Ta2O5 Coating as an HF Barrier for Improving the Electrochemical

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Ta2O5 Coating as an HF Barrier for Improving the Electrochemical Cycling Performance of High-Voltage Spinel LiNi0.5Mn1.5O4 at Elevated Temperatures Liubin Ben, Hailong Yu, Yida Wu, Bin Chen, Wenwu Zhao, and Xuejie Huang ACS Appl. Energy Mater., Just Accepted Manuscript • Publication Date (Web): 21 Sep 2018 Downloaded from http://pubs.acs.org on September 21, 2018

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ACS Applied Energy Materials

Ta2O5 Coating as an HF Barrier for Improving the Electrochemical Cycling Performance of High-Voltage Spinel LiNi0.5Mn1.5O4 at Elevated Temperatures Liubin Ben†, ‡, §, Hailong Yu†, ‡, §, Yida Wu†, ‡, Bin Chen†, ‡, Wenwu Zhao†, ‡ and Xuejie Huang†, ‡,*

† Beijing National Laboratory for Condensed Matter Physics, Institute of Physics, Chinese Academy of Sciences, Beijing 100190, China ‡ University of Chinese Academy of Sciences, Beijing, 100190, China § These authors contributed equally to this work.

* Email: [email protected] (XJ.H)

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Abstract The high-voltage spinel LiNi0.5Mn1.5O4 cathode material suffers from the rapid degradation of electrochemical cycling performance at elevated temperatures, which prevents its successful commercialization. Herein, we show that coating the surface of this material with Ta2O5, which has high resistance against hydrofluoric acid (HF) attack, is an effective way to improve its electrochemical cycling performance. A Ta2O5-coated LiNi0.5Mn1.5O4 half-cell shows a capacity retention of ~93% and a coulombic efficiency of ~98% after 100 cycles at 55 °C, compared to the corresponding values of ~76% and ~95% measured for the bare LiNi0.5Mn1.5O4 half-cell. The detailed structural analysis of the Ta2O5-coated LiNi0.5Mn1.5O4 shows that a small amount of Ta5+ ions diffuse into the 16c site on the cathode surface during the coating process, as directly observed by Cs corrected scanning transmission electron microscopy. The modification of the LiNi0.5Mn1.5O4 surface with Ta5+, together with the residual Ta2O5 coating, stabilizes the surface structure during cycling, leading to reduced Ni and Mn dissolution as well as formation of the solid electrolyte interface (SEI). In contrast, LiNi0.5Mn1.5O4 coated with HF scavengers, such as Al2O3, shows only limited improvement in cycling performance after prolonged cycling at 55 °C, due to the consumption of the surface coating by reaction with HF, which leaves LiNi0.5Mn1.5MnO4 unprotected against HF attack.

Keywords: Lithium ion batteries, LiNi0.5Mn1.5O4, Coating, Ta2O5, Scanning Transmission Electron Microscopy (STEM)


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Introduction Lithium ion batteries find wide application in portable electronic devices, electric vehicles, and hybrid electric vehicles due to their high energy density and environmental friendliness.1-2 The ever-increasing capacity requirements of these applications result in a strong demand for cathode and anode materials with high specific capacities.3-8 A possible way to achieve a high specific capacity is to increase the operating voltage of the cathode materials, e.g., to ≥ 4.5 V for LiCoO29-10 and lithium-rich layered materials with LiMO2-Li2MnO3 (M = Ni, Mn, etc.) general formula,11-13 ~4.7 V for LiNi0.5Mn1.5O4 (LNMO),14-15 and ~5 V for LiMPO4 (M = Ni, Co, Mn).16-18 However, cathode materials cycled to high voltage exhibit rapid performance degradation, particularly at elevated temperatures, which represents the main issue hindering their commercial application. The degradation of the cycling performance is generally attributed to the structural distortion of the cathode materials and to parasitic reactions between cathode and electrolyte due to the presence of HF.19-22

A common approach to improve the cycling performance at elevated temperatures involves coating the surface of cathode materials with various metal oxides,23-31 phosphates,32-34, solid electrolytes35-37 and other materials.38-39 The surface region in contact with the electrolyte, where lithium ions are reversibly intercalated/deintercalated, is considered the key region involved in the degradation process. Coating the surface of cathode materials with metal oxides is a common approach, due to the facile preparation methods involved. For example, Al2O3 coatings were reported to improve the electrochemical cycling performance of layered LiCoO2,23,25,31,40-41 spinel LiMn2O4,42 LiFePO4,43 etc. The improved cycling performance has generally been attributed to the consumption of trace amounts of HF present in the electrolyte via reactions between the coating metal oxide and HF, leading to stabilization of the surface structure and limiting the reactions between cathode material and electrolyte. However, most metal oxides have achieved only marginal improvements in the cycling performance of LNMO at a high operating voltage of ~4.7 V, which prevented its commercial applications. ACS Paragon Plus Environment


