Oxygen-Release-Related Thermal Stability and Decomposition

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Oxygen-Release-Related Thermal Stability and Decomposition Pathways of LixNi0.5Mn1.5O4 Cathode Materials Enyuan Hu, Seong-Min Bak, Jue Liu, Xiqian Yu, Yongning Zhou, Steven N. Ehrlich, Xiao-Qing Yang,* and Kyung-Wan Nam* Brookhaven National Laboratory, Upton, New York 11973, United States S Supporting Information *

ABSTRACT: The thermal stability of charged cathode materials is one of the critical properties affecting the safety characteristics of lithium-ion batteries. New findings on the thermal-stability and thermal-decomposition pathways related to the oxygen release are discovered for the high-voltage spinel LixNi0.5Mn1.5O4 (LNMO) with ordered (o-) and disordered (d-) structures at the fully delithiated (charged) state using a combination of in situ time-resolved X-ray diffraction (TRXRD) coupled with mass spectroscopy (MS) and X-ray absorption spectroscopy (XAS) during heating. Both o- and dLixNi0.5Mn1.5O4, at their fully charged states, start oxygenreleasing structural changes at temperatures below 300 °C, which is in sharp contrast to the good thermal stability of the 4Vspinel LixMn2O4 with no oxygen being released up to 375 °C. This is mainly caused by the presence of Ni4+ in LNMO, which undergoes dramatic reduction during the thermal decomposition. In addition, charged o-LNMO shows better thermal stability than the d-LNMO counterpart, due to the Ni/Mn ordering and smaller amount of the rock-salt impurity phase in o-LNMO. Two newly identified thermal-decomposition pathways from the initial LixNi0.5Mn1.5O4 spinel to the final NiMn2O4-type spinel structure with and without the intermediate phases (NiMnO3 and α-Mn2O3) are found to play key roles in thermal stability and oxygen release of LNMO during thermal decomposition. KEYWORDS: Li-ion battery, safety, high voltage spinel, in situ X-ray diffraction, EXAFS

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INTRODUCTION 4V-spinel LiMn2O4 is one of the important cathode materials for lithium-ion batteries (LIBs) intended as the power source of electric vehicles due to the straightforward synthesis, high ionicand electronic- conductivity, excellent thermal stability, and low cost.1,2 Recently, one of its derivatives, LiNixMn2−xO4, denoted as LNMO, (where x is around 0.5) has attracted a lot of research attention as a promising high-energy density cathode material based on its higher operating voltage at ∼4.7 V vs Li+/ Li compared to the parent material, LiMn2O4.3 On the other hand, the poor cycle and calendar life of LNMO, especially at elevated temperatures, still remains one of the major challenges in its widespread usage. Extensive research has addressed some key factors determining its capacity and rate performance, such as cation ordering,4−7 route of synthesis,8,9 stoichiometry,10,11 heat treatment,12 particle morphology,13,14 particle size,12,15 transition metal substitutes,16,17 and the Li-insertion/deinsertion mechanism5,18,19 of this material. Among these factors, the heat treatment procedure is especially important because LNMO is generated in two different kinds of structures, depending on the annealing procedure adopted after the 900 °C solid-state reaction. Reportedly, spinel-structured LNMO, formed through the solid-state reaction, loses oxygen between 700 and 900 °C, causing the partial transformation of a spinel phase to a rock-salt phase. The majority of the rock-salt © 2013 American Chemical Society

structure can be transformed back to a spinel structure during cooling through the material regains oxygen below 700 °C. However, if the cooling is not slow enough, the amount of oxygen regained is low and a significant amount of the rock-salt impurity phase will be retained. Under such conditions, the disordered phase is formed, in which Mn4+ and Ni2+ cations are randomly distributed in the transition metal layer, as shown in the upper right in Figure 1, with the space group Fd3̅m (to be referred as d-LNMO).11 On the other hand, if the cooling of the material is slow enough (in some cases, a constant temperature-annealing duration below 700 °C is used), the reincorporation of oxygen is almost complete, and the rock-salt impurity residues will be negligible. Under the latter conditions, the ordered phase is formed in which Mn4+ and Ni2+ cations are located at certain different sites in the crystal lattice as shown in Figure 1 lower right, with the space group P4332 (to be referred as o-LNMO). Although most researchers presently believe that the cycling performance of the disordered d-LNMO is better than that of the ordered o-LNMO, there are more and more reports demonstrating that the o-LNMO can display a satisfactory cycling performance, while preserving other Received: October 15, 2013 Revised: November 25, 2013 Published: December 10, 2013 1108

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during heating in terms of changes in crystal structures, chemical compositions, and valence of the transition metals. The in-depth understanding of the LiNi0.5Mn1.5O4 materials gained from this study may provide valuable information for developing advanced cathode materials with improved thermal stability and durability over the state-of-the-art materials. Correlation between the structural changes and the oxygen release during thermal decomposition investigated in this study may also provide some insight on understanding the (de)lithiation mechanism of Li-rich layer-structured cathode materials (e.g., Li1.2Ni0.2Mn0.6O2) where the electrochemically driven oxygen evolution was observed at room temperature by high voltage charging.31

Figure 1. The crystal structure of LNMO with the top view of transition metal layer shown on the right. The upper right part illustrates the d-LNMO case, in which Ni and Mn are randomly distributed in the layer. The lower right part illustrates the o-LNMO case, in which ordering is formed between Ni and Mn. The red spheres indicate lithium ions residing in tetrahedral sites.

