Surface Chemistry Consequences of Mg-Based Coatings on LiNi0

Article pubs.acs.org/JPCC

Surface Chemistry Consequences of Mg-Based Coatings on LiNi0.5Mn1.5O4 Electrode Materials upon Operation at High Voltage Gabriela Alva,†,‡ Chunjoong Kim,†,§ Tanghong Yi,† John B. Cook,†,⊥ Linping Xu,†,¶ Gene M. Nolis,† and Jordi Cabana*,†,§ †

Environmental and Energy Technology Division, Lawrence Berkeley National Laboratory, One Cyclotron Road, Berkeley, California 94720, United States ‡ College of Chemistry, University of California, Berkeley, California 94720, United States § Department of Chemistry, University of Illinois Chicago, Chicago, Illinois 60607, United States S Supporting Information *

ABSTRACT: LiNi0.5Mn1.5O4 epitomizes the challenges imposed by high electrochemical potential reactivity on the durability of high energy density Li-ion batteries. Postsynthesis coatings have been explored as a solution to these challenges, but the fundamentals of their function have not been ascertained. To contribute to this understanding, the surface of LiNi0.5Mn1.5O4 microparticles was modified with Mg2+, a coating component of literature relevance, using two different heat treatment temperatures, 500 and 800 °C. A combination of characterization tools revealed that Mg2+ was introduced mainly as an inhomogeneous MgO coating in the sample treated at 500 °C, and into the spinel lattice at the subsurface of the particles at 800 °C. Comparing the properties of these two different materials with an unmodified baseline afforded the opportunity to evaluate the effect of varying surface chemistries. Coulometry in Li metal half cells was used as a macroscopic measure of side reactions at the electrode−electrolyte interfaces. This magnitude was comparable in all the materials at room temperature. In contrast, a significant drop in efficiency was observed in the untreated material when the cycling temperature was raised to 50 °C, but not in the modified materials. The origin of the reduced reactivity of the materials after introducing Mg-based modifications was evaluated by probing the chemical changes at the Ni−O bonds using soft XAS. Taken together, the results of this study revealed that incorporation of Mg stabilizes highly oxidized Ni−O species, which can be related to the better stability toward the electrolyte. They point to a pathway toward the guided design of efficient surface modifications to yield battery electrode materials with increased stability against the electrolyte.

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INTRODUCTION In recent years, concerns have arisen regarding the extensive use of fossil fuels. Some of these concerns include dependence on foreign countries for oil, its limited availability, the volatility of gasoline prices, and the possible pernicious effect of excessive greenhouse and polluting gas emissions into the atmosphere.1 As such, significant research efforts have been directed at developing energy storage technologies that can penetrate into the transportation market and be an enabling piece of a smart grid based on renewable energies, thus displacing fossil fuels from their current central role in energy policy.2 Due to their high energy density compared to other technologically mature options, batteries based on the Li-ion concept, where Li+ ions shuttle between solid intercalation electrodes through a liquid electrolyte, are at the forefront of these efforts. However, their metrics still fall short of the requirements for economic viability.3 As a result, vigorous activities are directed at developing new materials that boost their energy storage capability.4 LiNi0.5Mn1.5O4 is viewed as an attractive active material for the positive electrode in Li-ion batteries.5 It has a spinel structure, which is stable to Li extraction and enables fast ion © 2014 American Chemical Society

diffusion, and reacts at an electrochemical potential (4.7 V vs Li+/Li0)6 that is 0.7−1.0 V higher than other commercial alternatives, such as LiCoO27 and LiFePO4.8 The electrochemical deintercalation of lithium is formally compensated by a two-electron redox reaction in Ni, which changes from Ni2+ to Ni4+.9 The high voltage of operation leads to an increase in energy density when used in a device. However, it also causes the electrolyte to anodically decompose on the surface of the electrode.10 The result of this side reaction is the formation of products that passivate the electrode, leading to severe losses of storage capacity after multiple cycles. In addition to this deleterious redox interaction, LiNi0.5Mn1.5O4 also suffers from corrosion driven by acidic impurities in electrolytes based on fluorinated salts such as LiPF6.11,12 These issues are aggravated when cycling is performed above room temperature,13,14 as can frequently occur in a vehicle. The chemical mechanisms that lie behind these undesired electrode−electrolyte interactions have not been clearly Received: January 10, 2014 Revised: April 30, 2014 Published: April 30, 2014 10596

dx.doi.org/10.1021/jp5003148 | J. Phys. Chem. C 2014, 118, 10596−10605

The Journal of Physical Chemistry C

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to be comparable to the literature, providing a relevant case study for our goals. Accurate coulometry measurements revealed the extent of irreversible side reactions associated with each material. The work focused on establishing correlations between chemistry and electrochemical performance, analyzing the effects of the coating on the material surface chemistry and morphology. The most important finding is that, based on soft XAS data collected with sensitivity to either the surface or bulk of the material, the presence of Mg2+ within the structure of the spinel surface could stabilize electron-depleted Ni−O bonds, especially upon high-temperature operation.