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In general, the metal oxides used as coatings can be classified into two types: HF scavengers and HF barriers. The former are constantly consumed by sacrificial reactions with trace HF to form oxide fluorides during cycling, whereas the latter may not react with HF but cover the surface of the cathode particles and retain their coverage and integrity.44 Recent computer simulations have suggested that HF barriers may be more effective than HF scavengers.44 Surprisingly, no detailed experimental studies on the resistance of metal oxides against HF attack have been reported to date. Furthermore, Ta2O5 coated high-voltage spinel LNMO cathode materials have not been studied in details, though a few Ta2O5 coated layered cathode materials have been reported.45-46 Thus, in this work, we initially examined the stability of several metal oxides (i.e., Al2O3, TiO2, Nb2O5, and Ta2O5) against HF attack, by immersing the oxide powders in dilute HF solution (20% v/v) for 7 days. The results showed that most metal oxides, particularly Al2O3, dissolved completely in the HF solution due to reaction with HF, suggesting that they can be classified as HF scavengers. However, Ta2O5 showed almost no dissolution after immersion in the HF solution; this was further confirmed by structural analysis, which revealed that Ta2O5 can be classified an HF barrier. The electrochemical cycling performance of Ta2O5 (HF barrier)coated spinel LNMO was investigated and compared to that of Al2O3 (HF scavenger)-coated LNMO. The effects of these two metal oxides on the cycling performances of the coated materials are discussed in detail. The results suggested that many oxides worked well with the normal-voltage cathode materials, showing significant improvement of cycling performance after coating, but exhibited very limited improvement for the high-voltage spinel LNMO after similar coating treatment. Some further considerations on the selection of suitable metal oxides for coating high-voltage cathode materials are given.

Experimental Section Sample Preparation. Bare LNMO powders were prepared by the conventional solid-state reaction method. Stoichiometric amounts of Li2CO3 (Alfa Aesar, 99.9%) and Ni0.25Mn0.75(OH)2 (Henan Kelong) ACS Paragon Plus Environment

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precursors were mixed thoroughly, and the resulting mixture was preheated at 500 °C for 5 h, followed by sintering at 900 °C for 12 h. Both above procedures were carried out in air. To prepare the Ta2O5and Al2O3-coated LNMO (Ta-LNMO and Al-LNMO, respectively) powders, LNMO and 2 wt% Ta2O5 (Alfa Aesar, 99.5%) or 2 wt% Al2O3 (Alfa Aesar, 99.5%) were mixed in stoichiometric ratios and sintered at 900 °C for 6 h in air. Stability of Metal Oxides in HF. The stability of the metal oxides against HF attack was tested by immersion of 5 g of nanosized oxides in 20 mL of dilute HF (20% v/v) for 7 days at room temperature (RT). Powder XRD Diffraction. A Bruker D8 Advance X-ray diffractometer with a Cu Kα radiation source (λ = 1.54056 Å) was employed to determine the crystal structure and phase composition of the samples. The diffractometer was equipped with a LynxEye detector and operated at 40 kV and 40 mA. The samples were scanned in the 2θ range of 10–100° at a scan rate of 0.02°/0.1s. All XRD patterns were refined using the Rietveld method, as implemented in the program TOPAS.47 Scanning Electron Microscopy (SEM) Characterization. The morphology and microstructure of the samples were characterized by SEM using a Hitachi SU-4800 instrument. Transmission Electron Microscopy (TEM) Characterization. TEM characterization was performed using a Philips CM200 electron microscope with a field emission gun operated at an acceleration voltage of 200 kV. Scanning Transmission Electron Microscopy (STEM) Characterization. STEM characterization was performed using a JEM-ARM 200F microscope operated at 200 kV. The attainable spatial resolution of the microscope was 80 pm, at the incident semi-angle of 25 mrad. X-ray Photoelectron Spectroscopy (XPS) Characterization. XPS measurements were performed using a spectrometer with Mg Kα radiation (ESCALAB 250, Sigma Probe, Thermo VG Scientific Co., Ltd.). All reported binding energies were corrected using the signal for carbon at 284.8 eV as an internal standard. The fitting of the peaks were performed with the CasaXPS software. The background was corrected using the Shirley method. ACS Paragon Plus Environment


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Metal Dissolution. The amount of dissolved metal ions was investigated by soaking the fully charged electrode in the electrolyte (10 mL) at 55 °C for 7 days, followed by determining the amount of metal ions in the electrolyte by inductively coupled plasma (ICP). Before the ICP measurements, the electrolyte was filtered with filter paper (pore size 0.2 µm).