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EXPERIMENTAL SECTION

Synthesis and Electrochemical Tests. The d-LNMO and oLNMO were synthesized by a conventional solid-state method. Li2CO3, MnO2, and NiO were mixed in stoichiometric amounts and ball-milled for 4 h. The mixture then was pelletized for further calcination. When d-LNMO was synthesized, the pellets were calcined at 900 °C for 24 h in air and then cooled to room temperature at a rate of 1 °C min−1. o-LNMO pellets were calcined at 900 °C for 24 h in air before annealing them at 700 °C for 24 h first, then at 650 °C for 24 h, and finally at 600 °C for 48 h. The pellets were cooled to room temperature at a rate of 1 °C min−1. The d- and o-LNMO cathodes were fabricated by slurry-casting 80% active material, 10% of carbon black (Chevron), and 10% of PVDF (Kureha) onto an Al foil current collector. 2032-type coin cells were assembled using a Li foil anode, a separator (Celgard), and an electrolyte composed of 1.2 M LiPF6 in EC (ethyl carbonate) and DMC (dimethyl carbonate) solution (3:7 EC/DMC volume ratio). For the electrochemical cycling test, the coin cells were charged at a C/10 rate, assuming a theoretical capacity of 140 mAh g−1. In situ TR-XRD Combined with Mass Spectroscopy. After the d- and o-LNMO samples were electrochemically delithiated (i.e., fully charged), the cells were disassembled and powder samples of the charged cathode materials were collected from the electrode. Approximately 3.5−4.0 mg of these powders were loaded into a glass capillary with an inner diameter of 0.7 mm in which one end of the capillary (i.e., the inlet) was connected to a He carrier gas source and the other end (i.e., the outlet) to a residual gas analyzer/mass spectrometer (RGA200, Stanford Research Systems). The TR-XRD data were collected using the thermal stage at beamline X7B (λ = 0.3196 Å) of the National Synchrotron Light Source (NSLS) in the transmission mode, while the MS signals were simultaneously collected as the sample was heated from room temperature to 375 °C at a rate of ∼2.0 °C min−1. Details of the preparation method were described in our previous publications.23,25 To ease comparisons with other published results using conventional X-ray sources, we plotted the XRD patterns using converted 2θ angles corresponding to the Cu Kα (λ = 1.54 Å) radiation wavelength. In situ X-ray Absorption Spectroscopy. In situ XAS experiments were carried out at beamline X18A of the NSLS in transmission mode using a Si (111) double-crystal monochromator detuned to 35− 45% of its original maximum intensity to eliminate the high-order harmonics. To obtain good-quality data, the XAS spectra were collected at room temperature after heating the fully delithiated (charged) cathode electrode to each of the predetermined high temperatures. Since most structural changes occurred during heating are irreversible, the spectra collected at room temperature basically reflect the sample’s state at high temperature before cooling. Mn and Ni K-edge XAS spectra collected at high temperature (370 °C) and at room temperature (25 °C) after cooling down (Supporting Information, Figure S1) were identical, confirming no structural changes occurred after cooling down the samples. The fully charged cathode electrodes were first washed by the DMC solvent and then mounted on the sample holder of a specially designed heating cell and heated under the protection of He gas. X-ray absorption near edge

properties like higher average operating voltage and larger capacity than those of the d-LNMO.6,20 Unlike the widely studied electrochemical performance and reaction mechanism, the thermal stability of LNMO, which could greatly impact the safety of LIBs, has received little attention. This lack of interest could probably be attributed to the assumption that the excellent thermal stability of the delithiated LNMO can be naturally inherited from its parent material LiMn2O4, for which only a subtle structural rearrangement takes place without the oxygen release up to 500 °C in the fully delithiated state.21,22 Therefore, LixMn2O4 has been regarded as a thermally safer cathode material than layered materials, such as LixCoO2 , Lix Ni0.8 Co0.15 Al0.05 O 2, and LixNi1/3Co1/3Mn1/3O2.23−25 All such structurally layered materials undergo a series of phase transitions with accompanied oxygen release below 300 °C in their charged states. However, for LNMO, what was overlooked is that when a quarter of the Mn is replaced by Ni, the thermodynamics of the material inevitably changes, yielding a very different thermal stability than its parent LiMn2O4. Unfortunately, little research has been published on the thermal stability of LNMO materials and their doped derivatives; the research focus has been on their reactivity with the electrolyte using calorimetric measurements thus far.26−29 There are very few studies correlating thermal stability neither with structural differences (ordered or disordered) nor with oxygen-releasing structure changes during heating for LNMO. The systematic studies we report in this paper are designed to reveal and clarify these correlations. To understand the thermal stability of both d- and o-LNMO in the delithiated state, we applied a combination of in situ synchrotron time-resolved X-ray diffraction (TR-XRD) coupled with mass spectroscopy (MS) and in situ X-ray absorption spectroscopy (XAS) during heating. This combination allowed us to simultaneously monitor, during thermal decomposition, the phase transformations (by TR-XRD)25 and the accompanying gas evolution (e.g., oxygen by MS)23 as well as the localand electronic-structural changes with an elemental selective capability (by XAS).30 Comparing the different structural behaviors between delithiated d- and o-LNMO during heating, we are able to discover some thermodynamic origins affecting the materials’ different electrochemical stabilities (i.e., cycle life). Through this systematic investigation, we demonstrate the mechanism of thermal decomposition and the oxygen release behavior of (electrochemically) delithiated d- and o-LNMO 1109

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structure (XANES) and extended X-ray absorption fine structure (EXAFS) data were analyzed by the ATHENA software package.32 The photoelectron energy origin, E0, is chosen at the first inflection point of the absorption edge jump. The extracted EXAFS signal, χ(k), was weighted by k3 to emphasize the high-energy oscillations and then Fourier-transformed in k-ranges of 3.0−13.5 Å−1 for Mn and 3.0−11.0 Å−1 for Ni using a Hanning window function to obtain the magnitude plots of the EXAFS spectra in R-space (Å). The Fourier-transformed peaks were not phase corrected, and thus the actual bond lengths are approximately 0.2−0.4 Å longer. The filtered Fourier transforms of EXAFS spectra were then fitted using theoretical single scattering paths generated with the FEFF 6.0 ab initio simulation code. Two model structures of LiM2O4 (M = Mn or Ni) and NiMn2O4 spinel phases were used for the fitting depending on the heating temperatures. In the case of the fitting with the LiM2O4 model, the Li contribution was ignored. The amplitude reduction factors, S02, were determined to be 0.70 for Mn and 0.80 for Ni from the preliminary fitting sessions, and then fixed during the final fitting unless noted otherwise. The same inner shell potential shift (ΔE) was shared for the M−O and M−M shells (M = Mn and Ni) whereas separate fitting parameters of the bond distance (R) and Debye− Waller factor (i.e., mean square disorder, σ2) were used for each shell. The coordination numbers were fixed to the crystallographic values unless otherwise noted. Detailed fitting ranges, parameters, and constrains used in the fitting are described in the Supporting Information (Figure S5−S7 and Table S2−S3).

Figure 2. Le Bail fitting of (a) charged d-LNMO and (b) charged oLNMO. The black circles, red lines, and blue lines correspond to the measured, calculated, and differentiated XRD patterns, respectively. Purple labels represent for the main phase peaks and green ones are for the rock-salt impurity. The Fd3m ̅ space group was used for the charged d-LNMO and the P4332 space group for the charged oLNMO. We note that the charged d-LNMO has more rock-salt impurities than charged o-LNMO, as shown in the insets.

