Combined Computational and Experimental Study of Li Exchange

Article pubs.acs.org/JPCC

Combined Computational and Experimental Study of Li Exchange Reaction at the Surface of Spinel LiMn2O4 as a Rechargeable Li-Ion Battery Cathode Masanobu Nakayama,*,†,§,∥ Hiroyuki Taki,⊥ Tomoaki Nakamura,† Satoshi Tokuda,† Randy Jalem,†,∥ and Toshihiro Kasuga‡ †

Department of Materials Science and Engineering and ‡Department of Frontier Materials, Nagoya Institute of Technology, Gokiso, Showa, Nagoya, Aichi 466-8555, Japan § Japan Science and Technology Agency, PRESTO, 4-1-8 Honcho Kawaguchi, Saitama 332-0012, Japan ∥ Unit of Elements Strategy Initiative for Catalysts & Batteries (ESICB), Kyoto University, Katsura, Saikyo-ku, Kyoto 615-8520, Japan ⊥ Department of Applied Chemistry, Tokyo Institute of Technology, Ookayama, Meguro, Tokyo 152-8555, Japan S Supporting Information *

ABSTRACT: Lithium manganese oxide (LiMn2O4) is regarded as an attractive positive electrode for rechargeable Li-ion batteries, in particular for the power source of electric vehicles, and there is thus an urgent need to improve its charge−discharge kinetics. In the current paper, the kinetics of a Li-ion exchange reaction at the interface between a LiMn2O4 cathode and nonaqueous liquid electrolytes is studied using experimental and computational techniques. Electrochemical ac impedance measurements showed two semicircles corresponding to the interfacial Li-ion exchange, and they are ascribed to the desolvation and lattice incorporation processes according to the adatom model [Bruce, P. G.; Saidi, M. Y. J. Electroanal. Chem. 1992, 322, 93]. To gain deeper insight into the latter process, delithiation from the electrode surface was simulated using density functional theory (DFT), and the DFT results were compared to the dependence of the ac impedance behavior on potential. We conclude that the chemical potential gradient is formed at the surface of positive electrodes, and the difference of the potential between the surface and the bulk corresponds to the activation energy of lattice incorporation.

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INTRODUCTION

exchange reactions) is still limited at the atomic and electronic levels. In general, the kinetics of electron transfer in soluble redox systems, such as the ferrocene/ferrocenium redox couple, have been well established, according to the theory underlying the Butler−Volmer-type charge-transfer reaction.6 In these reactions, electron transfer is assumed to occur at the electric double layer formed at the part of the liquid electrolyte adjacent to the surface of the electron-conductive solid-state electrode. The electron transfer reaction, however, takes place inside the solid-state electrode materials of LIB (oxidizing Co3+ to Co4+ for the LixCoO2 system), so that the activation energy for electron transfer would be smaller than in the case of the soluble system. On the other hand, Li+ ion exchange occurs at the heterophase interface of the solid electrode and liquid electrolyte and is suggested to be the rate-determining process for the interfacial exchange reaction.

Electrochemical lithium intercalation is a reaction with particular significance to high-energy-density electrode materials used in Li-ion batteries (LIBs).1−5 Thus, materials that possess lithium intercalation sites have attracted considerable attention owing to both the range of practical battery applications and the fundamental interest in the electrochemical electron/ion exchange reaction in the bulk of the electrode and at the interface of the electrode and electrolyte. This reaction involves a relatively simple mechanism: (1) the host crystal structure of the electrode remains almost unchanged from before to after the electrochemical reaction, and (2) the free energy of the reaction, the molar amount of reacted species, and kinetic parameters can be obtained easily by measuring cell voltage and electric current. There have been numerous reports on the bulk properties of the crystalline phases of LIB electrodes, such as their phase stability, crystal and electronic structures, and ion-diffusion properties. However, our understanding of the acquired knowledge on kinetic properties at the electrode/electrolyte interface, (i.e., the kinetics of fundamental electron/ion © 2014 American Chemical Society

Received: September 12, 2014 Revised: November 3, 2014 Published: November 5, 2014 27245

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size can be accounted for by such an increase in the relative volume of the surface region. The aim of the present study is to clarify the atomic-level description of the lattice incorporation mechanism for the surface Li+ exchange reaction by assuming the chemical potential slopes at the surface region.