Results Resistance of Metal Oxides to HF Attack The resistance to HF attack of several metal oxides, i.e., Al2O3, TiO2, Nb2O5, and Ta2O5, was initially studied by immersion of the corresponding oxide powders (5 g, with an average particle size of ~100200 nm) in dilute HF solution (20 mL, with a concentration of 20% v/v) for 7 days at RT, and the results are shown in Figure 1a and 1b. As expected, the Al2O3 powders rapidly reacted with the dilute HF solution, resulting in complete dissolution after immersion for 1 h, in agreement with previous reports.25,41 The TiO2 powders reacted with the HF solution at a somewhat slower rate, but still completely dissolved after immersion for 7 days. The Nb2O5 powders also reacted to some extent with the dilute HF solution, resulting in partial dissolution after immersion for 7 days. Surprisingly, Ta2O5, whose structure is similar to that of Nb2O5, showed much stronger resistance to HF attack. After 7 days of immersion in the dilute HF solution (Figure 1b), the Ta2O5 oxide powders showed almost no evidence of dissolution, suggesting that Ta2O5 may have the highest resistance against attack by HF among the oxides investigated in this work.

The Ta2O5 oxide powders were removed from the dilute HF solution after 7 days of immersion, for the subsequent XRD phase analysis. The amount of Ta2O5 powders after drying was ~5g, similar to that before the immersion test. All XRD peaks in Figure 1c were associated with pure Ta2O5, and no clear peaks associated with tantalum fluoride were observed. These results suggest that Ta2O5 is a good HF ACS Paragon Plus Environment

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barrier. Therefore, Ta-LNMO was prepared via solid-state reaction to investigate the effects of the Ta2O5 coating on the electrochemical cycling performance. Al2O3, which is a good HF scavenger, was also used to coat LNMO (forming Al-LNMO) for comparison.

Figure 1. Immersion of Al2O3, TiO2, Nb2O5, and Ta2O5 nanosized powders in dilute HF solution for (a) 1 h and (b) 7 days at RT. (c) XRD patterns of pristine Ta2O5 powders and of Ta2O5 powders after immersion for 7 days in dilute HF solution. The black and red dashed lines in panels (a) and (b) indicate the residual oxides in the dilute HF solution. Reference Ta2O5 (PDF 71-0639)

Morphology and Structure of the bare LNMO and Ta-LNMO The morphologies of the bare LNMO (Figure 2a1) and Ta-LNMO (Figure 2b1) powders shown in the SEM images were generally similar, with an average size of ~1-2 µm for the primary particles. The SEM images show that the surface of the Ta-LNMO powders was covered with many small particles associated with Ta2O5, while the bare LNMO powders exhibited a smooth surface. TEM image further suggests that the surface of Ta-LNMO was covered with Ta2O5, similar to that of Al-LNMO covered with Al2O3 (Supporting Information Figure S1). The crystal structures of the bare LNMO and TaLNMO powders (Figures 2a2 and 2b2, respectively) were similar, with a Rietveld-refined lattice parameter (a) of 8.1671(2) Å for the former, compared to 8.1651(3) Å for the latter. The refined results also confirm the presence of ~3.36 wt% rocksalt LixNiyO in the Ta-LNMO powders (Figure 2b2), ACS Paragon Plus Environment


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compared to ~1.65 wt% for the bare LNMO powders (Figure 2a2).48 It should be noted that a higher amount of rocksalt LixNiyO was observed in Ta-LNMO compared to that of previously reported TiO2coated LNMO, possibly due to the higher oxidation state of Ta5+ than Ti4+,30 in good agreement with first-principles computer simulation studies.49 The XPS investigation of the surface electronic structure, e.g., the oxidation state of Mn, provides important information on the possible diffusion of Ta5+ into the spinel LNMO structure during the high-temperature processing. The fitted Mn 2p XPS results for the bare and Ta-coated LNMO powders are shown in Figures 2a3 and 2b3, respectively. The detailed fitting procedures can be found elsewhere.50-51 These results confirmed that the Mn4+ content in the surface region of the Ta-LNMO powders was only ~59.68%, whereas that of Mn3+ was ~40.32%, compared to ~74.49% and ~25.51% Mn4+ and Mn3+ contents for the bare LNMO powders, respectively. This indicated that a small amount of Ta5+ diffused into the structure of LNMO, particularly in the surface region, during the hightemperature processing, and the excess charge was compensated by the reduction of Mn4+ to Mn3+ or lower. This is similar to the process occurring in many other metal oxide-coated cathode materials.30,42,49 The diffusion of Ta5+ into the Ni/Mn sites is also responsible for the formation of the rocksalt structure, as highlighted by the XRD results. The migration sites of Ta5+ are difficult to determine based on XPS fitted results observed here, due to the higher oxidation state of Ta5+ compared to Ni2+ and Mn4+. Migration of Ta5+ into the spinel structure to replace Ni2+ or Mn4+ may introduce additional electrons in the structure, resulting in the reduction of Mn4+ ions, as observed by XPS. However, since the amount of rocksalt impurities increased after Ta2O5 coating as observed by XRD above, it is likely that Ta5+ migrated into the Ni2+ site. Furthermore, recent computer simulations have suggested that Ta5+ ions prefer to migrate to the Mn site in the LNMO, since the energy required for migration to this site is 0.2 eV lower than that for the Ni site.49