ascertained, but the excellent stability of the spinel structure to cycling turns LiNi0.5Mn1.5O4 into a model material to understand them. In the absence of this information, the design of strategies to alleviate this issue has focused on placing electronically insulating solid barriers between the electrode surface and the liquid electrolyte. This strategy has been applied to a number of positive electrode materials of intense technological interest,15−20 including LiNi0.5Mn1.5O4.12,21,22 These barriers are thought to serve a dual purpose.17,23 They prevent direct contact between acidic impurities and oxide and, by virtue of their insulating character, they effectively screen the electrochemical potential at the electrode, so that electrolyte oxidation is not induced. Another possibility is that these solid barriers act as a scavenger of acidic impurities instead of a physical barrier. For instance, Sclar et al. did a comparative study of LiNi0.5Mn1.5O4 coated with MgO, which is a Lewis base, and ZnO, which is amphoteric, and consistently saw better performance with the former.12 A significant proportion of studies in the literature propose employing materials that do not contain lithium and, thus, should be insulating to lithium transport as well. In principle, impeded ion transfer kinetics would be detrimental to operation, yet these same reports concluded that coating the LiNi0.5Mn1.5O4 particles with oxide layers formulated as ZnO, ZrO2, or MgO can help alleviate the cycling losses induced by negative processes at the electrode− electrolyte interface, even at high temperature.12,17,24−26 Formation of phases different from these upon chemical contact with the electrode surface during processing and/or cycling could result in coatings that contain lithium,27,28 thus providing a possible mechanism to explain this apparent contradiction. Generally speaking, understanding the chemistry of oxidized species on the electrode surface and the function of coatings would enable a rational approach to the compositional and morphological design of tailored composites, irrespective of the active material, where the inactive and insulating fractions are reduced to the minimum required for complete elimination of undesired reactivity. Some authors have also suggested that the ions in the coating incorporate into the lattice of the active material.29 However, no direct evidence is available in the literature of how coatings modify the surface chemistry of the oxide vis-à-vis the bonding nature at the redox active electronic states. In the case of LiNi0.5Mn1.5O4 and a variety of layered oxides, these states are defined by the Ni−O chemical bond. Changes in covalence as electrons are removed from these bonds can be most effectively probed by X-ray absorption spectroscopy (XAS) at the Ni LII,III- and O K-edges (i.e., using soft X-ray beams),30 which directly measure the unoccupied density of states at the transition metal 3d and O 2p bands.31 No studies could be found in the literature where this tool has been used to evaluate the role of coatings on the reactivity of electrode materials upon lithium removal. In a previous study, we established that micrometer-scale, octahedral LiNi0.5Mn1.5O4 particles offer better extended performance than nanoscale counterparts, primarily because they minimize the electrolyte−electrode interface.32 However, electrolyte decomposition still occurred and severely compromised performance when the best samples were used in a full Li-ion device against graphite. Thus, coatings would seem to be necessary for further improvement. This work focuses on Mgbased modifications of the surface of these LiNi0.5Mn1.5O4 particles because they are a popular choice by battery engineers. The electrochemical performance of the electrodes was found

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EXPERIMENTAL SECTION LiNi0.5Mn1.5O4 was synthesized by using a solid state reaction between LiOH·H 2 O (98%, EMD Millipore) and Ni 0 . 2 5 Mn 0 . 7 5 (OH) 2 mixed in stoichiometric ratio. Ni0.25Mn0.75(OH)2 was synthesized by using a coprecipitation reaction. In a typical synthesis, a mixed metal solution was produced by dissolving nickel sulfate hexahydrate (NiSO4· 6H 2 O, 4.7523 g, 18.1 mmol) and manganese sulfate monohydrate (MnSO4·H2O, 9.1679 g, 54.2 mmol) in 50 mL of deionized water. The metal sulfate solution was added dropwise into a 400 mL continuously stirred (400 rpm) aqueous solution of sodium hydroxide (NaOH, 11.55 g, 0.289 mol) at a rate of 1 mL/min, using a Heidolph PD5201 peristaltic pump. Afterward, the dispersion was centrifuged to recover the powder, which was subsequently washed 3 times with water. In between each wash, the powder was sonicated for 5 min to redisperse the precipitate. Finally, the powder was left in an oven to dry at a temperature of 150 °C for 12 h. The lithium and metal hydroxide mixture was calcined at 900 °C for an hour in air with a heating and cooling rate of 5 deg/min to obtain the final product. At this temperature, the material shows Ni and Mn cationic disordering, which leads to the best performance.33 This baseline powder will be referred to as BL hereafter. To introduce Mg on the surface of LiNi0.5Mn1.5O4, a 1:4 (mol:mol) mixture of Mg(NO 3 ) 2 ·6H 2 O and LiNi0.5Mn1.5O4, which theoretically corresponds to a 5 wt % MgO coating, was added to 20 mL of oleylamine (70%, Aldrich) and sonicated for 30 min. The mixture was then brought to 140 °C and stirred at 900 rpm for an hour, using an oil bath. Afterward, the powder was recovered by centrifugation, washed with a mixture of ethanol and hexane, and dried overnight at 70 °C. Finally, the powder was annealed in a furnace for 12 h at the desired temperature, 500 or 800 °C, with a ramp rate of 5 deg/min. The lower annealing temperature was selected to induce decomposition of the Mg-based precursor, based on thermal analysis (Figure S1, Supporting Information). The temperature of 800 °C was also chosen to avoid Ni/Mn ordering occurring at ∼700 °C.33 Such ordering can negatively affect the electrochemical properties of LiNi0.5Mn1.5O4 with respect to the highly disordered baseline,33 thereby introducing changes that would be unrelated to the presence of surface modifications. These samples will be identified as Mg500 and Mg800, respectively, in the paper. The composition and crystallinity of the samples prepared were evaluated by X-ray diffraction (XRD). Patterns were collected between 10° and 80°, 2θ, at a rate of 0.02 deg/min, using a Bruker D2 Phaser diffractometer operating at Cu Kα radiation (λ = 1.5418 Å). The particle size and morphology resulting from the synthesis were analyzed by using scanning electron microscopy (SEM) in a JEOL 7500F operated at 1 kV and 20 mA in gentle beam mode. Soft X-ray absorption 10597



