Unusual Spinel-to-Layered Transformation in LiMn2O4 Cathode

Sep 21, 2017 - Distorted surface regions (5–6 nm) with an unusual layered-like structure on LiMn2O4 cathode material were directly observed after it...
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Unusual Spinel-to-Layered Transformation in LiMn2O4 Cathode Explained by Electrochemical and Thermal Stability Investigation Liubin Ben,†,‡ Hailong Yu,†,‡ Bin Chen,†,‡ Yuyang Chen,†,‡ Yue Gong,†,‡ Xinan Yang,†,‡ Lin Gu,†,‡,§ and Xuejie Huang*,†,‡ †

Beijing National Laboratory for Condensed Matter Physics, Institute of Physics, Chinese Academy of Sciences, Beijing, 100190, China ‡ School of Physical Sciences, University of Chinese Academy of Sciences, Beijing, 100190, China § Collaborative Innovation Center of Quantum Matter, Beijing 100190, China S Supporting Information *

ABSTRACT: Distorted surface regions (5−6 nm) with an unusual layered-like structure on LiMn2O4 cathode material were directly observed after it was cycled (3−4.9 V), indicating a possible spinel-to-layered structural transformation. Formation of these distorted regions severely degrades LiMn2O4 cathode capacity. As we attempt to get a better understanding of the exact crystal structure of the distorted regions, the structural transformation pathways and the origins of the distortion are made difficult by the regions’ nanoscopic size. Inspired by the reduction of Mn4+ to Mn3+ in surface electronic structures that might be associated with oxygen loss during cycling, we further investigated the atomic-level surface structure of LiMn2O4 by heat-treatments between 600 and 900 °C in various atmospheres, finding similar surface spinel-tolayered structural transformation only for LiMn2O4 heat-treated in argon atmosphere for a few minutes (or more). Controllable and measurable oxygen loss during heat-treatments result in Mn3+ for charge compensation. The ions then undergo a disproportionation reaction, driving the spinel-to-layered transformation by way of an intermediate LiMn3O4-like structure. The distortion of the surface regions can be extended to the whole bulk by heat-treatment for 300−600 min, ultimately enabling us to identify the bulk-level structure as layered Li2MnO3 (C2/m). This work demonstrates the critical role of Mn3+ in controlling the kinetics of the structural transformation in spinel LiMn2O4 and suggests heat-treatment in argon as a convenient method to control the surface oxygen loss and consequently reconstruct the atomic-level surface structure. KEYWORDS: lithium ion battery, spinel LiMn2O4, scanning transmission electron microscopy (STEM), defect-spinel LiMn3O4, layered Li2MnO3

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INTRODUCTION

cycling performance, because LMO is the parent of many highvoltage spinel cathode materials. Recent studies of cathode materials via advanced electron microscopy techniques, e.g., scanning transmission electron microscopy (STEM), have significantly improved our understanding of the surface’s atomic-level structure,20−26 structural evolution during electrochemical cycling,27−32 and stabilization of surface structure by modifying the cathode material.33−40 A recent review has been given by Lin Gu et al.26 In particular, layered Li2MnO3 reportedly shows migration of Mn ions from octahedral sites in the transition metal layers to tetrahedral sites in the lithium layers to form distorted regions resembling a spinel structure during cycling.30,31 Also, similar structural distortions are commonly observed in lithium-rich layered

Lithium-ion batteries have wide applications in portable electronics and potential applications in electric vehicles and hybrid electrical vehicles because they have both high operating voltage and high energy density.1−5 The most interesting cathode materials for lithium-ion batteries are spinel LiMn2O4 (LMO)6−11 and its derivatives, e.g., LiNi0.5Mn1.5O4,12−15 due to their low cost, nontoxicity, and the natural abundance of Mn, compared to the Co required for commercial LiCoO2. However, the cycling performance of these cathode materials degrades due to structural distortion during electrochemical cycling, particularly at elevated temperatures, hindering successful commercial application.16−19 Efforts to improve the cycling performance have recently focused on stabilizing the surface structure and minimizing reactions between the surface and the electrolyte.4,5,11 Understanding the surface structure of LMO during cycling is the key to effectively improve the © 2017 American Chemical Society

