sphere and another sphere or atom was less than the arithmetic mean of their van der Waals radii. So, for example, two spheres were considered to overlap if their centers were within 1.0 Å of each other while for ruthenium (which has a van der Waals radii of 1.90 Å), a sphere was considered overlapping with it when their interatomic distance was less than 1.45 Å (or [1.90 Å + 1.00 Å]/2). The environment of each sphere added to the working cell was characterized based on the number and types of atoms within twice the distance used to denote an overlap. If two or more metal atoms (Ru/Sn) were within the van der Waals radius of a sphere, it was considered to be in “bulk metal”, while if two or more oxygen or sulfur atoms were within the radius, it was considered “bulk O/S”. Otherwise, the spheres were denoted to be at the “interface”. A more detailed description of the method is given in previous work.64 The measure of the free volume in the different regions is given in Figure 7 for RuO2, SnO2, and SnS2 with respect to the number of lithium atoms. For RuO2, the available space in the bulk metal remains relatively low at all lithium amounts while the space near bulk tin increases with higher lithium sorption for SnO2 and SnS2. The free volume in bulk oxide/sulfide remains fairly constant with respect to lithium loading for all three materials, although the free volume in SnS2 appears to be higher than the two oxides even at low lithiation. The space available in the interface increases for RuO2 when 40 or more lithium atoms (≥5:1 ratio of Li/Ru) are present, which is consistent with an interfacial storage mechanism. For SnO2 and SnS2, when 64 or more lithium atoms (≥8:1 ratio of Li/M) are present, the space available in the interface increases also. The increase in free volume in bulk metal for the tin-based materials is consistent with an alloying mechanism, as suggested previously,46 but the free volume at the interface between the

lithium content while the Ru−Li distance initially decreases and then remains constant, suggesting that the added Li are entering the interfacial region between the ruthenium and the Li2O phases formed during conversion. In contrast, an alloying mechanism has been proposed for tin-based anodes (see eq 2).31,46 However, the similarity between the trends in Ru−O, Sn−O, and Sn−S distances, as well as the trends in the M−Li distances at high capacities suggests that similar mechanisms may be responsible for the “superstoichiometric” capacity in all three materials. We will examine this aspect in greater detail below. III.D. Volume Profiles. As pointed out earlier, the practical application of many new electrode materials has been hampered by large volume fluctuations with lithiation.3,37,38,88 Therefore, the percentage of volume change as a function of lithiation could be a vital measure in assessing structural stability/integrity of an electrode material. The ratios of the volume of converged lithiated structures used to calculate the voltage profiles, with the volume of the crystalline working cells before lithiation, are given in Figure 6. SnO2 has the largest

Figure 6. Percent volume change for RuO2, SnO2, and SnS2 as number of lithium atoms absorbed increases.

volume expansion, but the rate of expansion of RuO2 is similar to SnO2. SnS2, though, expands at a far lesser degree than the other two, being around two-thirds as much as the other two systems at the highest capacity studied. In addition to having less overall volume expansion than the other two materials, SnS2 shows exceptionally small volume expansion at low lithium content. This has also been reported in the GGA-PBE results of Liu et al., which showed an 11.74% expansion at x = 1,89 compared to the 19% expansion shown in 2041


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Figure 7. Number of spheres in the different regions, as the number of lithium atoms increases in RuO2, SnO2, and SnS2. The legends are explained in text.