Bruce and Saidi suggested a two-step model for such ion exchange across an interface.7 First, in the adatom process, solvated ions in the electrolyte near the electrode surface are subjected to partial desolvation, resulting in formation of an adatom on the electrode surface. Then, during the incorporation process (also called “lattice incorporation”), the remaining bound solvent molecules are released from the adatom, and the surface ion (Li+) becomes fully incorporated into the electrode material. We previously confirmed this twostep mechanism by using electrochemical impedance spectroscopy (EIS) for Li+ insertion into perovskite-type La1/3NbO3.8 Abe and colleagues have carried out particularly in depth experimental and computational studies on the adatom process for various materials, showing that solvation/desolvation of Li+ is one of the key factors affecting the overall reaction kinetics in LIBs (e.g., refs 9−13). A clear atomic-level understanding of the mechanism for incorporating the Li+ into the electrode lattice, however, remains elusive, which prevents rational design of the surface of electrode materials. For the current study, we have attempted to clarify the factors affecting the lattice incorporation process at the surface of the solid electrode. We have based our work on the report by Wang et al.,14,15 in which they calculated the Li+ removal energies, that is, the chemical potential of Li, in the bulk and surface Li sites of olivine-type cathode materials by using firstprinciples density functional theory (DFT) and a slab model. The chemical potential difference between the bulk and surface sites, here called the “chemical potential slope”, may form the energy barrier and hence control the Li+ exchange reaction at the interface between the electrode and electrolyte. Therefore, we infer that the chemical potential difference influences the lattice incorporation process of the Li+ exchange reaction. To verify this idea, both computational and experimental approaches are combined in this paper. The interfacial Li+ exchange reaction was investigated using spinel-type LiMn2O4, which belongs to one of the three most extensively studied families of cathode materials for LIBs. This oxide was chosen due to its high safety, low cost, and abundant availability.16−20 In addition, particle-size-dependent voltage profiles of nanosized LixMn2O4 electrodes have been reported.21,22 According to these reports, typical LiMn2O4 with a particle size of more than 1 μm shows a redox potential ranging from 3.9 to 4.2 V, while nanosized particles show capacities below and above the 3.9−4.2 V range at the early and late stages of the charge/discharge reaction, with the corresponding capacity increasing monotonically as the size of the nanoparticles is reduced. A similar trend was also reported for LixCoO2 electrodes by Ohkubo et al.23 According to the laws of thermodynamics, the redox potential does not depend on particle size, but on the structure and composition of the bulk material. Therefore, the reported particle size-dependent phenomena may relate to surface structure or properties. For example, surface redox potential differs from the bulk potential due to the amorphization or reconstruction of the surface crystal structure. If one assumes that formation of a surface region involves structural disordering that extends to depths of a fixed number of nanometers (say, 1 nm) below the surface regardless of the particle size, then the molar fraction of total LiMn2O4 located in the surface region would increase as the particle size is reduced and would increase especially sharply for nanoparticles. (see Figure S1 in Supporting Information and Figure 4 in ref 21). Therefore, the increase of the capacity at voltages outside of the 3.9−4.3 V range by reducing particle