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RESULTS AND DISCUSSION X-ray Diffraction and Electrochemical Characterization of Disordered (d)- and ordered (o)-LNMO. To ensure that the samples used in this study were consistent with materials reported in the literature, electrochemical cycling and XRD Rietveld refinement were carried out first;33−35 the results are shown in the Supporting Information (Figure S2 and Table S1). They demonstrate that well-defined d- (disordered) and o(ordered) LNMO were successfully synthesized. In addition to the main plateau around 4.7 V, the d-LNMO features another short plateau around 4.0 V whereas the o-LNMO does not have it (Supporting Information, Figure S2), confirming the formation of representative spinel LNMO with disordered and ordered structures, respectively.11,17,36 The weak superlattice peaks arising from Ni/Mn ordering are clearly visible, besides the main spinel peaks in the XRD pattern of the oLNMO (Supporting Information, inset of Figure S2c), but not in the d-LNMO. The o- and d-LNMO are illustrated in Figure 1. In the lower right of this figure, the crystal structure of oLNMO with space group P4332 is shown, wherein Ni and Mn ions in the transition-metal layer occupy the octahedral 4b and 12d sites, respectively, forming an ordered structure. In the upper right of Figure 1, for d-LNMO with space group Fd3̅m, the Ni and Mn ions are randomly distributed on these octahedral sites. Rietveld refinement of the XRD patterns indicated the existence of rock-salt impurity phase, Ni6MnO8,10 in both d- and o-LNMO, in addition to the main spinel structure. However, the amount of this rock salt phase is much less in o-LNMO (a weight faction of 0.83%) than in d-LNMO (a weight fraction of 3.61%), agreeing well with results reported by others.17,37 The lattice parameter of d-LNMO was refined to be 8.206 Å, larger than the 8.171 Å of o-LNMO likely due to larger percentage of Mn3+ in d-LNMO than in o-LNMO. Before discussing the TR-XRD result upon heating, structures of both d- and o-LNMO in the fully charged state (4.9 V vs Li/Li+) were analyzed using room-temperature XRD. The results are shown in Figure 2. The basic structures of the disordered (Fd3̅m) and ordered (P4332) were preserved after

charge but with reduced lattice parameters of 8.043 Ǻ for dLNMO and 8.039 Ǻ for o-LNMO, reflecting the lattice contractions caused by the oxidation of Ni2+ to Ni4+. It is interesting to note that cation ordering in o-LNMO was also preserved even after the full delithiation (charge), as evidenced by the superlattice peaks (Supporting Information, Figure S3). The larger amount of the rock-salt impurity phase in d-LNMO than in o-LNMO in the charged samples is also apparent, as depicted in the insets of Figure 2a,b, revealing the electrochemically inactive nature of this impurity phase. In situ TR-XRD Combined with Mass Spectroscopy. To monitor the structural changes and release of gaseous oxygen during thermal decomposition, we used TR-XRD combined with simultaneously measured mass spectroscopy (TR-XRD/ MS) for both charged d- and o-LNMO during heating up to 375 °C; the results are shown in Figure 3. The phase transitions and accompanying oxygen release can be tracked closely during heating. There was clear evidence of thermal-decompositionrelated phase transitions accompanied with noticeable oxygen release below 250 °C for both d-LNMO and o-LNMO. These results are in sharp contrast to the conventional spinel LixMn2O4 at the fully charged (x = ∼0.0) state, where no oxygen-release peak was observed below 375 °C (Supporting Information, Figure S4). This tells us that when 25% of the Mn in conventional spinel LiMn2O4 is replaced by Ni, the thermal stability of this high voltage LNMO is greatly changed for both d-LNMO and o-LNMO. Therefore, their related structural changes during heating need to be studied thoroughly. Before detailed data analysis, it is evident from Figure 3 that the charged d- and o-LNMO mainly follow a similar route of structural change during heating: Above 100 °C, all XRD peaks for both shifted toward low angles as the temperature rose, indicating the expansion of the lattice parameter; above 200 °C, the appearance and growth of peaks indexed as (220) and (422) were observed, suggesting that the transition metals migrated from octahedral sites to tetrahedral ones in the spinel structure;25 they both experienced major structural changes 1110

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Figure 3. In situ XRD patterns combined with simultaneously measured mass spectroscopy (MS) data that trace the release of gaseous oxygen of charged (a) d-LNMO and (b) o-LNMO during heating up to 375 °C. Left, in situ XRD patterns; right, profile of oxygen release.

to be analyzed were formed at low temperatures, therefore are full of defects, making it difficult to fit the atomic positions and site occupancies using the Rietveld method. The second is based on the fact that the XRD data we collected during heating were not corrected for the preferred orientation (without sample spinning) that is required for Rietveld. Considering that the phase identification and matching of lattice parameters were most important to us in this study, the Le Bail fitting method is a better choice. We attempted profile matching of the XRD pattern collected at 375 °C for the oLNMO with the structure database to identify the structure and phases (Figure 4).

accompanied by oxygen release, followed by the formation and growth of new phases. At 375 °C, the shape of the XRD patterns for the d- and o-LNMO are similar, but the intensities of the peaks differ. The most interesting difference between charged d-LNMO and o-LNMO is the onset temperature of oxygen release and the related phase transitions. The charged dLNMO started to release oxygen at a temperature as low as ∼210 °C, accompanying an abrupt change of the XRD patterns (Figure 3a), whereas this temperature for charged o-LNMO was pushed higher to ∼250 °C (Figure 3b). This clearly demonstrates that o-LNMO is structurally (i.e., thermodynamically) more stable than d-LNMO at the fully charged (delithiated) state. As mentioned, d-LNMO contains more rock-salt impurity than the o-LNMO in both the pristine and charged samples (Supporting Information, Figure S2; Figure 1). In our previous study on thermal stability in the layered material systems, we found that the rock-salt phases on the surface of the particle can act as nucleation centers for the phase transitions upon thermal decomposition.38 The fact that d-LNMO with a larger amount of rock-salt impurity phase has poorer thermal stability than o-LNMO suggests that this impurity phase may play an important role in forming nucleation centers, so facilitating the structure degradation and accelerating phase-transformation and oxygen release at lower temperatures for d-LNMO. Another interesting observation is the preservation of superlattice peaks during heating for the charged o-LNMO up to the onset temperature of oxygen release (Supporting Information, Figure S3). Although the peak width broadened and its position shifted to lower angles as the temperature increased, this characteristic (110) superlattice peak still is seen up to 235 °C, just before oxygen is released. As expected, this peak is not seen in the d-LNMO, either in the pristine or in the charged sample (Supporting Information, Figure S3). We analyzed phase evolution during thermal decomposition in more detail using the Le Bail fitting method.39 There were two major reasons for our preference for Le Bail fitting over Rietveld refinement. First, the structures

Figure 4. Le Bail fitting of XRD pattern collected at 375 °C for charged o-LNMO using three phases with their PDF cards shown above. (a) NiMnO3, (b) α-Mn2O3, and (c) NiMn2O4. Their structures are illustrated on the right side. 1111