III.F. Lithium Environmental Analysis. Using criteria similar to those to characterize free volume, the lithium environment is analyzed and described in Figure 9. If two or more metal (Ru/Sn) atoms were within the arithmetic mean of the van der Waals radii of the lithium and the specific metal atom, then the environment of the lithium atom was characterized as being in “bulk metal”. If two or more oxygen or sulfur atoms were within the specified distance, the lithium environment was characterized as “bulk O/S”, and the remaining lithium atoms were characterized as “interface”. As evident from Figure 9, for all systems, the number of lithium atoms in the bulk O/S increases steadily until the full conversion limit of 4:1 Li/M ratio or 2:1 Li/O/S ratio is reached (32 Li in our cells) while hardly any lithium appears to enter the interfacial region. This is consistent with the conversion mechanism. At higher capacities, the lithium content in the bulk metal region increases rapidly for the SnO2 and SnS2 systems, showing an alloying mechanism. For RuO2, there is a smaller, but noticeable increase in the number of lithium atoms in the bulk region. The criterion for atoms to be considered bulk metal, that two or more metal atoms be within the van der Waals radius of a lithium atom, will encompass lithium atoms on the surface of the ruthenium island. For all three materials, there is also a signif icant increase in lithium sorption into the interfacial region as the number of lithium atoms increases beyond the conversion limit. Thus, we conclude that two mechanisms for “superstoichiometric” capacity (i.e., beyond a 4:1 Li/M ratio) are present in SnO2 and SnS2: an alloying mechanism is primary for these two, but also common for all three materials is an interfacial storage mechanism, especially at higher capacity. A color-coded map of the different regions for the highest lithiated form/structure (80 Li) for each material is presented in Figure 10. The dark blue regions represent bulk metal regions (within the van der Waals radius of two or more metal atoms), red is for bulk O/S region (within the van der Waals radius of two or more O/S atoms), and green is for the interfacial region. It clearly shows that interfacial accumulation is more prevalent in RuO2 than the tin-based systems. Both SnO2 and SnS2 show a much smaller accumulation of lithium atoms in the interfacial regions, with SnO2 having a relatively larger interfacial region than SnS2. This representation shows that even for the tin-based materials that operate primarily by alloying, a small interfacial storage of lithium could be present at the highest capacities.

metal and Li2O/Li2S phases continues to increase as well, albeit at higher lithium content. This effect, which is related to interfacial storage of lithium (see below), has been described for the conversion material RuO2 in our previous work,64 but is reported here for the first time for the alloying materials SnO2 and SnS2. The large free volume originally present in bulk sulfur in SnS2 compared to the bulk oxygen in SnO2 is probably responsible for the lower percentage of volume expansion in SnS2 with respect to lithiation (Figure 6). It is apparent that the O/S region’s free volume changes only slightly with lithiation and serves as a buffer for total volume expansion. Since sulfur has a larger volume (and it stays constant), the total volume percentage does not change as much with lithiation for SnS2. The degree of expansion of the core Ru/Sn islands (bulk metal regions) can be reflected in the radius of gyration, Rg, of the structures as shown in Figure 8. The quantity Rg2 is defined as

Figure 8. Radius of gyration (Rg) in Å of the RuO2, SnO2, and SnS2 island structures with increasing lithium absorption.

the average of the squared distances of metal atoms from their center of mass (COM): Rg2 =

1 Nmetal

Nmetal

∑ i=1

(ri − rCOM)2

(5)

where Nmetal is the number of metal atoms. The tin radius of gyration increases the most for SnO2, while for a somewhat lesser degree for SnS2. This is consistent with the lower overall volume expansion for the SnS2 system. The ruthenium radius of gyration decreases initially with lithiation, followed by a small expansion at higher lithium content. This is expected as no significant Li/Ru alloying occurs. 2042


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Figure 9. Number of Li absorbed in the different regions, as the number of lithium atoms increases in RuO2, SnO2, and SnS2 (panels a, b, c, respectively). The legends are explained in text.

ranged from −0.2e to 0.03e at in the interfacial region. These results are consistent with the conclusion that, at high capacities, lithium is stored in almost elemental form in the interfacial regions.

IV. CONCLUSION Using ab initio computational modeling, we have investigated the lithiation mechanism and structural changes in RuO2, SnO2, and SnS2 due to lithiation. The computational voltage curves qualitatively agree with experimental discharge curves. Bond, free volume, and lithium environmental analyze carried out for the structures showed that, up to a certain capacity, all systems followed a conversion mechanism in which MO2 or MS2, with M representing ruthenium or tin, converted into pure M and Li2O or Li2S. This mechanism was prevalent until a 4:1 Li/M ratio was reached. Beyond this limit, the tin-based systems accommodate additional lithium through an alloying mechanism, which was not present in RuO2. The sorption of lithium into RuO2 beyond the 4:1 Li/M conversion ratio was found to be at the interface between the ruthenium and Li2O phases.64 It was mentioned in connection to Figure 1 that the maximum capacity of ∼2100 mAh/g recorded for SnO2 exceeds the combined conversion and alloying limit of 1494 mAh/g. To the best of our knowledge, the analysis presented in this paper is the first suggestion that lithium storage at the phase boundary between the Li−Sn alloy and Li2O or Li2S may account for the observed extra capacity in tin-based anode materials. This explanation for extra capacity is entirely based on properties of the electrode material because our simulations do not include the solid−electrolyte interface (SEI).91−95 Modeling the SEI using first-principles DFT would be extremely challenging because of the complexity of the system that needs to be modeled and also the lack of information about the detailed composition of the SEI for a given electrolyte and electrode material. It should be assumed that the SEI layer contains some lithium which makes a non-negligible contribution to the observed total capacity. However, experiments and computations suggest that lithium in the SEI cannot account for all of the superstoichiometric capacity nor the voltage at which the extra capacity is observed. Evidence in favor of interfacial lithium storage within the electrode material as the explanation for the excess capacity has been accumulating since the idea was first suggested in 2003 by Balaya et al.60 who recorded a lithium capacity of 1130 mAh/g for RuO2. Computational support for this suggestion was offered in 2006 using an idealized model for a Ti/Li2O