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METHODS Experiment. A powder sample of LiMn2O4 was prepared by a conventional solid-state reaction based on previously reported methods.24−26 A mixture of stoichiometric amounts of Li2CO3 and Mn2O3 (which was prepared by preheating of MnCO3 at 600 °C for 36 h) was used as the starting material. The mixture was heated at 750 °C for 12 h in air and then cooled at a rate of 0.5 °C/min. Phase identification was carried out by powder Xray diffraction (XRD) analysis using Cu Kα radiation (data not shown). Electrochemical measurements for LixMn2O4 (x ≤ 1) were made using a three-electrode cell. Li foil (Aldrich) was used as both counter and reference electrodes. Except when otherwise indicated, electrolyte solutions consisted of 0.2 M LiClO4 in anhydrous ethylene carbonate (EC) and diethylene carbonate (DEC), and these chemicals were purchased from Tomiyama Pure Chemical Co., Ltd. The working electrode consisted of a mixture of 90 wt % perovskite powders, 7 wt % acetylene black as a current collector, and 3 wt % poly(tetrafluoroethylene) (PTFE) binder. Before potentiostatic ac impedance measurements, galvanostatic charge−discharge tests for two cycles were conducted over a potential range between 3.5 and 4.5 V at a C/10 rate. The ac impedance measurements for the LixMn2O4 electrode were performed at various voltages, after potentiostatic relaxation down to a C/200 rate current. The impedance spectra were measured at frequencies ranging from 10−2 to 105 Hz at 20 mV in amplitude with the aid of a VMP-3 electrochemical interface (Bio-Logic). Cells were prepared in an Ar-filled glovebox. The electrochemical measurements were carried out in an airtight vessel filled with Ar gas, which allowed the cell to be removed from the glovebox. The vessel was placed in an incubator to control the experimental temperature. Computations. The computations were performed using the spin-polarized generalized gradient approximation (GGAPBEsol)27,28 and GGA + U methods,29,30 using the plane-wave basis set and the projector-augmented wave (PAW) method31,32 as implemented in the Vienna ab initio simulation package (VASP).33,34 The effective U parameter for Mn 3d electrons was set to 5.0 eV in reference to previous reports.35,36 At first, we calculated the structures of bulk LiMn2O4 and Mn2O4 under the condition of full structural relaxation (i.e., allowing for the change of the cubic lattice parameter and the internal atomic positions). Then, the final energies of the optimized geometries were recalculated so as to correct for changes in the plane-wave basis during relaxation. For the purpose of comparison, all the calculations were performed in a cubic unit cell with Fd3̅m symmetry that contained 8 formula units of LiMn2O4 and of Mn2O4 using 5 × 5 × 5 k-point meshes. After determining the lattice parameters of pure LiMn2O4 and Mn2O4, the surface structures were simulated using the slab technique, in which a set of infinite layers separated by vacuum layers are repeated periodically along the surface normal.37 The slabs were constructed such that their two sides are symmetrically equivalent and can be mapped into each other 27246

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by an inversion or mirror-type symmetry operation in the middle of the slab. Low-index facets of the (001), (011), and (111) surfaces were calculated in the current work as done previously.38,39 In this study, charge-balanced stoichiometric slabs were treated by adjusting only the number of surface atoms while satisfying the above-mentioned symmetry constraint. A total of 16 termination models were constructed for LiMn2O4 and Mn2O4 as follows: Li- and MnO-terminated {001} surfaces, LiMnO- and MnO-terminated {011} surfaces, and Mn-, O-, LiMn-, and inverse spinel-type Li-terminated {111} surfaces (see Figures S2−S4). Note that undercoordinated Mn ions are swapped with tetrahedral Li ions in the inverse spinel-type {111} surface model.39 All the bulk and slab calculations were performed taking the symmetry into account during ionic relaxation, which prevents the transition to a lower symmetry phase for the bulk and cancels the polarity.

Table 1. Calculated Surface Energies of Various Facets and Terminations for LiMn2O4 and Mn2O4 Compoundsa compounds

facet

termination

surface energy (meV Å−2)

LiMn2O4

100

Li MnO LiMnO MnO Mn O LiMn Li (inverse-spinel) Li MnO LiMnO MnO Mn O LiMn Li (inverse-spinel)

35 84 72 73 85 105 44 21 74 106 109 102 87 126 85 67

110 111

Mn2O4

110

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111

RESULTS Computational Results. The calculated cubic lattice parameters for bulk LiMn2O4 and Mn2O4 are 8.324 and 8.170 Å, respectively, and averaged intercalation voltage between the two compositions (LiMn2O4 → Li + Mn2O4) is 3.87 V, showing reasonable agreement with previous computational36 and experimental studies (experimental cubic lattice parameters are equal to 8.2455 and 8.029 Å for LiMn2O4 and Mn2O4, respectively, and the voltage is 4.02 V26,40,41). The obtained lattice parameters for the ab-plane of the slab (parallel to each slab model) were kept fixed, and the coordinates of only the ions were allowed to relax. The thickness of the vacuum layer was set at ∼10 Å in the present slab simulation. The surface energies, Esurf, can be determined for the stoichiometric slab simulation using the equation

a

Detailed atomic arrangements for each model are shown in Figures S2−S4 (Supporting Information).