Received: July 31, 2017 Accepted: September 21, 2017 Published: September 21, 2017 35463

DOI: 10.1021/acsami.7b11303 ACS Appl. Mater. Interfaces 2017, 9, 35463−35475

Research Article

ACS Applied Materials & Interfaces cathode materials with the general formula xLi2MnO3·(1 − x)LiMO2 (Mn = Ni, Co, Mn, etc.), resulting in gradual formation of bulk-level spinel phase after prolonged cycling.35−38 Spinel cathode materials, though believed to be structurally stable, have been reported to show migration of Mn ions from octahedral sites to lithium tetrahedral sites in the fewnanometer surface region during cycling, resulting in a LiMn3O4-like defect-spinel structure.29,32 Interestingly, when spinel LMO is charged to high voltage even for one cycle, a severe structural distortion forms on the surface whose exact crystal structure is not clear.28 These distortions observed by STEM, though small, may contribute significantly to the degradation of cathode cycling performance. Leifer et al. also observed a possible layered phase in high-voltage cycled LiMn2O4 via Raman spectroscopy.41 However, since the distorted regions are usually nanoscopic, the exact structure and the migration pathways of ions during structural transformation are not readily observed, particularly for spinel cathode materials. Comprehensive understanding of the distorted surface structure is desperately needed in order to effectively improve the cycling performance of spinel cathode materials. If the atomic-level distorted surface regions observed during electrochemical cycling of LMO are associated with a structural transformation, it must be driven by thermodynamics and kinetics that we would like to ascertain.42 Unfortunately, no correlation between this structural transformation that occurs during electrochemical cycling and a thermodynamically induced structural transformation has yet been established. Furthermore, the arrangement of atoms in small distorted atomic-level regions observed by STEM may be associated with several possible structures viewed along particular directions; for example, a rocksalt-like arrangement of atoms (occupation of the 16c site of spinel)25,29,32 can be associated with MnO (Fm3m ̅ ) viewed along the [110] direction or Li2MnO3 (C2/m) along the [111] direction. Due to the size of the distorted regions after cycling, it is usually impossible to observe the arrangement of atoms along several crystallographic directions. Furthermore, previous thermodynamic studies of LMO have also resulted in much controversy.43−47 Two groups, Tsuji et al. and Thackery et al., separately observed bulk-level structural transformation in pristine LMO during heating to 1000 °C in air,43,44 but Massaroti et al. found only slight migration of Mn ions into the lithium tetrahedral sites and no change of the general spinel structure in their samples.45 Schilling et al. and MacNeil et al. respectively observed that electrochemically delithiated samples showed a structural transformation from spinel λ-MnO2 to beta-MnO2 at ∼350 °C.46,47 It is clear that establishing a correlation between the atomic-level structural transformation and a thermodynamically induced bulk-level structural transformation is critical to full understanding of the structural stability of cathodes. In the present work, we investigated the general and atomiclevel structure of LMO cathode materials after high-voltage (3− 4.9 V) and normal-voltage (3−4.3 V) electrochemical cycling, via a combination of techniques: STEM, X-ray photoelectron microscopy (XPS), X-ray diffraction (XRD), etc. To reveal the local and general structural transformations of LMO and their thermodynamic origins, we proceeded to process LMO between 600 and 900 °C in air, oxygen, and argon atmospheres, respectively. Spinel LMO showed stable structure during heattreatment in air or oxygen up to 900 °C, but severe structural distortion in argon even at 600 °C for a few minutes. The

structural transformation during electrochemical cycling was compared with that during heat-treatment. The thermodynamic and kinetic origins of the structural transformations in LMO are discussed in detail. We also provide further suggestions for applying heat-treatment in argon to reconstruct the surface of a cathode material.