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Figure 2. Microstructure, crystal structure, and surface electronic structure of bare LNMO and TaLNMO powders. SEM images of (a1) bare LNMO and (b1) Ta-LNMO powders. XRD pattern and Rietveld refinement of (a2) bare LNMO and (b2) Ta-LNMO powders. Fitting of Mn 2p XPS peaks for (a3) bare LNMO and (b3) Ta-LNMO powders. The wRp, Rp and gof values are 9.182%, 6.564%, and 1.306 for LNMO and 9.182%, 6.564%, and 1.306 for Ta-LNMO, respectively.

Atomic-Scale Structure of the bare LNMO and Ta-LNMO STEM measurements were performed to inspect in detail the crystal structure of the bare LNMO and Ta-LNMO surface, and the results are shown in Figure 3 and 4, respectively. The typical STEM-highangle annular dark-field (STEM-HAADF) image of the bare LNMO powders (Figure 3c1) shows the arrangement of Ni and Mn atomic columns associated with the spinel structure, viewed along the perpendicular direction to the {110} crystallographic planes, as also indicated in the schematic structural models of Figure 3a and 3b. The assignment of the contrast observed in the STEM-HAAD image to the spinel crystal structure can be found in the literature.15,30,42,51-55 The detailed analysis of the enlarged surface (Figure 3c2) and subsurface regions (Figure 3c3) reveal the arrangement of contrast associated with Ni/Mn atomic columns in a typical diamond configuration, which is similar to the standard spinel structure and previous reported works.15,30,55



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Figure 3. STEM images of bare LNMO. Schematic illustrations of (a) {110} planes of the spinel structure and (b) atomic arrangement of the spinel structure viewed perpendicular to the {110} planes, showing the Ni(Mn) diamond configuration. (c1) Surface and subsurface regions of bare LNMO viewed perpendicular to the {110} planes, and enlargements of (c2) surface and (c3) subsurface regions.

The typical STEM HAADF image of the Ta-LNMO powders (Figure 4a1) also shows the arrangement of Ni and Mn atomic columns associated with the spinel structure, viewed along the perpendicular direction to the {110} crystallographic planes. The detailed analysis of the surface and subsurface regions, shown in Figures 4a2 and 4a3, reveals the presence of extra contrast in the center of the Ni(Mn) diamond (16c site), which is stronger in the surface region. This is significantly different from the image of bare LNMO (Figure 3), which shows no contrast in this site.15,30,55 The occupation of the 16c site suggests a transformation from the spinel to the rocksalt structure induced by migration of Ta5+ into the structure,

49

as shown in Figure 4b, in good agreement with the XRD and XPS results discussed

above (Figure 2). The distortion of the surface structure was also confirmed by the fast Fourier transform (FFT) of the image, which shows additional {0 0 2} reflections (indicated by red circles) associated with the migration of transition metal (TM) ions, Figure 4c.55-56 The STEM-energydispersive spectroscopy (STEM-EDS) analysis of the surface elements, shown in Figure 4d, confirms that the surface was rich in Ta and Ni but poor in Mn and O, indicating that Ta5+ ions may diffuse into the structure during the coating process. ACS Paragon Plus Environment

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Figure 4. STEM images of Ta-LNMO. (a1) Surface and subsurface regions of Ta-LNMO viewed perpendicular to the {110} planes, and enlargements of (a2) surface and (a3) subsurface regions. (b) Schematic illustration of the atomic arrangement of the spinel structure viewed perpendicular to the {110} planes, with additional atoms at the center of the Ni(Mn) diamond. (c) FFT of the STEMHAADF image showing additional {0 0 2} reflections should be absent from the spinel structure, indicated by red circles. (d) STEM-EDS profiles of O, Ta, Mn, and Ni elements along the line in panel (a1).