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A rough identification suggests that it constitutes a combination of the spinel phase with much larger lattice parameter than the pristine LiNi0.5Mn1.5O4 and an ilmenite NiMnO3 phase.40 However, two-phase fitting yielded some lattice parameters that do not make much sense. Therefore, we integrated a third phase, α-Mn2O3, which is commonly seen in the decomposition of Ni and Mn containing oxides.41−45 The three-phase fitting gave us satisfactory results. The spinel phase, having much larger lattice parameter than the pristine LNMO’s one, is confirmed to be NiMn2O4. Structures of these three phases are shown in the right panel of Figure 4. Ilmenite NiMnO3 is a derivative of the corundum structure, with Ni2+ and Mn4+ in alternating layers.40 Therein, the Ni- and Mnoctahedrons share faces but the space above and below this face-sharing pair is vacant to accommodate the octahedron shifted by electrostatic repulsion between cations. α-Mn2O3 is related to the CaF2 structure, from which it may be derived by removing one-quarter of the anions (e.g., oxygen in the case of α-Mn2O3) followed by some atomic rearrangement.46 The 6coordinated Mn atoms are of two types, as indicated in Figure 4, where one type is shown in polyhedra and the other in bonding sticks. Spinel NiMn2O4 differs from LiNi0.5Mn1.5O4 in the way that its tetrahedral sites are occupied by transition metal (mainly Mn2+) ions, whereas in LiNi0.5Mn1.5O4, the sites are exclusively occupied by lithium ions. However, because NiMn2O4 formed at low temperatures, a large deficiency of cations and defects existed in both d- and o-LNMO samples. In this paper, we refer the low-temperature-formed NiMn2O4 phase as the NiMn2O4-type spinel to emphasize characteristic occupation of its tetrahedral sites by transition metals. To have a better idea of the thermal decomposition mechanism and the pathway of charged LNMO, the evolution of each phase identified above by Le Bail fitting was plotted as functions of temperature (Figure 5) by integrating the intensity of the characteristic peak representing each phase. These intensities, expressed in arbitrary units, can only be used to indicate the increase and decrease of each individual phase during heating and not for comparison with other phases. It is fairly straightforward to select the (220) peak for identifying the NiMn2O4-type spinel. However, for the NiMnO3 and α-Mn2O3, it was very difficult to discern a well-resolved single characteristic peak in either of them. Accordingly, we chose the peak around 50° that is a combination of (20−4) peak of the NiMnO3 and (431) peak of α-Mn2O3. In this way, qualitative information about how these two phases change during heating was obtained. As shown in Figure 5a, during heating of d-LNMO, the NiMnO3 and α-Mn2O3 phases first increased, and then decreased while the NiMn2O4-type spinel that grew slowly first, began growing rapidly at the expense of the decreasing NiMnO3 and α-Mn2O3 phases. This finding indicates the important roles of the intermediate phases, NiMnO3 and α-Mn2O3, during the thermal decomposition of dLNMO. In contrast, for o-LNMO, the NiMn2O4-type spinel phase grew faster first while the formation of the NiMnO3 and α-Mn2O3 phases was insignificant; this was followed later a slower rate of growth of the NiMnO3 and α-Mn2O3 phases evolved with increasing temperatures, indicating the less important roles of the intermediate phases in the case of oLNMO. In both cases, the growth of the NiMnO3 phase follows the process of oxygen release, implying a close relationship between them. This may not be surprising, as is evident later in the X-ray absorption section, that rapid oxygen release is accompanied by the rapid reduction of Ni from Ni4+

Figure 5. Evolution of characteristic peaks of different phases for charged (a) d-LNMO and (b) o-LNMO. Corresponding oxygen release profiles (red circle) are plotted for comparison. NiMnO3 peaks and α-Mn2O3 peaks are not well resolved and they are integrated together.

to Ni2+. Indeed, the abundance of Ni2+ favors the formation of NiMnO3 phase as it adopts the Ni2+ and Mn4+ cation configuration. In situ X-ray Absorption Spectroscopy (XAS). Complementary to the information about the average changes in crystal structure and phase transitions as well as gas evolution during thermal decomposition obtained from in situ TR-XRD/MS, XAS can provide information about the changes in local- and electronic-structure surrounding the absorbing atoms in an elemental-selective way. This is very helpful in identifying which element is mainly responsible for thermal instability. Ni and Mn K-edge X-ray absorption near edge structure (XANES) data for delithiated d- and o-LNMO during heating up to 370 °C are shown in Figures 6 and 7, respectively. The most notable change in Ni K-edge XANES spectra (Figure 6) is the inhomogeneous edge-shifting to lower energy (i.e., the decrease in average oxidation state of Ni); this sudden jump of edge shift divides the series of spectra into two parts. It occurs at some temperature between 220 and 270 °C for the charged dLNMO, and between 270 and 320 °C in the charged o-LNMO. This feature is clearly shown in Figure 6c where the changes in edge position are plotted as a function of temperature. The edge positions indicated by dotted lines in Figure 6c for the reference compounds with known oxidation states of Ni4+, Ni3+, and Ni2+ demonstrate the reduction path of Ni from Ni4+ to Ni2+ in both charged d- and o-LNMO upon thermal decomposition. The rapid reduction in the oxidation state of Ni (i.e., edge jump to lower energy) is well synchronized with the onset of the massive release of oxygen seen at ca. 200 °C for d-LNMO, and ca. 250 °C for o-LNMO in satisfying electriccharge neutrality during the course of the thermal decomposition. This result also confirms that the delithiated o-LNMO better resists oxygen release and structural decomposition 1112

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during heating than does its delithiated disordered counterpart (Figure 6c). The Mn K-edge XANES results depicted in Figure 7 showed only a moderated edge shift to lower energy for both samples, even during oxygen release, compared with the Ni K-edge results. The Mn reduction commenced at higher temperatures and ended with lesser extents than did the Ni reduction, demonstrating the structural-stabilizing role of Mn as we previously reported.25 For example, for charged d-LNMO at 270 °C, the average oxidation state of Ni was reduced significantly from Ni4+ to states close to Ni3+ (Figure 6c) while, at the same temperature, the average oxidation state of Mn was maintained close to its initial state of Mn4+ (Figure 7c). At 370 °C, after major oxygen release, the average oxidation state of Mn in both charged d- and o-LNMO only fell to somewhere between Mn3+ and Mn4+,while that for Ni was reduced to close to Ni2+. This finding suggests that the release of oxygen was initiated by the Ni4+ to Ni2+ reduction for LNMO upon heating, so resulting in highly disordered and defective structures and decomposition compounds, which, in turn, triggered the reduction of Mn4+ ions. Figure 7c shows that the charged o-LNMO has a higher average oxidation state of Mn than the charged d-LNMO at 370 °C, confirming the better thermal stability of the former than the latter in the charged state. The intensity variation of the pre-edge peak in the XANES spectra, shown in the insets of Figures 6 and 7, offers further insights into the changes in the local symmetry of Ni and Mn ions in the charged LNMO during thermal decomposition. The pre-edge peak is associated with the dipole-forbidden 1s → 3d electronic transition, which is partially allowed as a quadrupole electronic transition when the 3d and 4p orbitals hybridize due to the lack of an inversion center symmetry environment or structural distortion in the local symmetry between transition metals and oxygen coordination.47 Therefore, the XANES spectra usually show a much stronger pre-edge peak intensity when transition metal ions occupy the tetrahedral sites rather than octahedral ones due to the lack of an inversion center in the tetrahedral coordination symmetry. The intensity of Ni preedge (insets of Figure 6) for both of the charged d- and oLNMO remains quite low at all temperatures, suggesting that the Ni ions mainly remain in the octahedral sites at all temperatures. In contrast, the intensity of Mn pre-edge (insets of Figure 7) for both o- and d-LNMO samples increased significantly with increasing temperatures. On the basis of analysis of the TR-XRD data, the NiMn2O4-type spinel was formed after oxygen release. The great enhancement in the intensity of the Mn pre-edge reveals that Mn cations, mainly Mn2+, are the occupants of tetrahedral sites in the newly formed NiMn2O4-type spinel phase during thermal decomposition. To study local structural variations around Ni and Mn during thermal decomposition, in situ extended X-ray absorption fine structure (EXAFS) spectra for both charged d- and o-LNMO during heating up to 370 °C were analyzed. The Fouriertransformed (FT) EXAFS spectra (k3-weighted in k-space but not phase-corrected FT, causing shorter bond lengths in the plots than for the real ones) at Ni and Mn K-edge are shown, respectively, in Figures 8 and 9. The FT magnitudes of pristine d- and o-LNMO and (stoichiometric) spinel NiMn2O4 are also plotted in the bottom panels as references. The first peak at approximately 1.5 Å in Figure 8 is assigned to the single scattering path from Ni to the closest oxygen atoms (i.e., the Ni−O bond). The second peak at about 2.4 Å is assigned to Ni