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EXPERIMENTAL SECTION

Sample Preparation. LiMn2O4 (LMO) was prepared for electrochemical studies by the conventional solid-state reaction. Li2CO3 (Alfa Aeasar) and EMD (Xiangtan Electrochemical Scientific Ltd.) were mixed (mole ratio Li:Mn = 1.05:2) and milled by mortar and pestle for several hours. The mixture was heated at 900 °C for 10 h in a tube furnace with flowing oxygen atmosphere followed by slow cooling to room temperature. For thermodynamic investigations, asprepared LMO samples were further heat-treated variously at 600− 900 °C for 10−600 min in air or argon atmosphere. XRD Measurements. XRD patterns were obtained using a Bruker D8 ADVANCE diffractometer with a Cu Kα radiation source (λ1 = 1.54056 Å, λ2 = 1.54439 Å). The diffractometer was equipped with a LYNXEYE detector and operated at 40 kV and 40 mA. Electrochemically cycled cathode samples were taken out of the cell inside an argonfilled glovebox and then washed several times in DMC to remove surface electrolyte. After drying in an argon-filled glovebox, the cathode samples were sealed in a custom-made airtight sample holder for XRD. All XRD patterns were refined using the Rietveld method as implemented in the program TOPAS.48 Electrode Preparation. LMO working electrodes for a half-cell were prepared by spreading a slurry of active materials (90 wt %), carbon black (5 wt %), and PVdF binder (5 wt %) on Al foil. The electrodes were then dried at 100 °C in vacuum for 10 h before use. The loading level for each cathode is ∼10 mg/cm2. A coin-cell type half-cell was assembled with a lithium metal disk as the counter electrode, and a Celgard polypropylene as the separator saturated with a 1 M LiPF6 solution in a mixture of ethylene carbonate (EC) and dimethyl carbonate (DMC) in a 1:1 ratio by volume (LP30) in the argon-filled glovebox. The discharge/charge measurements were carried out on a Land BT2000 battery test system (Wuhan, China) at a current rate of C/5 at room temperature (C/5 refers to one Li insertion into LMO per formula unit in 5 h). STEM Characterization. STEM characterization was performed using a JEM-ARM 200F transmission electron microscope operated at 200 kV. The attainable spatial resolution of the microscope is 80 pm at the incident semiangle of 25 mrad. Electrochemically cycled cathode samples were taken out of the cell in an argon-filled glovebox and then washed several times in DMC to remove surface electrolyte. After drying in the argon-filled glovebox, the cathode samples were placed on copper grids, still in the argon-filled glovebox. The samples were then transferred to the STEM without exposure to air, through a custom-made mobile airlock. All STEM image simulations were based on the fast-Fourier-transforms multislice approach.49 The simulated STEM images were performed with uniform parameters, which include an accelerating voltage of 200 kV and a specimen thickness of 50 nm. The simulations with thickness ranging from 20 to 60 nm exhibit no qualitative difference. The incident semiangle is 25 mrad, acceptance semiangle is 12−25 mrad, Cs value is 0.001 mm, and defocus is −0.5 nm. X-ray Photoelectron Spectroscopy. The X-ray photoelectron spectroscopy (XPS) spectra were recorded with a spectrometer having Mg Kα radiation (ESCALAB 250, Sigma Probe, Thermo VG Scientific Co. Ltd.). All binding energies reported were corrected using the signal for carbon at 284.8 eV as an internal standard. The peak-fitting and quantitative evaluation was performed with the CasaXPS software. The background was corrected using the Shirley method. Thermal Stability Measurement. The thermal stability of LMO was studied via thermal gravimetric (TG) differential scanning calorimetry (DSC). A sealed stainless steel pan with a gold-plated copper seal (30 μL volume) was used to collect about 4 mg of each 35464

DOI: 10.1021/acsami.7b11303 ACS Appl. Mater. Interfaces 2017, 9, 35463−35475

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Figure 1. Charge−discharge profiles for the 1st, 7th, 17th, and 100th cycle of (a) a normal-voltage-cycled (3−4.3 V) and (b) a high-voltage-cycled (3−4.9 V) LMO half-cell. Capacity retention and Coulombic efficiency of (c) a normal-voltage-cycled (3−4.3 V) and (d) a high-voltage-cycled LMO half-cell.