Cycling Performance of the bare LNMO and Ta-LNMO half-cells at RT The electrochemical cycling performance of the bare LNMO and Ta-LNMO half-cells was initially investigated at room temperature, and the results are shown in Figure 5. Both half-cells showed excellent capacity retention and coulombic efficiency at room temperature, in agreement with a previous work.15 The plateaus at ~4.7 and ~4.0 V (Figures 5a and 5b) are associated with the Ni2+/Ni4+ and Mn3+/Mn4+ redox couples, respectively.57-58 The ~4 V plateau and the associated discharge capacity were more pronounced for the Ta-LNMO half-cell (15.8 mAh/g) than for the bare LNMO (5.6 mAh/g) and Al-LNMO (8.9 mAh/g, Supporting Information, Figure S2a) half-cells, due to diffusion of highvalence Ta5+ ions into the structure during the synthesis, resulting in the reduction of Mn4+ to Mn3+ for charge compensation. The initial discharge capacity of the Ta-LNMO half-cell was 129.5 mAh/g, slightly lower than that (132.6 mAh/g) of the bare LNMO half-cell, but similar to that of the Al-LNMO half-cell (128.3 mAh/g, Supporting Information, Figure S2a). ACS Paragon Plus Environment


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The room-temperature capacity retention and coulombic efficiency of the bare and Ta2O5-coated LNMO half-cells are shown in Figures 5c and 5d. At room temperature, the discharge capacity retention of the bare LNMO and the Ta-LNMO half-cells after 100 cycles (0.2 C) were ~94% and ~97%, respectively (Figure 5c). In the first cycle, the coulombic efficiencies of the bare LNMO and TaLNMO half-cells were ~93% and ~94%, respectively, and increased to ~98-99% after few initial cycles (Figure 5d). The Al-LNMO half-cell also showed good cycling performance at room temperature (Supporting Information, Figure S2b). This suggests that the electrolyte is relatively stable during cycling at room temperature.15,59 In addition, the bare LNMO and Ta-LNMO half-cells exhibit similar rate capacities (Supporting Information, Figure S3). The Li-ion diffusion coefficients (DLi+) calculated from the GITT data during both charge and discharge processes of LNMO and Ta-LNMO half-cells were also very similar (Supporting Information, Figure S4)

Figure 5. Electrochemical cycling performance of bare LNMO and Ta-LNMO half-cells at room temperature. Charge-discharge curves for the 1st, 50th, and 100th cycles of (a) bare LNMO and (b) TaLNMO half-cells. (c) Capacity retention of bare LNMO and Ta-LNMO half-cells. (d) Coulombic efficiency of bare LNMO and Ta-LNMO half-cells.

Cycling Performance of the bare LNMO and Ta-LNMO half-cells at 55 °C ACS Paragon Plus Environment

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In the high-temperature cycling tests, the bare LNMO and Ta-LNMO half-cells were initially cycled at room temperature for five cycles (formation cycles) at 0.1 C and then transferred to a 55 °C oven for the subsequent cycles. The results (Figure 6) show that the Ta-LNMO half-cell exhibited a much superior capacity retention and coulombic efficiency than the bare LNMO one. The initial discharge capacity of the latter was 133.4 mAh/g, and decreased to 101.8 mAh/g after 100 cycles at 55 °C, with a capacity retention of only ~76% (Figures 6a and 6c). In contrast, the discharge capacity of the TaLNMO half-cell was 131.5 mAh/g, and decreased to 122.3 mAh/g after 100 cycles, indicating a capacity retention of ~93% after 100 cycles (Figures 6b and 6c). The average coulombic efficiency of the Ta-LNMO half-cell (~98%) was also higher than that of the bare (~95%) LNMO half-cell at 55 °C (Figure 6d). In contrast, the Al-LNMO half-cell only showed a limited improvement of the cycling performance at 55 °C, with a capacity retention of only ~81% after the 100th cycle (Supporting Information, Figure S2c). Note that the coulombic efficiency of the Ta-LNMO at 55 oC cycling was still not 100%, which is attributed the fact that the surface of the LNMO was not fully covered with Ta2O5 and the charged Ta-LNMO still oxidized the electrolyte during cycling. Further optimization of the coating procedure is under investigation. Nevertheless, the Ta-LNMO material also showed better cycling performances at elevated temperatures compared to the previously reported Ti-LNMO.30