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spectroscopy (XAS) measurements were carried out at beamline 10−1 at the Stanford Synchrotron Radiation Lightsource (SSRL, Menlo Park, CA). Data were obtained at a spectral resolution of ∼0.1 eV both in total electron yield (TEY) and fluorescence yield (FY) modes to compare differences in bulk and surface of the material. The angle of detection in FY mode was set to 70° so as to minimize selfabsorption while still being bulk sensitive. The beam footprint on the sample was less than 1 mm2. X-ray photoelectron spectroscopy (XPS) was conducted at 15 kV, using a Sigma Probe with monochromatic X-ray source (Al Kα). Each sample was characterized before and after surface etching by ion sputtering for 5 min. Composite electrodes were prepared by mixing LiNi0.5Mn1.5O4 and carbon black in a 6 wt % polyvinylidene difluoride (PVDF) solution in 1-methyl-2-pyrrolinone (NMP), to reach an active material:binder:carbon weight ratio of 80:10:10. The slurry was mixed until homogeneity was reached, and was subsequently cast onto Al foil with a doctor blade. The electrode was left to dry under an infrared lamp for 30−60 min before being placed in a vacuum oven at 110 °C overnight. The material loading of the prepared electrodes was approximately 7 mg/cm2. Electrochemical cycling was carried out in two-electrode coin-type cells. Lithium metal was used as both the counterand pseudoreference electrode and a solution of 1 M LiPF6 dissolved in a 1:1 volume mixture of ethylene carbonate:diethyl carbonate (Novolyte Technologies) was used as the electrolyte. Cells were fabricated in an argon-filled glovebox. Electrochemical measurements were conducted with a VMP3 potentiostat at room temperature and 50 °C. Galvanostatic cycling was conducted at a current rate of C/10, where C was defined as 147 mAh/g (theoretical capacity for full delithiation), and cutoff voltage window of 3.5−5 V. Three to five replicas were tested at each condition to collect statistically meaningful results. Error bars were subsequently added to the corresponding data sets. Rate capability experiments were carried out by cycling the samples at an increasing rate of C/10 to 5C, at room temperature and 50 °C. The specific discharge capacities were calculated as the measured fraction of LiNi0.5Mn1.5O4 in the electrode (∼80% for BL, and ∼76% for Mg500 and Mg800). Electrodes charged to 5 V at C/10 rate for XAS measurements were carefully harvested in the glovebox, rinsed with dimethylcarbonate, and transferred into the vacuum chamber for measurement.

Figure 1. XRD patterns of (a) BL, (b) Mg500, and (c) Mg800. A zoom-in of the region around the (004) reflection is provided to highlight the peaks corresponding to rock salt-type secondary phases (indicated by an arrow).