Figure 10. Color-mapped regions of the highest lithiated (80 Li) structures of RuO2, SnO2, and SnS2.

III.G. Bader Charge Analysis. We analyzed how the charge distribution of the systems change for select lithiation levels utilizing Bader charge partitioning scheme.85,86 This method partitions a molecular space into atomic fragments such that the electron density gradient vanishes at every point on the dividing surfaces. This approach has been used to describe the evolution of charges in lithiated RuO2 in earlier work.64 The average charges for each ionic (atomic) species in the working cell and their various lithiated forms in the first discharge curve are provided in Table 2. The charges for Sn decrease drastically in Table 2. Average Bader Charges of Atomic Species in the Different Lithiated SnO2 and SnO2 Systems 0 Li

32 Li

Sn O Li

2.510 −1.255

0.112 −1.365 0.670

Sn S Li

1.239 −0.619

0.083 −0.764 0.371

48 Li SnO2 −0.059 −1.395 0.579 SnS2 0.065 −0.743 0.237

64 Li

80 Li

−0.134 −1.442 −0.047

−0.320 −1.518 −0.028

−0.507 −0.903 0.289

−0.176 −0.657 0.149

both SnO2 and SnS2 as more lithium atoms are added, becoming negative at the highest lithium content. We also examined the individual charges on the lithium atoms depending on its environment. Lithium atoms in the Li2O phase had charges between 0.6e and 0.8e, while lithium alloyed with tin had fairly neutral charges, ranging between −0.1e and 0.2e. Lithium atoms in the interfacial region between Li2O and the Li−Sn alloy phase had charges around 0.3e. For the SnS2 system, the lithium charges in Li2S ranged from 0.3e to 0.6e and 2043


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The Journal of Physical Chemistry C interface.96 In 2009, the analysis of RuO2 lithiation by Bekaert et al.62 using magic-angle spinning 6Li NMR indicated that the lithium incorporated at the highest capacitieswhich is also the lithium that is removed first in the charging cyclehad different characteristics from that in the SEI layer. The analysis in ref 64 presented the first computational evidence based on a realistic model of lithiated RuO2 in support of the interfacial storage mechanism. Recently, Kim et al., using a combination of experimental methods, has also concluded that the extra capacity of RuO2 is due to “Li storage in the grain boundary between nanosized Ru metal and Li2O”.77 The present work offers further support to this conclusion. Moreover, the evidence presented here for interfacial storage of lithium in conversion as well as alloying materials suggests that this mechanism for excess capacity may be quite universal in highcapacity oxides and sulfides and possibly other materials.