Figure 6. In situ Ni K-edge XANES spectra of fully charged (a) dLNMO and (b) o-LNMO during heating up to 370 °C. Insets show the detailed feature of pre-edge region. (c) Variations of the Ni K-edge positions (defined as the energy at half height of the energy step at main edge).

Figure 7. In situ Mn K-edge XANES spectra of fully charged (a) dLNMO and (b) o-LNMO during heating up to 370 °C. Insets show the detailed feature of pre-edge region. (c) Variations of the Mn Kedge positions (defined as the energy at half height of the energy step at main edge).

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to the nearest transition metal atoms (Ni and Mn) occupying the octahedral sites (i.e., Ni−Moct). The third peak at about 3.2 Å in the blue curve for the NiMn2O4 is assigned to Ni to the second nearest transition-metal atoms (mostly Mn) occupying the tetrahedral sites (i.e., Ni−Mntet). It can be seen that the Ni−O bond length contracted from the pristine to fully charged state (25 °C), indicating the Ni2+ to Ni4+ oxidation. This Ni−O bond length abruptly expanded between 220 and 320 °C upon oxygen release, indicating a marked reduction of Ni4+. The second peak follows the same trend as the first one. Interestingly, the abrupt expansion in bond length for both Ni−O and Ni−Moct occurred between 220 and 270 °C for the disordered LNMO (black curve) but was pushed to higher temperatures of between 270 and 320 °C for the ordered one (red curve). Following this abrupt expansion, the Ni−Mntet peak emerged together, with other two characteristic peaks (marked by a blue + sign), pointing to the formation of the NiMn2O4-type spinel phase and the migration of Mn cations (mostly Mn2+) to the tetrahedral sites. For the Mn K-edge results shown in Figure 9, there were negligible alterations in the Mn−O bond length from the pristine to fully charged state, indicating that Mn in both such samples was at the Mn4+ oxidation state. For the spectral changes during heating, the expansion of the Mn−Moct bond length and the emergence of the Mn−Mntet (together with additional characteristic peaks marked by a blue + sign) between 220 and 320 °C followed the same trend as the Ni Kedge results in Figure 8, revealing the lower temperature of formation of the NiMn2O4-type spinel in the d-LNMO than in the ordered one. However, the changes in the Mn−O bond length during heating are not as clear as those in Figure 8 for Ni−O, due to the complication of the varying sites of Mn (in different Mn-containing phases such as LNMO, NiMnO3, αMn2O3, and NiMn2O4-type spinel, as well as at the octahedral and tetrahedral sites in NiMn2O4). This agrees well with our XANES results and clearly demonstrates the better thermal stability of o-LNMO than of d-LNMO. It is worthwhile to mention that the Ni−Mtet and Mn−Mtet peaks could be observed clearly only if the tetrahedral sites are occupied by heavy atoms (such as Mn and Ni but not by Li). The emergence and growth of these peaks at around 3.2 Å in Figures 8 and 9 coincide with the growth of the (220) peak in XRD, and the growth of pre-edge intensity of Mn K-edge XANES depicted in Figure 7. All of them consistently demonstrate the migration of transition-metal cations (mostly Mn2+) to the tetrahedral sites during heating. EXAFS fitting analysis was performed further to monitor the structural parameter changes during thermal decomposition in a quantitative manner. Due to the limited number of fitting variables allowed in the fitting according to the Nyquist formula,48 it was very challenging to include all the phases identified from the TR-XRD analysis such as NiMnO3, αMn2O3, and NiMn2O4, during the fitting, especially in the temperature ranges where more than two phases coexist (e.g., intermediate phases). Instead, we chose one of the model structures between spinel LiM2O4 (M = Ni or Mn) and NiMn2O4 based on the goodness of fit quality determined after the preliminary fitting session using both models for each spectrum. In general, the spinel LiM2O4 model provided better fitting results (i.e., lower EXAFS R-factor and chi square) than the NiMn2O4 model at temperatures below oxygen-release, whereas the fitting with the spinel NiMn2O4 model was better at temperatures above it. As seen in the good agreement

Figure 8. Fourier-transformed magnitude of Ni K-edge EXAFS spectra without phase correction for the fully charged d-LNMO (black) and oLNMO (red) after heating at various temperatures in comparison with the reference NiMn2O4 spinel, pristine d- and o-LNMO (bottom).

Figure 9. Fourier-transformed magnitude of Mn K-edge EXAFS spectra without phase correction for the fully charged d-LNMO (black) and o-LNMO (red) after heating at various temperatures in comparison with the reference NiMn2O4 spinel, pristine d- and oLNMO (bottom).

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Figure 10. Refined EXAFS structural parameters for the (a) Ni and (b) Mn K-edge results of the fully charged d-LNMO (black) and o-LNMO (red) after heating at various temperatures; bond lengths for (i) the first M−O shell and (ii) the second M−M shell; Debye−Waller factors for (iii) the first M−O shell and (iv) the second M−M shell, where M = Ni and Mn. Solid circle/square, fitted by using spinel Li(M2)octO4 model (M = Ni or Mn); hollow circle/square, fitted by using spinel NiMn2O4 (i.e., (Mn)tet(NiMn)octO4) model. Refined Ni2+−O bond length (2.06 Å) for the reference NiMn2O4 spinel is marked with a dashed line in the panel (a)-(i).