SEM images were taken to observe the surface of the coated and uncoated samples (Figure 2). Particle size distributions were relatively wide, formed by octahedra ranging from approximately 0.5 to 3 μm for all three samples. The surfaces of BL particles were smooth and clean. In contrast, Mg500 showed a thin, rough coating layer over the surface of the particles with a few unevenly distributed agglomerates throughout. Energy-dispersive X-ray spectroscopic imaging was coupled with SEM (SEM-EDS) to investigate the presence of Mg. Elemental maps of Mg, Mn, and Ni collected for this sample material can be found in Figure S2 (Supporting Information), parts b−d, respectively. They show that Mg was present throughout the material. XPS also confirmed the presence of Mg, Ni, and Mn (Figure S3, Supporting Information); the spectra were consistent with Mg2+ 35 and Ni2+,36,37 whereas the exact oxidation state of the Mn was difficult to accurately ascertain at this resolution.38,39 EDS line scans (Figure 3) showed measurable contents of Mg throughout the surface, with spikes associated with areas that contain agglomerates. Mg800 showed no visible sign of such coating or agglomerates, the appearance of the crystals thus resembling pristine LiNi0.5Mn1.5O4. Nonetheless, SEM-EDS maps (Figure S4, Supporting Information) and XPS (Figure S3, Supporting Information) revealed the presence of Mg on the sample surface. This phenomenon can be explained by the reaction between the small MgO domains in Mg500 and the spinel surface, inducing incorporation of Mg2+ into the subsurface structure. While the Mg 1s, Ni 2p, and Mn 2p XPS were essentially the same for Mg500 and Mg800, sputtering the surface revealed an upward shift in binding energy of Mn in the latter, which could be due to differences in formal valence states between surface and bulk,39 supportive of the existence of a reaction between coating and spinel surface upon high-temperature annealing. Electrochemical Properties. Parts a and b of Figure 4 show plots for specific discharge capacity (mAh/g) and Coulombic efficiency (%) vs cycle number, respectively, at room temperature for the three samples, BL, Mg500, and Mg800. The corresponding voltage vs specific capacity profiles during the first cycle at room temperature can be found in Figure S5, Supporting Information. No significant changes in the electrochemical signals was found after post-treating with Mg2+ species; these signals were found to be typical of

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RESULTS Crystal-Chemical and Morphological Characterization. The XRD patterns obtained from all three samples are given in Figure 1. The major peaks in all patterns could be indexed with a simple spinel structure (space group Fd3̅m). The spinel peaks positions were fit, resulting in lattice parameters for BL, Mg500, and Mg800 materials that were measured to be 8.171(3), 8.170(4), and 8.168(7) Å, respectively, within error of each other. Very subtle additional reflections at 43.8°, 2θ, corresponding to a rock salt-type oxide containing Ni and Mn33 were found in both BL and Mg500. The intensity of this peak noticeably increased in Mg800, its line shape becoming broader and more complex. This change is an indication of the reaction between the coating and the spinel surface. However, little change would be expected in the cell parameters if Mg2+ incorporated into the lattice in the place of Ni2+, as both ions have similar ionic radii.34 10598



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Figure 3. (a) SEM images of a single Mg500 crystal, where the lines with an arrow indicate the EDS line scan location and direction; (b) EDS line graphs, where the black line corresponds to oxygen, red is magnesium, green is manganese, and blue is nickel.

Figure 2. Scanning electron micrographs of (a) BL, (b) Mg500, and (c) Mg800.

LiNi0.5Mn1.5O4 showing cationic disorder in the octahedral sites.40 The reactions between electrode and electrolyte are especially prominent in the initial cycles, followed by stabilization through passivation of the electrode surface.32,41 In this case, we verified in selected cases that they tended to stabilize after just 10 cycles, which is the number shown in the figures. The Coulombic efficiency was used as a direct measure of the extent of irreversible electrolyte decomposition, whereas the cycle retention was an indication of both impedance growth and active material loss due to dissolution. BL displayed the highest specific discharge capacity (∼128 mAh/g), followed by Mg800 (∼118 mAh/g), and then Mg500 (∼112 mAh/g). The presence of Mg-rich surface modifications was thus detrimental to this figure of merit. Interestingly, despite the Mg contents being the same in both modified samples, the capacity decrease was not, suggesting that the modifications had different nature and role on the properties. The Coulombic efficiency was

Figure 4. (a) Specific discharge capacity and (b) Coulombic efficiency vs cycle number of a series of Li metal cells (3−5 replicas) containing the different LiNi0.5Mn1.5O4 prepared in this study as working electrodes, cycled at room temperature at C/10.

almost identical for BL and Mg800, at about 92−93%, while Mg500 displayed lower efficiency averages, at 85%. 10599



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Since the severity of parasitic reactions increases with temperature,42 cycling was also performed at 50 °C. Specific discharge capacity and Coulombic efficiency vs cycle number graphs at 50 °C for samples BL, Mg500, and Mg800 are shown in Figure 5, parts a and b, respectively. A statistically significant

Figure 6. Normalized discharge capacity at multiple rates of Li metal cells containing the different LiNi0.5Mn1.5O4 prepared in this study as working electrodes, cycled at (a) room temperature and (b) 50 °C.

temperature, BL exhibited the best rate capability of the three samples with a retention of 70% at 5C with respect to C/10, followed by Mg800 then Mg500, with retentions of 63% and 52%, respectively. However, when testing was carried out at 50 °C, Mg800 displayed the best rate capability when going from C/10 to 5C, with retention of 62%, followed by BL then Mg500, with 48% and 42%, respectively. Thus, going from room temperature to 50 °C, BL lost about 22% rate capability by the rate of 5C. In contrast, Mg800 was hardly affected at all by temperature, consistently giving the same approximate capability for each rate under both conditions. Mg500 represents an intermediate case; its rate capability degraded at 2C and 5C with respect to room temperature, but significantly less than BL. In general, the electrochemical results are in line with the literature, whereby coatings are generally found a beneficial effect on performance.16,22 The system selected was thus a relevant choice for an in-depth study of the chemical consequences of such coatings. Soft X-ray Absorption Study. The measurement of Coulombic efficiencies, especially at 50 °C, clearly highlights that the electrode−electrolyte interfacial chemistry is affected by the Mg-based modifications. However, the nature of the chemistry of the active material and how it is affected by both the electrochemical reaction and chemical modifications remains unclear. Defining this chemistry could provide a mechanism that explains the behavior of the samples. Several authors have hinted at the reaction between the coated species and the surface of the active oxide.16,43 Further, Xiong et al. proposed the existence of Al3+ incorporation into the structure