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(13) Hayner, C. M.; Zhao, X.; Kung, H. H. Materials for Rechargeable Lithium-Ion Batteries. Annu. Rev. Chem. Biomol. Eng. 2012, 3, 445−471. (14) Jung, H.; Park, M.; Yoon, Y.-G.; Kim, G.-B.; Joo, S.-K. Amorphous Silicon Anode for Lithium-Ion Rechargeable Batteries. J. Power Sources 2003, 115, 346−351. (15) Cho, J. Porous Si Anode Materials for Lithium Rechargeable Batteries. J. Mater. Chem. 2010, 20, 4009−4014. (16) Ohara, S.; Suzuki, J.; Sekine, K.; Takamura, T. A Thin Film Silicon Anode for Li-Ion Batteries Having a Very Large Specific Capacity and Long Cycle Life. J. Power Sources 2004, 136, 303−306. (17) Valvo, M.; Lafont, U.; Munao, D.; Kelder, E. M. Electrospraying-Assisted Synthesis of Tin Nanoparticles for Li-Ion Battery Electrodes. J. Power Sources 2009, 189, 297−302. (18) Winter, M.; Besenhard, J. O. Electrochemical Lithiation of Tin and Tin-Based Intermetallics and Composites. Electrochim. Acta 1999, 45, 31−50. (19) Kim, H.; Cho, J. Template Synthesis of Hollow Sb Nanoparticles as a High-Performance Lithium Battery Anode Material. Chem. Mater. 2008, 20, 1679−1681. (20) Pralong, V.; Leriche, J. B.; Beaudoin, B.; Naudin, E.; Morcrette, M.; Tarascon, J. M. Electrochemical Study of Nanometer Co3O4, Co, CoSb3 and Sb Thin Films toward Lithium. Solid State Ionics 2004, 166, 295−305. (21) Hamon, Y.; Brousse, T.; Jousse, F.; Topart, P.; Buvat, P.; Schleich, D. M. Aluminum Negative Electrode in Lithium Ion Batteries. J. Power Sources 2001, 97−98, 185−187. (22) Huggins, R. A. Lithium Alloy Negative Electrodes. J. Power Sources 1999, 81−82, 13−19. (23) Obrovac, M. N.; Christensen, L.; Le, D. B.; Dahn, J. R. Alloy Design for Lithium-Ion Battery Anodes. J. Electrochem. Soc. 2007, 154, A849−A855. (24) Tirado, J. L. Inorganic Materials for the Negative Electrode of Lithium-Ion Batteries: State-of-the-Art and Future Prospects. Mater. Sci. Eng., R 2003, 40, 103−136. (25) Besenhard, J. O.; Yang, J.; Winter, M. Will Advanced LithiumAlloy Anodes Have a Chance in Lithium-Ion Batteries? J. Power Sources 1997, 68, 87−90. (26) Huggins, R. A. Polyphase Alloys as Rechargeable Electrodes in Advanced Battery Systems. J. Power Sources 1988, 22, 341−350. (27) Zhang, W.-J. Lithium Insertion/Extraction Mechanism in Alloy Anodes for Lithium-Ion Batteries. J. Power Sources 2011, 196, 877− 885. (28) Cabana, J.; Monconduit, L.; Larcher, D.; Palacín, M. R. Beyond Intercalation-Based Li-Ion Batteries: The State of the Art and Challenges of Electrode Materials Reacting through Conversion Reactions. Adv. Mater. 2010, 22, E170−E192. (29) Emsley, J. Nature’s Building Blocks: An A-Z Guide to the Elements; Oxford University Press: Oxford, U.K., 2011. (30) Besenhard, J.; Yang, J.; Winter, M. Will Advanced Lithium-Alloy Anodes Have a Chance in Lithium-Ion Batteries? J. Power Sources 1997, 68, 87−90. (31) Courtney, I. A.; Dahn, J. Electrochemical and in Situ X-Ray Diffraction Studies of the Reaction of Lithium with Tin Oxide Composites. J. Electrochem. Soc. 1997, 144, 2045−2052. (32) Li, N.; Martin, C. R.; Scrosati, B. A High-Rate, High-Capacity, Nanostructured Tin Oxide Electrode. Electrochem. Solid-State Lett. 2000, 3, 316−318. (33) Idota, Y.; Kubota, T.; Matsufuji, A.; Maekawa, Y.; Miyasaka, T. Tin-Based Amorphous Oxide: A High-Capacity Lithium-Ion-Storage Material. Science 1997, 276, 1395−1397. (34) Courtney, I. A.; Dahn, J. R. Electrochemical and in Situ X-Ray Diffraction Studies of the Reaction of Lithium with Tin Oxide Composites. J. Electrochem. Soc. 1997, 144, 2045−2052. (35) Kim, C.; Noh, M.; Choi, M.; Cho, J.; Park, B. Critical Size of a Nano SnO2 Electrode for Li-Secondary Battery. Chem. Mater. 2005, 17, 3297−3301.

AUTHOR INFORMATION

Corresponding Authors

*Phone: +1-318-257-4304. Fax: +1-318-257-2562. E-mail: [email protected]. *Phone: +1-318-257-2345. Fax: +1-318-257-3823. E-mail: [email protected]. Notes

The authors declare no competing financial interest.