Figure 2. (a) Typical STEM-HAADF image of normal-voltage-cycled LMO (16th cycle, fully charged). (b, c) Enlarged images of the surface and subsurface regions, corresponding respectively to the red and blue boxes in panel (a). (d1) Crystal structure of the defect-spinel LiMn3O4 viewed along the [110] direction. Mn atoms (purple) occupy both octahedral and tetrahedral sites. O atoms are depicted in red. (d2) Corresponding simulated STEM-HAADF image. (e, f) Line profiles corresponding to the red lines in panel (b), blue lines in panel (c) and purple lines in (d2), respectively. (Schematic lattice structures are overlaid in panels a−c.) sample. The scan rate was 10 °C/min and the atmosphere was a mixture of argon (80%) and oxygen (20%).

influence of electrolyte decomposition at high voltage.28 In previous work using the same electrolyte, we observed excellent room temperature cycling performance between 3.5 and 4.9 V for LiNi0.5Mn1.5O4, suggesting the electrolyte to be relatively stable up to 4.9 V.32,39 Results of the high-voltage cycling of the LMO half-cell are shown in Figure 1, along with those of normal-voltage cycling, for comparison. The normal-voltage-cycled LMO half-cell shows expected excellent cycling performance between 3 and 4.3 V at room temperature (Figure 1a and c), in agreement with previous reports.28,29 The initial discharge capacity is ∼110 mAh/g, which decreases slightly to ∼104 mAh/g after the 100th cycle, indicating good capacity retention of ∼95%. The average Coulombic efficiency is ∼99% (Figure 1c). In contrast, the LMO half-cell shows severe degradation after 3−4.9 V cycling for only ten cycles (7th−16th cycle) (Figure 1b and d),

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RESULTS Electrochemical Cycling of LMO. LMO cathode material generally shows stable cycling performance between 3 and 4.3 V, but only ∼80% of the lithium ions are utilized. More lithium ions can be used for intercalation if the upper cutoff voltage is higher than 4.3 V, but in this case, capacity degrades severely. To investigate the influence of cycling voltage on performance and structure, a LMO half-cell was cycled between 3 and 4.3 V for the first six cycles to get stabilized and then cycled between 3 and 4.9 V for ten cycles (7th−16th cycle). After the highvoltage (3−4.9 V) cycling, the upper cutoff voltage was lowered back to 4.3 V in the following cycles. The maximum voltage (4.9 V) applied during the high-voltage cycling was slightly lower than that in our previous work, in order to minimize the 35465

DOI: 10.1021/acsami.7b11303 ACS Appl. Mater. Interfaces 2017, 9, 35463−35475

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Figure 3. (a) Typical STEM-HAADF image of high-voltage cycled LMO (16th cycle, fully charged). (b, c) Enlarged images of the surface and subsurface regions, corresponding respectively to the red and blue boxes in panel (a). (d1) Crystal structure of the layered (e.g., Li2MnO3, C2/m) viewed along the [100] direction. Mn atoms (purple) occupy octahedral sites. O atoms are depicted in red. (d2) Corresponding simulated STEMHAADF image. (e, f) Line profiles corresponding to the red lines in panel (b), blue lines in panel (c), and black lines in panel (d), respectively. (Schematic lattice structures are overlaid in panels a−c.)