Figure 5. (a) Specific discharge capacity and (b) Coulombic efficiency vs cycle number of a series of Li metal cells (3−5 replicas) containing the different LiNi0.5Mn1.5O4 prepared in this study as working electrodes, cycled at 50 °C at C/10.

decrease of the average discharge capacities, to 105−110 mAh/ g, from those seen at room temperature was noted for BL. These average values appeared to slightly decrease for Mg500 and Mg800 as well, but were still found to be within statistical error of the data at room temperature. All three average discharge capacities at 50 °C also fell within error of each other. These results are somewhat different from trends in the literature, in which the baseline material still produced a higher capacity than the coated material for the first few cycles.12,25,26 Despite the decrease in discharge capacity, the Coulombic efficiencies of Mg800 and Mg500 remained almost the same as those seen at room temperature, while the Coulombic efficiency of BL was much lower and very unsteady; the highest value was ∼82%, and the average was below 80%, ∼10% lower than what was seen at room temperature. Rate discharge capability tests were conducted and the various discharge capacities were normalized to those at C/10. This procedure highlighted the transport properties in each sample. Normalized rate capability comparisons of the samples at room temperature and 50 °C are shown in Figure 6, parts a and b, respectively, for rates ranging from C/10 through 5C. For both room temperature and 50 °C, there was an expectable and gradual decrease in capacity as the rate increased. At room 10600



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of LiMn1.97Ti0.03O4 as a result of the reaction with an Al2O3 coating, and hypothesized that the modification of the Mn−O bonding by Al 3+ was at the origin of the improved electrochemical properties.27 However, no direct chemical evidence of the effect of these ions on cycling was provided. Here, the electronic structure was directly probed at the most relevant states in the electrochemical reaction of LiNi0.5Mn1.5O4, those resulting from Ni−O bonds, using soft Xray absorption spectroscopy, a highly sensitive tool to the d states of the transition metal.31,44 The three samples were compared before and after full lithium deintercalation, at different temperatures. The Ni LII,III- and O K-edges spectra for all samples are shown in Figures 7 and 8, respectively. The

Figure 8. XAS at the O K-edge, collected in (a) TFY and (c) TEY mode, of BL (black), Mg500 (blue), and Mg800 (red) in their pristine state (marked as “p”), harvested from a Li metal cell charged to 5 V at room temperature (c), and at 50 °C (c_50). Pre-edge regions of TEY and TFY are shown in (c) and (d), respectively.

eV (LIII region), as well as a doublet centered at 869 and 870 eV (LII region). These signatures are consistent with the presence of Ni2+.45 No differences were observed between the data collected in total electron yield (TEY) and fluorescence yield (FY) modes, which indicates that surface and bulk were chemically identical. In turn, the broad peaks centered at 531.5 and 533.7 eV in the O K-edge pristine spectra were assigned to the unoccupied O 2p orbitals hybridized to either Ni or Mn 3d states, with the broader features at higher energy being due to bonding to 4s/4p states.46 When comparing the spectra of BL and Mg500, new features could be found at ∼542, 548, and 559 eV (see “*” in Figure 8a,c). These features are reminiscent of pure MgO spectra collected by others.47,48 However, a contribution from O 2p states in the spinel structure generated by the introduction of Mg2+ cannot be discarded, especially since very subtle changes are introduced when Mg−O bonds are part of a rock salt or spinel framework.49 Indeed, Mg−O signals were also observed in the O K-edge spectra of Mg800. The signals were very weak in the TEY data, which could be considered as the result of a decrease in concentration of Mg at the surface as it migrated into the subsurface or the bulk of the particle. The overall trends after full delithiation were similar in all cases. The changes in the Ni LIII region were dominated by the appearance of a large peak at 854.2 eV accompanied by an overall shift of the center of gravity toward higher energy,

Figure 7. XAS at the Ni LII,III-edge, collected in (a) TFY and (b) TEY mode, of BL (black), Mg500 (blue), and Mg800 (red) in their pristine state (marked as “p”), harvested from a Li metal cell charged to 5 V at room temperature (c), and at 50 °C (c_50), respectively.

pristine samples were labeled as “p”, whereas the fully delithiated samples are referred to as “c” and “c_50”, depending on whether the reaction was performed at room temperature or 50 °C, respectively. In all cases, the Ni LII,III spectra of the pristine samples showed a very prominent peak at 852 eV, a smaller one at 853.8 10601