Figure 6. Electrochemical cycling performance of the bare LNMO and Ta-LNMO half-cells at 55 °C. Charge-discharge curves for the 1st, 50th, and 100th cycles of (a) bare LNMO and (b) Ta-LNMO halfACS Paragon Plus Environment


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cells. (c) Capacity retention of bare LNMO and Ta-LNMO half-cells. (d) Coulombic efficiency of bare LNMO and Ta-LNMO half-cells. Mn3+/Mn4+ Redox Couple during Cycling at RT and 55 °C A small amount of oxygen was lost by the spinel LNMO cathode material during cycling at high voltage; although this amount was too small to be measured directly, it could be estimated from the capacity associated with the Mn3+/Mn4+ redox couple.60-61 During cycling at RT, both the bare LNMO and Ta-LNMO half-cells showed limited increases in capacity associated with the Mn3+/Mn4+ redox couple: 1.7 mAh/g (from 5.6 to 7.3 mAh/g, as shown in Figure 7a1) for the former and 0.2 mAh/g (from 15.8 to 16.0 mAh/g, see Figure 7b1) for the latter. However, during cycling at 55 °C, the bare LNMO half-cell showed a sharp increase in the capacity associated with the Mn3+/Mn4+ redox couple, particularly in the initial cycles. The capacity increased from 6.3 to 8.6 mAh/g in the initial ~20 cycles, then gradually increased to 9.8 mAh/g after 80 cycles (Figure 7a2), for a total increase of 3.5 mAh/g. In contrast, the Mn3+/Mn4+-associated capacity remained almost constant for the Ta-LNMO half-cell (Figure 7b2), with a total increase of 0.9 mAh/g (from 15.9 to 16.8 mAh/g). The change of the capacities associated with the Mn3+/Mn4+ redox couples during cycling at 55 oC could also be observed from the dQ/dV curves as shown in the Supporting Information Figure S5. The Al-LNMO half-cell also showed a steady increase in the capacities associated with the Mn3+/Mn4+ redox couple during cycling at room temperature, along with a sharp increase after a few cycles at 55 °C (Supporting Information, Figure S2d). The changes in the capacities associated with the Mn3+/Mn4+ redox couple for Ta-LNMO suggest that structural distortions and the associated oxygen evolution are effectively reduced during cycling at elevated temperatures.


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Figure 7. Discharge capacity associated with the Mn3+/Mn4+ redox couple during cycling of the bare LNMO half-cell at (a1) room temperature and (a2) 55 °C, and during cycling of the Ta-LNMO half-cell at (b1) room temperature and (b2) 55 °C.

XPS Analysis of Solid Electrolyte Interface (SEI) after Cycling Detailed information on the electrolyte decomposition and formation of the SEI, particularly at 55 °C, was obtained by XPS analysis of the electrodes after cycling. The XPS results for the surface components of the cycled LNMO and Ta-LNMO electrodes (after 20 cycles at 55 °C) are shown in Figure 8. All XPS spectra were adjusted and aligned with respect to the binding energy and intensity of the C 1s signal originating from carbon black (C-C), located at ~284.8 eV (Figures 8a1-8d1).

In general, both the LNMO and Ta-LNMO electrodes showed an increase in C, P, O, and F peaks during cycling at 55 °C, indicating that the electrodes were covered with a gradually thickening surface film. The C 1s peaks at ~285.7, ~289.5, ~286.5, and ~290.8 eV (Figures 8a1-8d1) indicate the surface presence of C in the SEI species (C-H, C-O, C=O, and CO32-, respectively).62-65 A comparison of the bare LNMO and Ta-LNMO pristine electrodes (Figures 8a1 and 8c1) showed that the intensity of the carbonate peaks associated with the SEI species increased upon cycling (Figures 8b1 and 8d1), suggesting the deposition of SEI species on the surface of the cycled LNMO electrode at elevated



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temperatures. After cycling, higher amounts of SEI species were observed on the surface of the bare LNMO electrode (Figure 8b1) than on that of the Ta-LNMO electrode (Figure 8d1).

No P 2p peaks were observed in the XPS curves of bare LNMO (Figure 8a2) and Ta-LNMO pristine electrodes (Figure 8c2); however, the P 2p peaks were observed at ~137.7, ~136.8, 135.7, and ~134.8 eV for the bare LNMO (Figure 8b2) and Ta-LNMO (Figure 8d2) cycled electrodes, which can be attributed to the presence of phosphates (LixPFy and LixPOyFz) caused by the increased decomposition of LiPF6.62,66-67 After cycling, the bare LNMO electrode showed a higher amount of phosphates than the Ta-LNMO one, in agreement with the changes in the carbonate peaks.