between the experimental data and the fitting (Supporting Information, Figures S5 and S6), this fitting approach effectively represented average local structural changes around Ni and Mn during thermal decomposition of d- and o-LNMO. Detailed fitting ranges, parameters, and constraints used in the fitting are described in the Supporting Information (Figure S5− S7 and Table S2−3). Figure 10 shows the best fitted structural parameters for the first two M−O and M−M shells (M = Ni or Mn) obtained by the EXAFS fitting including bond length (R) and Debye− Waller (D-W) factor (i.e., mean square disorder, σ2). Lattice parameter a of both as-delithiated spinel samples, calculated by using the estimated bond lengths for the second M−M bonds (M = Ni or Mn), showed very good agreement with the value calculated by TR-XRD, confirming the delithiated samples studied using TR-XRD/MS and XAS are in the same state of charge (see the Supporting Information for the detailed calculation). In the case of the d-LNMO, the average local structure around Ni was better fitted by using the original spinel LiNi2O4 model than the spinel NiMn2O4 one up to 220 °C (vice versa above 220 °C), while the o-LNMO was better described with its original spinel framework of LiNi2O4 up to 270 °C. Note that exclusive octahedral site occupancy of Ni ions was assumed in the case of the fitting with the spinel NiMn2O4 model. As the temperature increases, the average Ni−O bond lengths (∼1.90 Å, Ni4+−O) for both delithiated dand o-LNMO expand close to 2.06 Å, a value which corresponds to the Ni2+−O bond in the reference spinel NiMn2O4 (Figure 10(a)-(i)). The bond length for the second Ni−M shell (Figure 10(a)-(ii)) also increases with temperature in a similar manner, indicating that the reduction of Ni ions from 4+ to 2+ occurs during thermal decomposition of delithiated LNMO. Unlike the monotonic increase of bond lengths, the D-W factors for the Ni−O and Ni−M bonds,

which represent the degree of disorder in each bond, begin to increase above 220 °C and reach their maximum at 270 °C for d-LNMO, and at 320 °C for o-LNMO, as the temperature increases. These temperatures for the maximum D-W factors are highly correlated with the oxygen release temperatures for d- and o-LNMO (Figure 5), revealing significant structural disordering around Ni ions in LNMO during thermal decomposition due to oxygen release from the oxide lattice. This also confirms that the reduction of Ni is mostly responsible for the oxygen release during thermal decomposition. On the other hand, formation of intermediate phases (NiMnO3 and α-Mn2O3) accompanied by oxygen release could also contribute to the increase of D-W factors, as only a single representative spinel NiMn2O4 model structure was considered in the fitting. It should be noted that the reduction of Ni and corresponding local structure changes (e.g., increase of bond lengths and D-W factors for the Ni−O and Ni−M bond) due to thermal decomposition are initiated at higher temperatures by ∼50 °C in the o-LNMO in comparison with the d-LNMO case, indicating better thermal stability of o-LNMO than dLNMO. As shown in Figure 10(b), local structure changes around Mn upon heating are not as noticeable as in the Ni case, verifying a better resistance of Mn than Ni to the structural degradation during thermal decomposition. For example, the average Mn−O bond length for both d- and o-LNMO does not show appreciable expansion at 370 °C, even after major oxygen release, while the average Ni−O bond lengths for both d- and oLNMO significantly expanded by ∼0.16 Å due to the reduction of Ni4+ to Ni2+ after heating at 370 °C. However, this does not necessarily mean that the average oxidation state of Mn remained unchanged as for Mn4+ after heating at 370 °C, which is in contradiction to the Mn K-edge XANES edge shift in Figure 7. Rather, it is due to the complication of the varying 1115

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Figure 11. Schematics of the proposed two phase-transition pathways of charged LNMO during heating.

tetrahedral sites, Mn distribution in the tetrahedral sites for the sample heated at 370 °C can be estimated by setting a variable, x, in (Mnx)tet(M3−x)octO4 (M = Ni and Mn) during the Mn Kedge EXAFS fitting. The best fitted x values were 0.94 (±0.47) for d-LNMO and 0.62 (±0.37) for o-LNMO (Supporting Information, Table S2−S3). Assuming that the Ni and Mn ratio is kept to 1:3 as in the as-delithiated case of (Li0.0)Ni0.5Mn1.5O4, the estimated composition of the samples after heating at 370 °C based on the x values corresponded to [Mn(II)0.94]tet[Ni(II)0.75Mn(IV)1.31]octO4 for d-LNMO and [Mn(II)0.62]tet[Ni(II)0.75Mn(IV)1.63]octO4 for o-LNMO, respectively, suggesting more structural analogy of d-LNMO at 370 °C to the stoichiometric NiMn2O4 spinel. This also reveals more Mn migration to tetrahedral sites in the case of d-LNMO than the o-LNMO case during thermal decomposition. Note that the refined x values may not represent a real composition because there should be some contributions from other phases such as LNMO, NiMnO3, and Mn2O3, which was not considered in the fitting. Rather, it is intended to compare the degree of Mn distribution in the tetrahedral sites between d- and o-LNMO at 370 °C. All these EXAFS results consistently demonstrate better thermal stability of o-LNMO than d-LNMO and a structure-stabilizing role of Mn, which agrees well with the TRXRD/MS and XANES results. Based on the experimental results described above, the two thermal decomposition pathways are qualitatively described in the following reactions and schematically described in Figure 11:

sites of Mn in different Mn-containing phases such as LNMO, NiMnO3, α-Mn2O3, and NiMn2O4-type spinel, as well as at the octahedral and tetrahedral sites in NiMn2O4. Indeed, we only considered an averaged Mn−O bond length during the fitting, which could lead to apparently insignificant expansion in the Mn−O bond length seen in Figure 10(b)-(i). We believe that some of the Mn4+ ions (but definitely not all of them) reduce to Mn3+ (in intermediate α-Mn2O3) and Mn2+ (in spinel NiMn2O4) during thermal decomposition as evidenced by the Mn K-edge XANES result (Figure 7). On the other hand, the average Mn−M bond length for both d- and o-LNMO expanded from 2.85 to 2.91 Å after heating at 370 °C in a similar manner seen in the Ni-M bond case but, to a lesser extent, than the Ni−M bond expansion from 2.85 to 2.95 Å. The D-W factors for the Mn−O and Mn−M show a gradual and small increase compared with the Ni case with increasing temperatures, suggesting less structural disordering around Mn compared with the Ni case during thermal decomposition. Unlike the Ni case, the changes of the D-W factors for Mn do not correlate with the oxygen-release behavior during heating, which confirms that Mn is not a main contributor for the oxygen-release during thermal decomposition. Ni and Mn Kedge EXAFS fitting of both d- and o-LNMO samples heated at 370 °C using a spinel NiMn 2 O 4 model (Supporting Information, Figure S7 and Table S2−S3) showed very good fit quality, suggesting the formation of a spinel NiMn2O4 type structure as the final phase upon thermal decomposition. However, the local structure around Ni and Mn in d-LNMO heated at 370 °C was more analogous than o-LNMO to that of the stoichiometric NiMn2O4 spinel structure (Supporting Information, Figure S7 and Table S2−S3). Because the Mn ions in the NiMn2O4 spinel occupy both octahedral and 1116

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Chemistry of Materials 1 1 1 NiMnO3 + Mn2O3 + O2 ↑ 2 2 2 heating 2 1 ⎯⎯⎯⎯⎯⎯→ Mn(Ni 0.75Mn1.25)O4 + O2 3 6

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elements with strong tetrahedral occupancy (e.g., Zn2+, V5+, and Cu2+) to block the pathways of cationic migration of Ni4+ and Mn4+ in LNMO could be a viable option for the improvement in thermal stability of this class of materials.