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ACKNOWLEDGMENTS This work is supported by the National Science Foundation through EPS-1003897 and EPS-1006891. Grants of computer time from the Louisiana Optical Network Initiative (LONI) are gratefully acknowledged. This research also used resources of the National Energy Research Scientific Computing Center, a DOE Office of Science User Facility supported by the Office of Science of the U.S. Department of Energy under contract no. DE-AC02-05CH11231.

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REFERENCES

(1) Goodenough, J. B.; Kim, Y. Challenges for Rechargeable Li Batteries. Chem. Mater. 2010, 22, 587−603. (2) Palacin, M. R. Recent Advances in Rechargeable Battery Materials: A Chemist’s Perspective. Chem. Soc. Rev. 2009, 38, 2565− 2575. (3) Tarascon, J. M.; Armand, M. Issues and Challenges Facing Rechargeable Lithium Batteries. Nature 2001, 414, 359−367. (4) Oyama, N.; Tatsuma, T.; Sato, T.; Sotomura, T. DimercaptanPolyaniline Composite Electrodes for Lithium Batteries with High Energy Density. Nature 1995, 373, 598−600. (5) Scrosati, B. Challenge of Portable Power. Nature 1995, 373, 557− 558. (6) Hassoun, J.; Reale, P.; Scrosati, B. Recent Advances in Liquid and Polymer Lithium-Ion Batteries. J. Mater. Chem. 2007, 17, 3668−3677. (7) Thomas, J. Lithium Batteries: A Spectacularly Reactive Cathode. Nat. Mater. 2003, 2, 705−706. (8) Buqa, H.; Goers, D.; Holzapfel, M.; Spahr, M. E.; Novák, P. High Rate Capability of Graphite Negative Electrodes for Lithium-Ion Batteries. J. Electrochem. Soc. 2005, 152, A474−A481. (9) Megahed, S.; Scrosati, B. Lithium-Ion Rechargeable Batteries. J. Power Sources 1994, 51, 79−104. (10) Wu, Y. P.; Rahm, E.; Holze, R. Carbon Anode Materials for Lithium Ion Batteries. J. Power Sources 2003, 114, 228−236. (11) Armand, M.; Tarascon, J. M. Building Better Batteries. Nature 2008, 451, 652. (12) Manthiram, A. Materials Challenges and Opportunities of Lithium Ion Batteries. J. Phys. Chem. Lett. 2011, 2, 176−184. 2044