The O 1s XPS spectra (Figures 8a3-8d3) display two main components: a peak at ~530 eV, characteristic of the O2- anions of LNMO,62-64,68 and peaks above 531 eV (CO32-, C-O-C, C-OH), assigned to the SEI species deposited on the surface of the electrodes.63 58 During cycling at 55 °C, the gradual increase in the peaks at energies greater than 531 eV further suggests increased coverage of the SEI species on the surface of the electrodes; the surface coverage was higher for the bare LNMO cycled electrode (Figure 8b3) than for the Ta-LNMO one (Figure 8d3). The F 1s peaks (Figures 8a4-8d4) are associated with poly(vinylidene fluoride) (PVDF, ~688 eV), phosphates (LixPFyOz and LixPFy, ~686-687.5 eV), and LiF (~685.5 eV).66,69 The relative intensities of the phosphate and LiF peaks also increased during cycling at 55 °C, due to the formation of surface SEI species. In addition, the XPS Ta 3d peaks for the pristine and cycled Ta-LNMO electrodes suggest that the amount of Ta element on the surface did not change significantly upon cycling (Figures 8c5 and 8d5). This is in contrast with the Al element on the surface of Al-LNMO during cycling, as will be discussed later.


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Figure 8. Surface XPS analysis of the bare LNMO and Ta-LNMO electrodes before and after the 20th electrochemical cycle at 55 °C. Surface C, P, O, and F contents for (a1-a4) pristine LNMO, (b1-b4) LNMO after 20 cycles at 55 °C, (c1-c4) pristine Ta-LNMO, and (d1-d5) Ta-LNMO after the 20th cycle. Surface Ta content for (c5) pristine and (d5) cycled Ta-LNMO electrodes. All peaks were adjusted according to the intensity of the C 1s peak (284.8 eV).

Discussion The high-voltage spinel LNMO cathode material exhibits excellent cycling performance at RT but shows severe degradation of cycling performance at elevated temperatures. The degradation has been attributed to distortion of the LNMO surface structure, decomposition of the electrolyte, and parasitic reactions between cathode and HF present in the electrolyte.19-22 In particular, the distortion of the surface structure has been directly observed by Cs corrected STEM, which showed migration of TM ions during charging to high voltage.15,30,51-54,56,70-78 This is similar to the processes widely observed in layered cathode materials, e.g., Li2MnO3 and Li2MnO3-LiMO2 (M = Ni, Co, Mn, etc.), which exhibit significant structural distortions from layered to spinel and finally to rocksalt structure.56,70-77 The migration of TM ions in LNMO during cycling is also related to their dissolution from the cathode structure to the electrolyte in the presence of trace amounts of HF.15,51,54 In this work, the amount of ACS Paragon Plus Environment


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dissolved Ni and Mn ions was estimated by storing the charged electrodes in the electrolyte for several days at 55 °C (Supporting Information, Figure S6), which revealed a fast increase in Ni and Mn dissolution, in agreement with previous reports.21,30 The structural distortion of LNMO during cycling at elevated temperatures is accompanied by the loss of a small amount of oxygen,60-61,79 which is expected to be highly oxidative and may be responsible for triggering a series of chain reactions and the formation of new HF.60-61,79-80 Although the small amount of oxygen evolved from the spinel structure is usually beyond the detection limit of mass spectroscopic techniques,60-61 the increase in capacity associated with the Mn3+/Mn4+ redox couple indicates loss of oxygen during cycling (Figure 7). The oxygen loss and associated structural reconstruction of bare LNMO and thickening of surface SEI during cycling suppress the diffusion of lithium, resulting in a fast increase in the impedance of the LNMO half-cell (Supporting Information, Figure S7).

The improvement of the electrochemical cycling performance of LNMO at 55 °C via surface coating of Ta2O5 is attributed to the stabilization of the surface structure by migration of a small amount of Ta5+ ions into the surface spinel structure, resulting in reduced Ni and Mn dissolution (Supporting Information, Figure S6). Although the overall crystal structure did not change significantly after the Ta2O5 coating, the Cs corrected STEM measurements clearly indicated the presence of Ta5+ in the 16c site of the spinel structure, particularly in the surface region (Figure 4), which is believed to limit the dissolution of Ni and Mn ions since their migration channels are preoccupied.30 The residual Ta2O5 on the surface of LNMO also played a significant role in the stabilization of the surface structure by providing resistance to HF attack, at variance with HF scavengers such as Al2O3, as discussed below. The stabilized surface structure leads to reduced oxygen evolution from the LNMO half-cell during cycling, as evidenced by the almost complete lack of increase in the capacity associated with the Mn3+/Mn4+ redox couple (Figure 7). Furthermore, the stabilized surface structure, reduced Mn dissolution, and reduced oxygen evolution result in a limited SEI formation (Figure 8) and a slow ACS Paragon Plus Environment