heating

Ni 0.5Mn1.5O4 ⎯⎯⎯⎯⎯⎯→

(1)

heating

Ni 0.5Mn1.5O4 ⎯⎯⎯⎯⎯⎯→ □1 − x Mnx(Ni 0.5Mn1.5 − x)O4 − 2x + xO2 ↑ heating 2 2 ⎯⎯⎯⎯⎯⎯→ Mn(Ni 0.75Mn1.25)O4 + O2 ↑ 3 3

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CONCLUSIONS

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ASSOCIATED CONTENT

Because LiMn2O4 has good thermal stability (stable at temperature as high as 375 °C based on our experimental result), LiNi0.5Mn1.5O4 (LNMO), in which only 25% of Mn is substituted by Ni, is naturally expected to inherit this merit. Unfortunately, this is not the case. On the basis of the combined results of in situ TR-XRD/MS and in situ XAS during heating presented in this paper, it was found that both disordered d-LNMO and ordered o-LNMO at their fully charged states can start oxygen-releasing structural changes at temperatures below 300 °C during heating, and the o-LNMO showed better thermal stability than did the d-LNMO. This better thermal stability of o-LNMO is accounted for by unraveling the two thermal-decomposition-pathways of charged LNMO. Pathway one involves the formation of ilmenite NiMnO3 and α-Mn2O3 that play a bridging role in forming the final NiMn2O4-type spinel. Occurrence of this reaction in the early stage of thermal decomposition is an indicator of the poor thermal stability as this reaction involves the destruction of original structure. Readily available Ni2+ in the impurity rocksalt phase and Ni/Mn segregation resulted from disordering in d-LNMO obviously facilitate such reaction. In contrast, oLNMO does not have the favorable environment for the pathway one to take place and adopts different pathway to begin with. For pathway two, Mn migrates to the tetrahedral sites within the original spinel framework to form NiMn2O4type spinel without formation of intermediate phases. This reaction happens at higher temperature than the reaction one and releases oxygen in a milder way (e.g., begun higher temperatures and spread in wider temperature ranges). Although many researchers believe that disordered LNMO materials have better electrochemical performance (e.g., cycle life and rate capability) than ordered LNMO, improving their thermal stability should be seriously considered. On the other hand, considering the newly reported good electrochemical performance6,20 and thermodynamic stability shown in this study for ordered LNMO, further extensive research on the ordered LNMO will be quite rewarding.

(2)

For the first pathway described in reaction 1, the delithiated LNMO first decomposes into two intermediate compounds, namely ilmenite NiMnO3 and α-Mn2O3; these two compounds further react to form the NiMn2O4-type spinel. On the basis of the proposed reaction 1, the first step (formation of NiMnO3 and α-Mn2O3) is the main contributor to the release of oxygen, being 3-fold larger than the second step (formation of the NiMn2O4-type spinel). The second pathway described in reaction 2 involves continuous Mn migration from octahedral sites to tetrahedral sites in the original spinel framework. As the structure framework is maintained throughout the process, this pathway described in reaction 2 features milder oxygen release spread in a wider temperature range. The NiMn2O4-type spinel formed at low temperatures is likely more defective than that formed at higher temperatures. From our previous discussion, it is clear that, in case of d-LNMO, the first pathway (reaction 1) is dominant, which causes the oxygen release process and formation of NiMn2O4-type spinel at lower temperatures than the o-LNMO case. For o-LNMO, the second pathway (reaction 2) is the major reaction to begin with and the first pathway (reaction 1) only kicks in in the later stage of the heating process, which triggered more rapid formation of NiMn2O4type spinel phase. These reaction pathways described here are simplified models, without considering the cation deficiency in each phase. The Mn(Ni0.75Mn1.25)O4 (i.e., Mn(II)tet(Ni(II)0.75Mn(IV)1.25)octO4) formula is used to show that tetrahedral sites are mainly occupied by Mn cations whereas the original Ni:Mn = 1:3 ratio is maintained in final structure. On the basis of the combined results of in situ TR-XRD/MS and in situ XAS during heating, we propose two possible mechanisms contributing to the better thermal stability of charged o-LNMO than charged d-LNMO. The first one is based on the restriction of the migration of transition metal cations during heating: the ordering of Ni and Mn in the transition-metal layer may increase the energy barrier for the movement of these cations and, therefore, push the formation of NiMn2O4-type spinel to higher temperatures. The second one is based on the restriction of the formation of NiMnO3 and α-Mn2O3 phases: the amount of electrochemically inactive rock-salt phase, providing the readily available Ni2+ in favor of NiMnO3 formation, is larger in the charged d-LNMO than in the charged o-LNMO. This NiMnO3 phase, along with the αMn2O3 phase, acting as intermediate phases for charged dLNMO, bridge the phase transformation from the initial charged d-LNMO to the final NiMn2O4-type spinel. In addition, the locally Ni2+-rich and Mn4+-rich regions in the dLNMO which were likely formed during thermal decomposition (XANES result in Figures 6 and 7) are also in favor of the formation of the NiMnO3 and α-Mn2O3 phases. These two proposed mechanisms suggest that increasing the degree of ordering and reducing the amount of the rock-salt impurity phase can be considered as effective approaches for improving thermal stability of spinel LiMxMn2−xO4-type (M: transition metals) cathode materials. In addition, doping of other

S Supporting Information *

In situ XANES spectra measured at 370 °C and at room temperature after cooling down, electrochemical profile and Rietveld refinement result for d- and o-LNMO; evolution of (110) peak for o-LNMO during heating; in situ XRD/MS for charged LixMn2O4; detailed description of EXAFS fitting analysis for d- and o-LNMO with heating temperatures. This material is available free of charge via the Internet at http:// pubs.acs.org.