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oxidized species than BL, especially when the reaction was performed at 50 °C. This observation was particularly quantifiable when comparing data at the O K-edge because neither of the BL spectra showed evidence of the 530.5 eV shoulder (Table 1). Among the samples with Mg-based modifications, the differences were more subtle, but Mg800 generally appeared to have higher ratios of oxidized species than Mg500 (Table 1 and in the Supporting Information Figures S7 and S8). Second, the TEY data of BL always revealed smaller peaks corresponding to oxidized species compared to FY (Figures 7 and 8, and Table 1). Again, this observation was most visible in the O K-edge data. Differences in these ratios were less obvious for Mg500 and Mg800, suggesting that the chemical gradient between bulk and surface was smaller. Finally, in the case of the Ni L-edge data of BL and Mg500, the 853.5 and 870.1 eV features were more pronounced at 50 °C than room temperature, at the expense of the peaks at higher energy (see the broken lines in Figure 7). These intermediate signals have been assigned by others to the presence of Ni3+, based on theoretical simulations of the spectra.45 Mg800 showed the highest spectral intensity at the highest energies (i.e., around 854 and 871 eV) when charged at 50 °C, which, in turn, was rather close to room temperature (Figure 7, and, for details, see Table S1 and Figure S7 in the Supporting Information). In other words, the Mg800 surfaces seemed to be the least reduced out of the three samples when the oxidation reaction was carried out at high temperature, consistent with the higher Coulombic efficiencies observed in Figure 5.

including a very minor shift of the signal at 852 eV and the rise of the background around 853.5 eV (see the broken line in Figure 7a). In turn, the LII doublet was replaced by a complex envelope with features at 869, 870.1 (see the broken line in Figure 7b), and 871.3 eV, also leading to a blueshift of the center of gravity. These observations are consistent with the oxidation of Ni during lithium deintercalation.44,45 These changes were accompanied by the rise of a shoulder at 530.5 eV (see the broken line in Figure 8b,d) and an upward shift of the center of gravity of the O 2p-Ni/Mn 4s/4p signals in the O Kedge spectra. The existence of new O 2p-Ni/Mn 3d states is a result of the increase in covalence of the Ni−O bond upon oxidation, which involves O taking an active role in the charge compensation mechanism.50 Indeed, signals at such low energies were observed in the O K-edge spectrum of Li2O2,51 supporting that they are due to electron-depleted oxide species. Similar observations can be found in literature reports of lithium deintercalation in layered oxides such as LiNi1/2Mn1/2O2.44,52 A detailed description of the bonding changes and charge compensation mechanisms is elusive, with different explanations put forth by different authors,44,52 and would probably require the support of computational simulations, especially in the case of Ni L-edges. However, it is generally agreed that the larger the portion of intensity at the 530.5 eV shoulder and the 854.2 features, the more oxidized the probed volume is. Upon close inspection, subtle, but important differences were found among spectra depending on the sample, the probed volumes (i.e., TEY and FY modes), and the temperature at which charging took place. Most visibly, discrepancies were found in the intensity ratios between the signals at 852 eV versus 854.2 eV, and between the 530.5 eV shoulder and the main pre-edge peaks in the Ni L- and O K-edge data, respectively. These ratios are representative of the degree of oxidation of the probed volume. They were quantified by spectral deconvolution of the whole Ni L-edge spectra, and of the pre-edge region, between 527 and 543 eV, in the O K-edge spectra. The results are presented in Table 1 and Table S1

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DISCUSSION Taken together, the XRD, XPS, SEM-EDS, and XAS results on the pristine samples indicate that Mg500 is formed by micrometric and submicrometric LiNi0.5Mn1.5O4 octahedral particles covered by uneven layers of MgO. In turn, Mg2+ was incorporated to a significant degree into the spinel surface and subsurface structure in Mg800. This incorporation must be compensated by either preferential Ni2+ segregation in the form of a rocksalt phase or formation of a Li and O defective spinel. The increase in the XRD peaks corresponding to a rocksalt structure when comparing Mg500 and Mg800 supports the first hypothesis and confirms that the chemical reaction between coating and spinel oxide happened. The Mg-based surface modifications slightly decreased the storage capacity of the active oxide at room temperature. The possible reason Mg500 had the lowest specific capacity is that the MgO coating and, especially, agglomerates present on the surface of the active material acted as a physical barrier that hindered both electron and Li transport. Similar observations were made by other groups when using Zr-based26 or Mg-based coatings on LiNi0.5Mn1.5O4.12 The utilization of Mg800, i.e., specific capacity, is higher than that of Mg500 in a statistically significant manner. This effect could be due to the incorporation of Mg into the spinel structure, which exposes a semiconducting surface that contains Li (i.e., ionic charge carriers), thus improving the charge transfer kinetics. In contrast, the lower capacity compared to BL was probably due to the enrichment of the spinel phase in Mg2+ at the expense of electrochemical active Ni2+, which segregates as a rocksalt phase that is known to have poor electrochemical properties.33 The result is a reduction of the theoretical capacity of the oxide.