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increase in the impedance of the Ta-LNMO half-cell (Supporting Information, Figure S7). Another advantage of the Ta2O5-coated LNMO is that Ta5+ has high oxygen affinity, which may stabilize the surface oxygen and contribute to the stabilization of spinel structure.81

HF barriers such as Ta2O5 are more effective than HF scavengers such as Al2O3 as coating materials for high-voltage spinel LNMO, in order to improve its cycling performance at elevated temperatures. Al2O3-coated LNMO only shows improved electrochemical cycling performance in the initial cycles (Supporting Information, Figure S2). This attributed to the sacrificial reaction of Al2O3 with HF, as observed by the XPS analysis which revealed that the Al content on the surface vanished after few cycles at 55 °C (Supporting Information, Figure S8). However, due to the high cycling voltage of LNMO, HF may be constantly generated by decomposition of the electrolyte, leading to attack of AlLNMO by HF and degradation of cycling performance after a few cycles, similar to that occurring in bare LNMO during cycling.19-22 Note that Al2O3 coating was reported to improve the cycling performance of LiMn2O4, which may be due to the limited amount of HF generated in the electrolyte during cycling at low voltage.42 Our results suggest that the coating oxides which works for lowvoltage cathode materials may not be suitable for the high-voltage ones. Figure 9 shows a schematic illustration of LNMO coated with HF barriers and scavengers.



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Figure 9. Schematic illustration of LNMO coated with HF barriers and scavengers, before and after electrochemical cycling. LNMO coated with HF barriers shows a stable surface structure and limited deposition of SEI after cycling because of high resistance barriers against HF attack. In contrast, coating with HF scavengers results in a damaged surface and significant amount of SEI after a few cycles due to consumption of scavengers via sacrificial reactions with HF.

Conclusion The effects of Ta2O5 (HF barrier) coating on the electrochemical cycling performance of high-voltage spinel LNMO at room temperature and 55 °C have been investigated in detail via a combination of techniques, including STEM, XPS, XRD, and other methods. Ta2O5 shows higher resistance to HF attack compared to other common oxide coatings (HF scavengers) such as Al2O3. During the surface coating of LNMO with Ta2O5 at high temperature, a small amount of Ta5+ ions migrate into the 16c site of the spinel structure, thereby transforming the spinel structure into a rocksalt-like structure, particularly in the surface region. The diffusion of the high-valence Ta5+ ions into the spinel structure is ACS Paragon Plus Environment

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charge-compensated by the reduction of Mn4+ to Mn3+. The Ta2O5-coated LNMO half-cell shows excellent cycling performance at room temperature and 55 °C, despite the slightly low initial discharge capacity. Ta2O5-coated LNMO shows capacity retention and coulombic efficiency of approximately 97% and 99%, respectively, after 100 cycles at room temperature, and ~93% and 98%, respectively, after 100 cycles at 55 °C. Without surface coating, the capacity retention and coulombic efficiency after 100 cycles at room temperature are approximately 94% and 98%, respectively, and they decrease to ~76% and ~95%, respectively, after 100 cycles at 55 °C. The significant improvement in the electrochemical cycling performance of Ta2O5-coated LNMO at 55 °C is attributed to the stabilized surface structure and the HF resistance of the Ta2O5 coating. Furthermore, compared to LNMO coated with HF scavengers such as Al2O3, the Ta2O5-coated LNMO shows much better performance after prolonged cycling at 55 °C, due to the effective protection of the surface against HF attack. In conclusion, our work suggests that the surface coating of LNMO with HF barriers is one of the most promising methods to improve the high-temperature cycling performance of cathode materials for next-generation lithium-ion batteries, used in electrical and hybrid electrical vehicles.

ASSOCIATED CONTENT Supporting Information Available: TEM images, Electrochemical cycling, Rate capability, GITT, dQ/dV, Ni and Mn dissolution, Electrochemical impedance spectra and XPS spectra.

AUTHOR INFORMATION Corresponding Authors *Email: [email protected] (X. J. Huang) ACS Paragon Plus Environment


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Notes The authors declare no competing finical interest.

ACKNOWLEDGEMENTS This work was supported by the National Key R&D Program of China (Grant No. 2016YFB0100300) and the Strategic Priority Research Program of the Chinese Academy of Sciences (Grant No. XDA09010000).

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