despite more lithium ions intercalating during high-voltage cycling, i.e., ∼120 mAh/g during charge and ∼115 mAh/g during discharge. The charge/discharge capacity decreases rapidly (17th−100th cycle) and the discharge capacity is only ∼84 mAh/g after the 100th cyclea capacity retention of only ∼76% (Figure 1d). Coulombic efficiency decreases (90−96%) during high-voltage cycling and gradually increases to ∼99% after the 40th cycle (Figure 1d). The fast degradation of cycling performance for the high-voltage-cycled LMO half-cell is in good agreement with previous reports.28 Since electrolyte decomposition is limited at room temperature, structural changes may play a significant role in the degradation of cycling performance. Crystal and Atomic-Level Structure of Electrochemically Cycled LMO. The general crystal structure of the cycled LMO after the 100th cycle remains cubic spinel without strong evidence of secondary phases, as observed by XRD (Supporting Figure S1). Lattice parameter a for the high-voltage-cycled LMO is ∼8.2085(3) Å, calculated by the Rietveld refinement method, which is similar to that of the normal-voltage-cycled LMO of ∼8.2098(1) Å but smaller than of that of pristine LMO of ∼8.2168(5) Å. The smaller lattice parameter a in the cycled LMO is partially attributed to loss of Mn ions.50 Further refinement of the site occupancy is not possible since the quality of XRD data is poor due to well-known surface coverage of electrolytes, binders, conducting additives, etc. Nevertheless, the XRD results here confirm that the spinel structure is generally stable even after prolonged high-voltage cycling. Thus, the degradation of cycling performance of LMO is likely associated with local structure. The local structure of LMO after electrochemical cycling was investigated by aberration corrected scanning transmission electron microscopy (STEM). The initial investigations started with the normal-voltage-cycled (3−4.3 V) LMO in the 16th cycle (fully charged). The results show no significant atomiclevel distortion in the bulk region of the cycled LMO even after prolonged cycling, as shown in SI Figure S2a, in accordance with the XRD results. Thus, the focus of the investigation of

local structure is on the surface region. A typical STEMHAADF image (Figure 2) shows an arrangement of Mn ions resembling the standard spinel structure viewed along the [110] crystallographic direction in the surface/subsurface region. The details of the assignment of Mn ions in the image to the standard spinel structure can be found in the simulated image (Supporting Figure S2c), and elsewhere.24,29 In the image, the subsurface region shows a structure similar to the bulk and the standard spinel, which can also be seen in the enlarged image (Figure 2b). The line profiles (Figure 2e1 and f1) confirm that the intensity at the Mn2 site is nearly half of that at the Mn1 site and there is no extra intensity at the lithium tetrahedral site, which is a typical feature of the arrangement of Mn ions in a standard LMO spinel structure.24,29 Fast Fourier transforms (FFTs) of the subsurface region (inset in Figure 2a) shows that reflections can be fully indexed to cubic Fd3̅m. The surface region (2−3 nm), however, demonstrates a slight migration of Mn ions into lithium tetrahedral sites to give a defect spinel LiMn3O4-like arrangement of Mn ions, as observed previously.25,29 Note that the size of the distorted surface region does not grow significantly, even in a sample cycled longer than the previous one. In addition, for an ideal LiMn3O4 structure, the intensity ratio of Mn2/Mn1 and Mn(t)/Mn1 is ∼0.5 in the line profile, as shown in the simulated Figure 2e3 and f3. The distorted LiMn3O4-like surface region shows similar Mn2/Mn1 (Figure 2e2), but much smaller Mn(t)/Mn1 (∼0.2) (Figure 2f2), indicating that the surface structure is in an intermediate stage of structural transformation from spinel LMO to defectspinel LiMn3O4. Note that, since the distorted surface region is only 2−3 nm, FFTs for this region do not provide clear information for further structural identification; see Supporting Figure S3. STEM-HAADF images of high-voltage-cycled LMO (16th cycle, fully charged) show large surface and subsurface regions to be severely distorted after ten cycles between 3 and 4.9 V (Figure 3), while the bulk region still retains the standard spinel structure (Figure S2b). The most salient feature is distorted surface regions with an unusual layered-like arrangement of Mn 35466