Table 1. Summary of the Results of the Deconvolution of Soft XAS Spectra Corresponding to Materials Charged to 5 V in a Li Metal Half-Cell at Different Temperaturea I854/I852 (Ni LIIIedge)

% I530.5 (O K pre-edge) sample

RT (FY)

RT (TEY)

50 °C (FY)

50 °C (TEY)

RT (TEY)

50 °C (TEY)

BL Mg500 Mg800

1.8 2.0 1.9

0 0.7 3.0

1.7 2.9 3.2

0 2.0 2.6

0.47 0.64 0.68

0.52 0.65 0.69

a

% I530.5 corresponds to the ratio of intensity of the shoulder at 530.5 eV over the total intensity in the pre-edge region, from 527 to 536 eV. The changes at the Ni LIII-edge were represented as the ratio between the peaks at 854 and 852 eV. See text for details on the assignment of these signals. TEY = total electron yield detector; FY = fluorescence yield detector; RT = room temperature.

(Supporting Information), whereas representative graphical examples are provided in Figure S6 (Supporting Information). The clearest and most important difference was found when comparing the surface-sensitive spectra taken with the TEY detector (Table 1 and Figures 7b and 8c,d, as well as Figures S7 and S8 in the Supporting Information). In general, the Mgcontaining materials showed more prominent signals due to 10602



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especially at 50 °C, is an indication that the presence of the alkaline earth ions stabilizes these species against reaction with the electrolyte. The mechanism is not clear at this point and should be the subject of follow-up studies, but we hypothesize that since Mg2+ has a much lower charge density than Ni4+, it increases the charge density at the O surface terminations by reducing the amount of oxidized Ni per O atom.

Increasing the cycling temperature to 50 °C affected the uncoated, BL electrode to significantly greater extent than the modified oxides. The Coulombic efficiency trends were used to quantify their effect on the electrochemical stability of the electrode−electrolyte interface. This value was very stable when going from room temperature to 50 °C in Mg500 and Mg800, but notably dropped for BL. This drop implies an increase in the deleterious processes at the electrode surface, such as the formation of insulating species from electrolyte decomposition. As a result of degraded interfacial transport, the utilization of the material with increasing rates deteriorated, but Mg500 and Mg800 showed stable performance. Thus, this result is consistent with the reduced degradation of the electrode surface due to the presence of Mg-based modifications, and shows it is crucial to maintain electrochemical stability at higher temperatures. The significantly worse rate capability of Mg500 compared to Mg800 was ascribed to the presence of insulating MgO domains, which introduce electrical polarization at high rates. It indicates that the formation of semiconducting surfaces that contain Li is more favorable than growing purely resistive layers that degrade the rate of charge transfer. Ultimately, the precise mechanism by which performance improves would require direct probing of the different resistances in the cell, isolating the oxide working electrode from the contributions of the Li metal counter electrode, i.e., in three-electrode configurations.53 The XAS results offer an interesting picture of the chemical nature of the Ni−O bonds before and after the electrochemical reaction. It is clear that, in the absence of modifications, the surface of the spinel oxide is less oxidized than the bulk, which indicates that it participates in the process of the oxidation of the electrolyte. While this difference between surface and bulk has been highlighted in the literature for other positive electrode materials based on Ni,44 the universality of this reactivity has now been established for LiNi0.5Mn1.5O4. Kawaura et al. recently compared XAS data at the Ni K-edge for the surface and the bulk of LiNi0.5Mn1.5O4 thin films after full Li extraction, and did not report differences.54 It is probable that the subtle, but clear changes observed here by soft XAS are more difficult to resolve at the Ni K-edge because of the formal participation of the O electronic states in the redox chemistry, as proved in this study, and the higher sensitivity of the Ni Ledges to the d states of the transition metal.44 The fact that the surfaces of BL were even less oxidized when the temperature was raised and, thus, the Coulombic efficiency decreased provides additional support to this explanation. Active participation of the electrode surface species in the mechanism of electrolyte decomposition on Ni-containing spinel electrodes has also been put forth by others in the literature.55,56 The XAS data here provide direct proof of this active role. The results also allow us to propose a likely mechanism of reaction, by which the spinel oxide participates in the decomposition reaction through the generation of highly polarizing Ni4+ ions. In turn, these ions extract electron density from O2−, which are the surface terminating ions in this morphology.57 Electronpoor O terminations would be prone to an electrophilic attack on the electron-rich alkylcarbonate molecules that constitute the electrolyte.58 The Mg-based surface modifications led to more oxidized surfaces than in the bare material, indicating that they were more stable toward decomposing the electrolyte. The fact that Mg800, the material with the highest amount of Mg2+ in the spinel surface and subsurface structure, showed the highest proportion of oxidized species in the TEY data,