Article

The Journal of Physical Chemistry C (36) Lou, X. W.; Wang, Y.; Yuan, C.; Lee, J. Y.; Archer, L. A. Template-Free Synthesis of SnO2 Hollow Nanostructures with High Lithium Storage Capacity. Adv. Mater. 2006, 18, 2325−2329. (37) Lou, X. W.; Deng, D.; Lee, J. Y.; Archer, L. A. Preparation of SnO2/Carbon Composite Hollow Spheres and Their Lithium Storage Properties. Chem. Mater. 2008, 20, 6562−6566. (38) Larcher, D.; Beattie, S.; Morcrette, M.; Edstrom, K.; Jumas, J.C.; Tarascon, J.-M. Recent Findings and Prospects in the Field of Pure Metals as Negative Electrodes for Li-Ion Batteries. J. Mater. Chem. 2007, 17, 3759−3772. (39) Greenaway, D. L.; Nitsche, R. Preparation and Optical Properties of Group IV−VI2 Chalcogenides Having the CdI2 Structure. J. Phys. Chem. Solids 1965, 26, 1445−1458. (40) Deshpande, N. G.; Sagade, A. A.; Gudage, Y. G.; Lokhande, C. D.; Sharma, R. Growth and Characterization of Tin Disulfide (SnS2) Thin Film Deposited by Successive Ionic Layer Adsorption and Reaction (Silar) Technique. J. Alloys Compd. 2007, 436, 421−426. (41) Ray, S. C.; Karanjai, M. K.; DasGupta, D. Structure and Photoconductive Properties of Dip-Deposited SnS and SnS2 Thin Films and Their Conversion to Tin Dioxide by Annealing in Air. Thin Solid Films 1999, 350, 72−78. (42) Schlaf, R.; Armstrong, N. R.; Parkinson, B. A.; Pettenkofer, C.; Jaegermann, W. Van Der Waals Epitaxy of the Layered Semiconductors SnSe2 and SnS2: Morphology and Growth Modes. Surf. Sci. 1997, 385, 1−14. (43) Momma, T.; Shiraishi, N.; Yoshizawa, A.; Osaka, T.; Gedanken, A.; Zhu, J.; Sominski, L. SnS2 Anode for Rechargeable Lithium Battery. J. Power Sources 2001, 97−98, 198−200. (44) Kim, T.-J.; Kim, C.; Son, D.; Choi, M.; Park, B. Novel SnS2Nanosheet Anodes for Lithium-Ion Batteries. J. Power Sources 2007, 167, 529−535. (45) Mukaibo, H.; Yoshizawa, A.; Momma, T.; Osaka, T. Particle Size and Performance of SnS2 Anodes for Rechargeable Lithium Batteries. J. Power Sources 2003, 119−121, 60−63. (46) Brousse, T.; Lee, S. M.; Pasquereau, L.; Defives, D.; Schleich, D. M. Composite Negative Electrodes for Lithium Ion Cells. Solid State Ionics 1998, 113−115, 51−56. (47) Seo, J.-w.; Jang, J.-t.; Park, S.-w.; Kim, C.; Park, B.; Cheon, J. Two-Dimensional SnS2 Nanoplates with Extraordinary High Discharge Capacity for Lithium Ion Batteries. Adv. Mater. 2008, 20, 4269−4273. (48) Zhong, H.; Yang, G.; Song, H.; Liao, Q.; Cui, H.; Shen, P.; Wang, C.-X. Vertically Aligned Graphene-Like SnS2 Ultrathin Nanosheet Arrays: Excellent Energy Storage, Catalysis, Photoconduction, and Field-Emitting Performances. J. Phys. Chem. C 2012, 116, 9319−9326. (49) Benavente, E.; Santa Ana, M. A.; Mendizábal, F.; González, G. Intercalation Chemistry of Molybdenum Disulfide. Coord. Chem. Rev. 2002, 224, 87−109. (50) Poizot, P.; Laruelle, S.; Grugeon, S.; Dupont, L.; Tarascon, J. M. Nano-Sized Transition-Metal Oxides as Negative-Electrode Materials for Lithium-Ion Batteries. Nature 2000, 407, 496−499. (51) Wang, X.; Gordon, R. G. High-Quality Epitaxy of Ruthenium Dioxide, RuO2, on Rutile Titanium Dioxide, TiO2, by Pulsed Chemical Vapor Deposition. Cryst. Growth Des. 2013, 13, 1316−1321. (52) Gujar, T.; Kim, W.-Y.; Puspitasari, I.; Jung, K.-D.; Joo, O.-S. Electrochemically Deposited Nanograin Ruthenium Oxide as a Pseudocapacitive Electrode. Int. J. Electrochem. Sci. 2007, 2, 666−673. (53) Jirkovský, J.; Makarova, M.; Krtil, P. The Effect of Coherent Domain Size on the Insertion Activity of Nanocrystalline RuO2. J. Electrochem. Soc. 2005, 152, A1613−A1619. (54) Patake, V.; Lokhande, C.; Joo, O. S. Electrodeposited Ruthenium Oxide Thin Films for Supercapacitor: Effect of Surface Treatments. Appl. Surf. Sci. 2009, 255, 4192−4196. (55) Lytle, J. C.; Rhodes, C. P.; Long, J. W.; Pettigrew, K. A.; Stroud, R. M.; Rolison, D. R. The Importance of Combining Disorder with Order for Li-Ion Insertion into Cryogenically Prepared Nanoscopic Ruthenia. J. Mater. Chem. 2007, 17, 1292−1299.