Figure 1. Computationally derived optimized morphologies of (a) LiMn2O4 and (b) Mn2O4 particles. In each case, {111} facets constitute the majority of the crystal surface and some {100} facets are present, based on the Wulff construction.

Esurf = 1/(2σ )[Eslab(Li nxMn2nO4n) − nE bulk (LixMn2O4 )] (x = 0, 1)

100

(1)

In this equation, Eslab and Ebulk refer respectively to the total energy of the slab and the bulk model, which were used for cubic LiMn2O4 and Mn2O4. Also, n and σ respectively stand for the mole number and the surface area of LiMn2O4 and Mn2O4 in the corresponding slab model. Table 1 lists the calculated surface energies for each facet with several termination models. Among the eight slab models tested, the (111) surface with a local inverse-spinel arrangement exhibits the lowest surface energy for both LiMn2O4 and Mn2O4, as reported by Karim et al.39 The optimized particle morphology is shown in Figure 1a,b and was estimated by applying the Wulff construction, in which the distance to the central point {hkl} from the origin is taken to be proportional to the surface energy.42 The facets of the (110) surface were not apparent due to its relatively large surface energy. The facets of the (111) surface, which is the most stable surface, cover the largest area of the optimized morphology in both compositions; these facets cover nearly the entire surface of LiMn2O4, as reported by Karim et al.39 On the other hand, our calculations indicate that the facets of the (100) surface cover 27% of the surface of the Mn2O4 composition. Therefore, we infer that the (111) and (100) facets form the dominant surfaces of the LiMn2O4 and Mn2O4 crystals. We investigated the redox potentials for electrochemical removal (uptake) of Li from (into) bulk and various surfaces of LiMn2O4 with the most stable termination. Note that two Li ions were removed from both the (111) and (100) surfaces to cancel the polarity for each slab. The bulk redox potential for

removal/uptake of Li was calculated according to a reported reaction, LiMn2O4 → Li0.5Mn2O4 + 0.5Li or Mn2O4 + 0.5Li → Li0.5Mn2O4, in which the lowest voltage was obtained within the GGA + U framework.43 The present calculation results are shown in Table 2. These redox potentials at both the (100) and (111) surfaces for LiMn2O4 are much smaller than that of the bulk, but are comparable to redox potential calculated for the Table 2. Calculated Redox Potentials for the Removal of Li from Various Facets of LiMn2O4 as Well as the Bulk Potential compounds

facets

voltage (V)

LiMn2O4

{100} {110}a {111} bulkb {100} {110} {111} bulkb

3.10 3.75 3.33 3.84 3.87 4.27 4.70 3.88

Mn2O4

a

Surface does not appear in the crystal derived from the Wulff construction shown in Figure 1. bVoltage was calculated by considering the reaction where a single Li ion was removed from (inserted into) a bulk-Li8Mn16O32 cell (bulk Mn16O32). 27247

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(110) surface. This difference in redox potentials between the surfaces and the bulk supports the report by Wang et al.14,15 Smaller voltages at the (100) and (111) surfaces than in the bulk would cause Li ions to be first removed from the LiMn2O4 particles at these surfaces, followed by removal of Li ions in the bulk in the charging process. In addition, our computational results also support previous experimental results that indicate that the capacity for charging at the relatively low voltage region (i.e., at 3.5−3.9 V) increases with decreasing particle size. Note that redox capacity is mainly observed from 3.9 to 4.2 V for spinel-type LixMn2O4. Taken together, these calculations suggest that there are more surface Li sites at lower redox potentials. Experimental Results. As shown in Figure 2, discharge occurs at a slow rate (0.05 C) following the first charging. No

Figure 3. Typical ac impedance plots at 4.0 V vs Li+/Li in the cell Li/ 0.2 M LiClO4 in GBL/LixMn2O4. Panel b shows a magnification of panel a.