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AUTHOR INFORMATION

Corresponding Authors

*K. N. E-mail: [email protected]. *X. Y. E-mail: [email protected]. Notes

The authors declare no competing financial interest. 1117

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(27) Patoux, S.; Sannier, L.; Lignier, H.; Reynier, Y.; Bourbon, C.; Jouanneau, S.; Le Cras, F.; Martinet, S. Electrochim. Acta 2008, 53, 4137. (28) Xiang, H. F.; Wang, H.; Chen, C. H.; Ge, X. W.; Guo, S.; Sun, J. H.; Hu, W. Q. J. Power Sources 2009, 191, 575. (29) Zhou, F.; Zhao, X. M.; van Bommel, A.; Xia, X.; Dahn, J. R. J. Electrochem. Soc. 2011, 158, A187. (30) Giorgetti, M. ISRN Mater. Sci. 2013, No. ID 938625. (31) Armstrong, A. R.; Holzapfel, M.; Novak, P.; Johnson, C. S.; Kang, S. H.; Thackeray, M. M.; Bruce, P. G. J. Am. Chem. Soc. 2006, 128, 8694. (32) Ravel, B.; Newville, M. J. Synchrotron Radiat. 2005, 12, 537. (33) LarsonA. C.Von DreeleR. B.Los Alamos National Laboratory Report LAUR, 1985. (34) Rietveld, H. M. J. Appl. Crystallogr. 1969, 2, 65. (35) Toby, B. H. J. Appl. Crystallogr. 2001, 34, 210. (36) Cabana, J.; Casas-Cabanas, M.; Omenya, F. O.; Chernova, N. A.; Zeng, D. L.; Whittingham, M. S.; Grey, C. P. Chem. Mater. 2012, 24, 2952. (37) Lee, E. S.; Nam, K. W.; Hu, E. Y.; Manthiram, A. Chem. Mater. 2012, 24, 3610. (38) Wu, L. J.; Nam, K. W.; Wang, X. J.; Zhou, Y. N.; Zheng, J. C.; Yang, X. Q.; Zhu, Y. M. Chem. Mater. 2011, 23, 3953. (39) Lebail, A.; Duroy, H.; Fourquet, J. L. Mater. Res. Bull. 1988, 23, 447. (40) Cloud, W. H. Phys. Rev. 1958, 111, 1046. (41) Tang, X. X.; Manthiram, A.; Goodenough, J. B. J. Less-Common Met. 1989, 156, 357. (42) Yabuuchi, N.; Lu, Y. C.; Mansour, A. N.; Chen, S.; Shao-Horn, Y. J. Electrochem. Soc. 2011, 158, A192. (43) Kovanda, F.; Grygar, T.; Dornicak, V. Solid State Sci. 2003, 5, 1019. (44) Deng, H. X.; Belharouak, I.; Sun, Y. K.; Amine, K. J. Mater. Chem. 2009, 19, 4510. (45) Geller, S. Acta Crystallogr., Sect. B: Struct. Crystallogr. Cryst. Chem. 1971, B 27, 821. (46) Wells, A. F. Structural inorganic chemistry, Oxford University Press: New York, 1984. (47) Yamamoto, T. X-Ray Spectrom. 2008, 37, 572. (48) Stern, E. A. Phys. Rev. B: Condens. Matter Mater. Phys. 1993, 48, 9825.

ACKNOWLEDGMENTS This work was supported by the U.S. Department of Energy, the Assistant Secretary for Energy Efficiency and Renewable Energy, Office of Vehicle Technologies under Contract Number DE-AC02-98CH10886. Use of the NSLS was supported by the U.S. Department of Energy, Office of Science, Office of Basic Energy Sciences, under Contract No. DE-AC02-98CH10886. The authors acknowledge technical support by the NSLS’s beamline scientists Dr. Sanjaya Senanayake at X7B and Drs. Nebojsa Marinkovic and Syed Khalid at X18A.

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REFERENCES

(1) Thackeray, M. M.; David, W. I. F.; Bruce, P. G.; Goodenough, J. B. Mater. Res. Bull. 1983, 18, 461. (2) Thackeray, M. M.; Johnson, P. J.; Depicciotto, L. A.; Bruce, P. G.; Goodenough, J. B. Mater. Res. Bull. 1984, 19, 179. (3) Zhong, Q. M.; Bonakdarpour, A.; Zhang, M. J.; Gao, Y.; Dahn, J. R. J. Electrochem. Soc. 1997, 144, 205. (4) Kim, J. H.; Myung, S. T.; Yoon, C. S.; Kang, S. G.; Sun, Y. K. Chem. Mater. 2004, 16, 906. (5) Kunduraci, M.; Amatucci, G. G. J. Electrochem. Soc. 2006, 153, A1345. (6) Ma, X. H.; Kang, B.; Ceder, G. J. Electrochem. Soc. 2010, 157, A925. (7) Shaju, K. M.; Bruce, P. G. Dalton Trans. 2008, 40, 5471. (8) Kim, J. H.; Myung, S. T.; Sun, Y. K. Electrochim. Acta 2004, 49, 219. (9) Kunduraci, M.; Al-Sharab, J. F.; Amatucci, G. G. Chem. Mater. 2006, 18, 3585. (10) Maeda, Y.; Ariyoshi, K.; Kawai, T.; Sekiya, T.; Ohzuku, T. J. Ceram. Soc. Jpn. 2009, 117, 1216. (11) Song, J.; Shin, D. W.; Lu, Y. H.; Amos, C. D.; Manthiram, A.; Goodenough, J. B. Chem. Mater. 2012, 24, 3101. (12) Ariyoshi, K.; Maeda, Y.; Kawai, T.; Ohzuku, T. J. Electrochem. Soc. 2011, 158, A281. (13) Cabana, J.; Zheng, H. H.; Shukla, A. K.; Kim, C.; Battaglia, V. S.; Kunduraci, M. J. Electrochem. Soc. 2011, 158, A997. (14) Chemelewski, K. R.; Lee, E. S.; Li, W.; Manthiram, A. Chem. Mater. 2013, 25, 2890. (15) Ivanova, S.; Zhecheva, E.; Stoyanova, R.; Nihtianova, D.; Wegner, S.; Tzvetkova, P.; Simova, S. J. Phys. Chem. C 2011, 115, 25170. (16) Liu, J.; Manthiram, A. Chem. Mater. 2009, 21, 1695. (17) Xiao, J.; Chen, X. L.; Sushko, P. V.; Sushko, M. L.; Kovarik, L.; Feng, J. J.; Deng, Z. Q.; Zheng, J. M.; Graff, G. L.; Nie, Z. M.; Choi, D. W.; Liu, J.; Zhang, J. G.; Whittingham, M. S. Adv. Mater. 2012, 24, 2109. (18) Ariyoshi, K.; Iwakoshi, Y.; Nakayama, N.; Ohzuku, T. J. Electrochem. Soc. 2004, 151, A296. (19) Kim, J. H.; Yoon, C. S.; Myung, S. T.; Prakash, J.; Sun, Y. K. Electrochem. Solid-State Lett. 2004, 7, A216. (20) Ariyoshi, K.; Yamatoa, R.; Makimura, Y.; Amazutsumi, T.; Maeda, Y.; Ohzuku, T. Electrochemistry 2008, 76, 46. (21) Schilling, O.; Dahn, J. R. J. Electrochem. Soc. 1998, 145, 569. (22) Wang, L.; Maxisch, T.; Ceder, G. Chem. Mater. 2007, 19, 543. (23) Bak, S. M.; Nam, K. W.; Chang, W.; Yu, X. Q.; Hu, E. Y.; Hwang, S.; Stach, E. A.; Kim, K. B.; Chung, K. Y.; Yang, X. Q. Chem. Mater. 2013, 25, 337. (24) Konishi, H.; Yuasa, T.; Yoshikawa, M. J. Power Sources 2011, 196, 6884. (25) Nam, K. W.; Bak, S. M.; Hu, E. Y.; Yu, X. Q.; Zhou, Y. N.; Wang, X. J.; Wu, L. J.; Zhu, Y. M.; Chung, K. Y.; Yang, X. Q. Adv. Func. Mater. 2013, 23, 1047. (26) Bhaskar, A.; Gruner, W.; Mikhailova, D.; Ehrenberg, H. RSC Advances 2013, 3, 5909. 1118

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