(56) Kuratani, K.; Kiyobayashi, T.; Kuriyama, N. Influence of the Mesoporous Structure on Capacitance of the RuO2 Electrode. J. Power Sources 2009, 189, 1284−1291. (57) Murphy, D.; Di Salvo, F.; Carides, J.; Waszczak, J. Topochemical Reactions of Rutile Related Structures with Lithium. Mater. Res. Bull. 1978, 13, 1395−1402. (58) Dalard, P.; Deroo, D.; Foscallo, D.; Mouliom, C. A Galvanostatic Study of Lithium Intercalation in RuO2. Solid State Ionics 1985, 15, 91−94. (59) Armand, M.; Dalard, F.; Deroo, D.; Mouliom, C. Modelling the Voltammetric Study of Intercalation in a Host Structure: Application to Lithium Intercalation in RuO2. Solid State Ionics 1985, 15, 205−210. (60) Balaya, P.; Li, H.; Kienle, L.; Maier, J. Fully Reversible Homogeneous and Heterogeneous Li Storage in RuO2 with High Capacity. Adv. Funct. Mater. 2003, 13, 621−625. (61) Delmer, O.; Balaya, P.; Kienle, L.; Maier, J. Enhanced Potential of Amorphous Electrode Materials: Case Study of RuO2. Adv. Mater. 2008, 20, 501−505. (62) Bekaert, E.; Balaya, P.; Murugavel, S.; Maier, J.; Ménétrier, M. 6 Li MAS NMR Investigation of Electrochemical Lithiation of RuO2: Evidence for an Interfacial Storage Mechanism. Chem. Mater. 2009, 21, 856−861. (63) Gregorczyk, K. E.; Liu, Y.; Sullivan, J. P.; Rubloff, G. W. In Situ Transmission Electron Microscopy Study of Electrochemical Lithiation and Delithiation Cycling of the Conversion Anode RuO2. ACS Nano 2013, 7, 6354−6360. (64) Hassan, A. S.; Navulla, A.; Meda, L.; Ramachandran, B. R.; Wick, C. D. Molecular Mechanisms for the Lithiation of Ruthenium Oxide Nanoplates as Lithium-Ion Battery Anode Materials: An Experimentally Motivated Computational Study. J. Phys. Chem. C 2015, 119, 9705−9713. (65) Blöchl, P. E. Projector Augmented-Wave Method. Phys. Rev. B: Condens. Matter Mater. Phys. 1994, 50, 17953−17979. (66) Kresse, G.; Joubert, D. From Ultrasoft Pseudopotentials to the Projector Augmented-Wave Method. Phys. Rev. B: Condens. Matter Mater. Phys. 1999, 59, 1758−1775. (67) Perdew, J. P.; Burke, K.; Ernzerhof, M. Generalized Gradient Approximation Made Simple. Phys. Rev. Lett. 1996, 77, 3865−3868. (68) Monkhorst, H. J.; Pack, J. D. Special Points for Brillouin-Zone Integrations. Phys. Rev. B 1976, 13, 5188−5192. (69) Kresse, G.; Furthmüller, J. Efficient Iterative Schemes for Ab Initio Total-Energy Calculations Using a Plane-Wave Basis Set. Phys. Rev. B: Condens. Matter Mater. Phys. 1996, 54, 11169−11186. (70) Baur, W. H.; Khan, A. A. Rutile-Type Compounds. IV. SiO2, GeO2 and a Comparison with Other Rutile-Type Structures. Acta Crystallogr., Sect. B: Struct. Crystallogr. Cryst. Chem. 1971, 27, 2133− 2139. (71) Haines, J.; Léger, J. M.; Schulte, O. Pa Modified Fluorite-Type Structures in Metal Dioxides at High Pressure. Science 1996, 271, 629−631. (72) Batzill, M.; Diebold, U. The Surface and Materials Science of Tin Oxide. Prog. Surf. Sci. 2005, 79, 47−154. (73) Pałosz, B.; Steurer, W.; Schulz, H. Refinement of SnS2 Polytypes 2H, 4H and 18R. Acta Crystallogr., Sect. B: Struct. Sci. 1990, 46, 449− 455. (74) Al-Alamy, F. A. S.; Balchin, A. A.; White, M. The Expansivities and the Thermal Degradation of Some Layer Compounds. J. Mater. Sci. 1977, 12, 2037−2042. (75) Bard, A. J.; Faulkner, L. R. Electrochemical Methods: Fundamentals and Applications; John Wiley & Sons: New York, 1980. (76) Aydinol, M. K.; Kohan, A. F.; Ceder, G.; Cho, K.; Joannopoulos, J. Ab Initio Study of Lithium Intercalation in Metal Oxides and Metal Dichalcogenides. Phys. Rev. B: Condens. Matter Mater. Phys. 1997, 56, 1354−1365. (77) Kim, Y.; Muhammad, S.; Kim, H.; Cho, Y.-H.; Kim, H.; Kim, J. M.; Yoon, W.-S. Probing the Additional Capacity and Reaction Mechanism of the RuO2 Anode in Lithium Rechargeable Batteries. ChemSusChem 2015, 8, 2378−2384. 2045