DOI: 10.1021/acsami.7b11303 ACS Appl. Mater. Interfaces 2017, 9, 35463−35475

Research Article

ACS Applied Materials & Interfaces ions. These regions can be clearly observed in the enlarged image (Figure 3c). The distance between the layers suggests the distorted regions have a d-spacing of ∼4.7 Å, similar to that of many layered cathode materials.22,27 Spots along the blue line, in particular, all show nearly the same contrast, as confirmed by the line profile (Figure 3e2 and f2), which is also in agreement with the simulated results, shown in Figure 3e3 and f3. Since the layered-like regions are small, no clear information could be obtained from FFTs (Supporting Figure S4). The LiMn3O4-like structure also presents in the image next to the layered-like region, as confirmed by the enlarged image in Figure 3b and line profiles in Figure 3e1 and f1, but the intensity ratio of Mn2/Mn1 and Mn(t)/Mn1 is much smaller than that of the standard LiMn3O4 shown in Figure 2, indicating Mn deficiency in this LiMn3O4 region, which will be explained later. Combined STEM observations of the surface region of cycled LMO (Figures 2 and 3) reveal the possible pathways of Mn ions during surface’s structural evolution induced by the increase of cycling voltage, as shown in Figure 4. During initial cycling between 3 and 4.3 V, Mn ions migrate from Mn octahedral sites in the standard LMO region (e.g., subsurface region of Figure 2) (Figure 4a) to lithium tetrahedral sites to form the LiMn3O4-like structure (e.g., surface region of Figure 2) (Figure 4b). On further increasing the voltage to 3−4.9 V, the surface LiMn3O4-like structure distorts by migration of Mn ions from the Mn(t) tetrahedral sites and Mn2 octahedral sites in some regions to fully occupy Mn2 octahedral sites in other regions nearby (possibly close to the outmost surface), resulting in a Mn-deficient LiMn3O4-like structure (e.g., subsurface region of Figure 3) (Figure 4c) and a newly formed layered-like structure (e.g., outermost surface region of Figure 3) (Figure 4d). All this suggests that the spinel-to-layered transformation is made possible via migration of Mn ions from octahedral sites to lithium tetrahedral sites, forming an intermediate LiMn3O4 structure, followed by a further shift of some of Mn ions from the LiMn3O4 back to the octahedral sites to form the layered-like structure. The movement of Mn ions is associated with oxygen loss and consequent existence of Mn3+ ions, which will be discussed later. Surface Electronic Structure of Cycled LMO. The surface electronic structure of the pristine and the electrochemically cycled LMO samples was investigated by X-ray photoelectron spectroscopy (XPS) which provides a surfacesensitive analysis of the presence of specific elements, particularly Mn ions. The XPS spectra for Mn 2p of the pristine and the cycled LMO samples (3−4.3 V and 3−4.9 V) are shown in Figure 5. All the spectra are normalized to Mn 2p3/2 for better comparison. Two typical peaks (Figure 5a) corresponding to spin−orbit splitting of Mn 2p3/2 and 2p1/2 are observed in the spectra for all samples.51 Fwhm’s (full width at half-maximum) larger than 3.5 eV appear in all the results for all the Mn 2p3/2 peaks, indicating the coexistence of mixed Mn3+ and Mn4+ ions. Fitting of Mn 2p3/2 peaks was performed to obtain the percentage of Mn3+ and Mn4+ ions and the details of which can be found elsewhere.29,51 If present at all, Mn2+ ions are relatively few, so fitting of Mn2+ ions is not included. The fitted parameters indicate that ∼48.31% Mn4+ and ∼51.79% Mn3+ ions are present on the surface of the pristine LMO (Figure 5b), consistent with the standard spinel LMO electronic structure. After normal-voltage cycling (3−4.3 V), the percentage of Mn4+ ions on the surface in the fully charged state (16th cycle) is not 100% but only ∼68.93%, while the percentage of Mn3+ ions present on the surface is ∼31.07%.