Figure 2. First discharge profiles of the Li/LiMn2O4 cell using the nonaqueous electrolytes (a) 0.2 M LiClO4 in EC:DEC (1:1 volume ratio) and (b) 0.2 M LiClO4 in GBL.

significant irreversibility was observed after the second cycle (not shown), indicating negligible effect arising from side reactions, except for an irreversible capacity loss of about 10 mA h g−1 at the first charge. Thus, we performed electrochemical ac impedance spectroscopy at various potentials; all measurements were conducted after the slow first charge/ discharge cycle unless mentioned otherwise. Typical impedance plots are shown in Figure 3 for the cell at 4.0 V vs Li+/Li. In the low-temperature region, two distinct semicircles were seen at the frequency range from 10 to 1 kHz (semicircle A) and from 0.1 to 10 Hz (semicircle B), shown in Figures 3a and 3b, respectively. The details of the small semicircle are hidden at the high-frequency region around 1 kHz, owing to the solid electrolyte interface (SEI) resistance;44 however, this issue is beyond the scope of the present paper. According to the adatom model,7,8 the semicircles at high- and low-frequency ranges could be attributed respectively to desolvation and lattice incorporation. In the former process, solvated Li ions in the electrolyte are (partially) desolvated and adsorbed on the electrode surface (adatom), whereas the latter process corresponds to the incorporation of the adatom into the electrode lattice. To confirm the validity of the present assignments for the two semicircles, activation energies were estimated by measuring the temperature dependence of impedance spectra at various electrochemical potentials. Inverse resistances, which are proportional to exchange current density, from both semicircles A and B are plotted as a function of reciprocal temperature (i.e., an Arrhenius plot). Typical results are displayed in Figure 4a, and a linear relationship is indeed observed, which follows the Arrhenius equation and allows activation energies to be straightforwardly derived. The

Figure 4. (a) Typical Arrhenius plots of interfacial resistances from semicircles A and B shown in Figure 3. The results are obtained at the potential of 3.95 V vs Li+/Li using the electrolyte 0.2 M LiClO4 in a mixture of EC and DEC. Panels b and c show the dependence of activation energy on potential for semicircles A and B, respectively.

activation energies for semicircles A and B are summarized in Figures 4b and 4c, respectively, in terms of potential and solvent selection. Note that we measured ac impedance spectra using two kinds of solvents, EC + DEC and γ-butyl lactone (GBL), due to the relatively large difference in their dielectric constants (90 and 40 for EC and GBL, respectively) despite their similar molecular structures. The desolvation energies of Li+ from EC and GBL were evaluated with the DFT along with the literature,45 and the calculated results indicate an approximately 0.1 eV higher desolvation energy in GBL than in EC (see Table S1 in the Supporting Information). Variation 27248

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Li+ vacancy formation energies between the outermost surface layer sites and the bulk sites. In Figure 5, the chemical potential at the surface is assumed to be higher (and the voltage is assumed to be lower) than in the bulk, which is consistent with our DFT calculations (Table 2). Nevertheless, the opposite situation is also conceivable, where the chemical potential of Li at the surface is lower than that at the bulk (see Figure S5 in the Supporting Information). In addition, we assumed that the chemical potential anomaly occurs not only at the outermost surface but continues through a surface layer on the order of a nanometer thick in this model. This is supported by the fact that the experimental voltage profile depended on the particle size, as mentioned in the Introduction, due to the rapid increase of the molar fraction of Li in the surface region that occurs upon reducing the size of the particle to a nanometer scale (see Figure S1 in the Supporting Information). Similarly, Moriwake et al. have recently reported that a voltage anomaly forms at the vicinity of the grain boundary of LiCoO2 based on the results of first-principles DFT.46 When charging starts, Li ions are removed from those surface regions such as the {100} surface in the present computation whose sites have higher chemical potentials (lower voltages) than do those in the bulk. Then, Li ions need to climb the energy barrier from the corresponding chemical potential to the top of chemical potential. As displayed in Figure 4c, the activation energy for the semicircle B is relatively low (∼20 kJ mol−1) at lower voltages ranging from 3.8 to 4.0 V, in which Li ions at the surface region are extracted. Note that corresponding capacity in this voltage region is rather small due to limited volumetric fraction of surface region in the particle. On the other hand, the observed activation energy for the semicircle B become higher (∼40 kJ mol−1) above 4.0 V, where Li ions in the bulk react, in comparison with the low voltage region. In addition, the activation energies are almost the same regardless of the selection of electrolyte solvent as mentioned above. Therefore, we conclude that the activation energies derived from semicircle B correspond to the chemical potential gradient formed at the electrode surface as shown in Figure 5. Note that we discussed only the onset of the charging process at the lower voltage region (