Article

The Journal of Physical Chemistry C (78) Wang, Y.; Lee, J. Y. Molten Salt Synthesis of Tin Oxide Nanorods: Morphological and Electrochemical Features. J. Phys. Chem. B 2004, 108, 17832−17837. (79) Navulla, A.; Stevens, G.; Kovalenko, I.; Meda, L. Hierarchical Columnar RuO2 Nanoplates and Their Improved Cycle Life Performance at High Capacity. J. Phys. Chem. C 2014, 118, 13382− 13386. (80) Meduri, P.; Pendyala, C.; Kumar, V.; Sumanasekera, G. U.; Sunkara, M. K. Hybrid Tin Oxide Nanowires as Stable and High Capacity Anodes for Li-Ion Batteries. Nano Lett. 2009, 9, 612−616. (81) Winter, M.; Brodd, R. J. What Are Batteries, Fuel Cells, and Supercapacitors? Chem. Rev. 2004, 104, 4245−4270. (82) Weppner, W.; Huggins, R. A. Determination of the Kinetic Parameters of Mixed-Conducting Electrodes and Application to the System Li3Sb. J. Electrochem. Soc. 1977, 124, 1569−1578. (83) Zhu, Y.; Wang, C. Galvanostatic Intermittent Titration Technique for Phase-Transformation Electrodes. J. Phys. Chem. C 2010, 114, 2830−2841. (84) Johannes, M.; Stux, A. M.; Swider-Lyons, K. E. What Makes a Ru-Based Compound a Viable Lithium-Ion Insertion Material? ECS Trans. 2007, 11, 7−11. (85) Bader, R. F. Atoms in Molecules; Oxford University Press: Oxford, U.K., 1990. (86) Henkelman, G.; Arnaldsson, A.; Jónsson, H. A Fast and Robust Algorithm for Bader Decomposition of Charge Density. Comput. Mater. Sci. 2006, 36, 354−360. (87) Jamnik, J.; Maier, J. Nanocrystallinity Effects in Lithium Battery Materials Aspects of Nano-Ionics.Part IV. Phys. Chem. Chem. Phys. 2003, 5, 5215−5220. (88) Liu, X. H.; Zheng, H.; Zhong, L.; Huang, S.; Karki, K.; Zhang, L. Q.; Liu, Y.; Kushima, A.; Liang, W. T.; Wang, J. W.; et al. Anisotropic Swelling and Fracture of Silicon Nanowires During Lithiation. Nano Lett. 2011, 11, 3312−3318. (89) Liu, Z.; Deng, H.; Mukherjee, P. P. Evaluating Pristine and Modified SnS2 as a Lithium-Ion Battery Anode: A First-Principles Study. ACS Appl. Mater. Interfaces 2015, 7, 4000−4009. (90) Batsanov, S. Van Der Waals Radii of Elements. Inorg. Mater. 2001, 37, 871−885. (91) Peled, E. Lithium Stability and Film Formation in Organic and Inorganic Electrolyte for Lithium Battery Systems. In Lithium Batteries; Gabano, J. P., Ed.; Academic Press: Waltham, MA, 1983. (92) Lithium-Ion Batteries: Solid-Electrolyte Interphase; Balbuena, P. B., Wang, Y. X., Eds.; Imperial College Press: London, 2004. (93) Xu, K. Nonaqueous Liquid Electrolytes for Lithium-Based Rechargeable Batteries. Chem. Rev. 2004, 104, 4303−4418. (94) Aurbach, D. Review of Selected Electrode−Solution Interactions Which Determine the Performance of Li and Li Ion Batteries. J. Power Sources 2000, 89, 206−218. (95) He, Y.-B.; Liu, M.; Huang, Z.-D.; Zhang, B.; Yu, Y.; Li, B.; Kang, F.; Kim, J.-K. Effect of Solid Electrolyte Interface (SEI) Film on Cyclic Performance of Li4Ti5O12 Anodes for Li Ion Batteries. J. Power Sources 2013, 239, 269−276. (96) Zhukovskii, Y. F.; Balaya, P.; Kotomin, E. A.; Maier, J. Evidence for Interfacial-Storage Anomaly in Nanocomposites for Lithium Batteries from First-Principles Simulations. Phys. Rev. Lett. 2006, 96, 058302.

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Comparison of Storage Mechanisms in RuO2, SnO2, and SnS2 for

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