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CONCLUSION The surface chemistry consequences of modifying LiNi0.5Mn1.5O4 using Mg2+ vis-à-vis electrochemical performance was studied. A Mg precursor was coated on micrometric spinel particles, followed by annealing at 500 or 800 °C. The results showed a rough coating with an inhomogeneous distribution of agglomerates, most likely MgO, for the former, which was modified after treatment at 800 °C through the incorporation of Mg2+ into the spinel subsurface. An electrochemical study showed that the presence of Mg does not improve performance at room temperature, but is crucial at higher temperature. Among the two modified materials, the best results were obtained following annealing at 800 °C, in terms of capacity, Coulombic efficiency, and rate capability at 50 °C. The electrochemical data, together with analysis of X-ray absorption spectra collected at the Ni LII,III- and O K-edges, revealed a possible mechanism for the improvement with chemical bonding underpinnings. It is based on the generation of highly electron depleted Ni−O species on the electrode surface as the culprit for electrolyte decomposition reaction. The incorporation of Mg2+ to the spinel surface and subsurface structure showed evidence of being a chemical stabilizer of these electron-depleted states, possibly by increasing the charge density around O surface terminations. It is the first time that a plausible chemical explanation, based on data with direct chemical insight, has been provided to the beneficial effect of coatings during cycling at high potential. We believe this work contains two important general lessons that could guide the rational design of stable high-voltage materials. First, coatings that are semiconducting in nature and contain lithium are preferred to reduce charge (ion and electron) transfer losses. In this sense, structures based on core−shell gradients of composition59 would seem to provide a way forward. Second, small amounts of a redox inactive, strongly ionic species such as Mg2+ on the surface of an active material could have a significant stabilizing effect by changing the nature of chemical bonding vis-à-vis possible redox activity with the electrolyte. These rules are thought to be of potential relevance to other phases of interest as positive electrodes in Liion batteries, including, but not limited, to LiNi0.5Mn1.5O4.

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ASSOCIATED CONTENT

* Supporting Information S

Figure S1, TGA and DSC for the as-prepared Mg-based precursor in the absence of spinel powder; Figure S2, results of SEM-EDS analysis of Mg500; Figure S3, results of XPS analysis of Mg500 and Mg800 as received and after sputtering the surface for 5 min; Figure S4, results of SEM-EDS analysis of Mg800; Figure S5, voltage vs specific capacity and dQ/dV traces during the first charge and discharge of Li metal cells containing BL, Mg500, and Mg800 as working electrodes; Figure S6, representative examples of the spectral deconvolutions of the O K pre-edge, Ni LIII, and LII regions of the spectra in Figures 7 and 8; Figure S7, overlaid Ni LII,III-edge spectra, shown in stacked form in Figure 7, collected with a TEY 10603



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detector, highlighting the differences in center of gravity between samples; Figure S8, zoom of the region around 531 eV in the O K-edge XAS data shown in Figure 8, highlighting the differences in intensity of the 530.2 eV shoulder depending on the sample; and Table S1, a listing of the results of the spectral deconvolution of the Ni LII,III-edge spectra. This material is available free of charge via the Internet at http://pubs.acs.org.

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AUTHOR INFORMATION

Corresponding Author

*E-mail: [email protected]. Tel: +1(312)-355-4309. Present Addresses ⊥

Department of Chemistry & Biochemistry, University of California Los Angeles, Los Angeles, CA 90095. ¶ Crystal, 6752 Baymeadow Drive, Glen Burnie, MD, 21060. Notes

The authors declare no competing financial interest.

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ACKNOWLEDGMENTS This work was supported by the Assistant Secretary for Energy Efficiency and Renewable Energy, Office of Vehicle Technologies of the U.S. Department of Energy (DOE) under Contract No. DE-AC02-05CH11231, as part of the Batteries for Advanced Transportation Technologies (BATT) Program. Portions of this research were carried out at the Stanford Synchrotron Radiation Lightsource, a Directorate of SLAC National Accelerator Laboratory and an Office of Science User Facility operated for the U.S. Department of Energy Office of Science by Stanford University. G.M.N. was supported by LBNL through the Science Undergraduate Laboratory Internship program from the DOE. The authors wish to thank Dr. Dennis Nordlund (SSRL) for his assistance during the XAS measurements, and Dr. Marca M. Doeff (LBNL) for valuable interactions. This document was prepared as an account of work sponsored by the United States Government. While this document is believed to contain correct information, neither the United States Government nor any agency thereof, nor the Regents of the University of California, nor any of their employees, makes any warranty, express or implied, or assumes any legal responsibility for the accuracy, completeness, or usefulness of any information, apparatus, product, or process disclosed, or represents that its use would not infringe privately owned rights. Reference herein to any specific commercial product, process, or service by its trade name, trademark, manufacturer, or otherwise, does not necessarily constitute or imply its endorsement, recommendation, or favoring by the United States Government or any agency thereof, or the Regents of the University of California. The views and opinions of authors expressed herein do not necessarily state or reflect those of the United States Government or any agency thereof or the Regents of the University of California.

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Surface Chemistry Consequences of Mg-Based Coatings on LiNi0

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