Figure 4. STEM-HAADF images and corresponding simulated crystal structures showing an increase of atomic-level structural distortion of spinel LMO in the surface and subsurface regions with increasing cycling voltage from 4.3 to 4.9 V. (a1) STEM-HAADF image showing arrangement of Mn atoms associated with the standard LMO spinel (e.g., subsurface region of Figure 2). (a2) Corresponding simulation of the standard LMO spinel crystal structure (b1) STEM-HAADF showing arrangement of atoms associated with the defect-spinel LiMn3O4-like structure after 3−4.3 V cycling (e.g., surface region of Figure 2). (b2) Corresponding simulation of the defect-spinel LiMn3O4-like crystal structure. (c1) STEM-HAADF image showing arrangement of atoms associated with the defect-spinel LiMn3O4-like structure after 3−4.9 V cycling (e.g., subsurface region of Figure 3). In this structure, the contrast associated with Mn(t) (purple cycles) and Mn2 (red cycles) is much weaker compared with the standard defectspinel LiMn3O4, suggesting Mn deficiency in the structure. (c2) Corresponding simulation of Mn-deficient LiMn3O4-like crystal structure. (d1) STEM-image showing the arrangement of atoms associated with the layered-like structure after 3−4.9 V cycling (e.g., surface region of Figure 3). (d2) Corresponding simulated layered-like structure.

When the cycling voltage is increased (3−4.9 V), the surface percentage of Mn4+ ions in the fully charged state (16th cycle) decreases to ∼61.27%, while that of Mn3+ ions increases to ∼38.83% (Figure 5d). The presence of Mn3+ ions on the surface of fully charged LMO particles suggests that oxygen loss may occur on the surface of LMO, charge-compensated by reduction of Mn4+ ions to Mn3+. The increase of Mn3+ ions with increasing cycling voltage further indicates that more oxygen may be lost with higher cycling voltage. However, direct measurement of the amount of oxygen loss during electrochemical cycling by other techniques such as differential 35467

DOI: 10.1021/acsami.7b11303 ACS Appl. Mater. Interfaces 2017, 9, 35463−35475

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Figure 5. (a) Mn 2P XPS spectra of pristine and cycled LMO. Fitted spectra of (b) pristine LMO, (c) normal-voltage-cycled (16th cycle, fully charged) LMO, and (d) high-voltage-cycled (16th, fully charged) LMO.

Figure 6. STEM-HAADF images of LMO after heat-treatment at (a) 900 °C for 10 h in oxygen, (b) 600 °C for 10 h in air, and (c) 900 °C for 10 h in air. (Schematic lattice structures are overlaid in panels a−c.)

Figure 7. (a) STEM-HAADF image of LMO-10m. (b, c) Enlarged images of the surface and subsurface regions, corresponding respectively to the red and blue boxes in panel (a). (d, e) Line profiles corresponding respectively to red lines in panel (b), blue lines in panel (c), and purple lines in Figures 2d2. (Schematic lattice structures are overlaid in panels a−c.) FFTs inserted in (a) indicate reflections (yellow boxes) associated with {0 0 2}spinel which should be symmetrically absent for spinel LMO structure.

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DOI: 10.1021/acsami.7b11303 ACS Appl. Mater. Interfaces 2017, 9, 35463−35475

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ACS Applied Materials & Interfaces electrochemical mass spectrometry is difficult, as only a small amount of oxygen is lost.52 Structure and Microstructure of Heat-Treated LMO. The investigations of electrochemically cycled LMO suggest that oxygen loss may play an important role in the structural transformation of LMO. Thus, further investigations of the surface atomic-level structure of LMO were performed via heattreatment at 600−900 °C in air and argon atmospheres, in an attempt to intentionally induce oxygen loss. The results show that spinel LMO is generally stable, and there is no strong evidence of atomic-level structural distortions even after heattreatment at 900 °C in oxygen (Figure 6a). For LMO heattreated in air between 600 and 900 °C (Figure 6b and c), it shows only slight surface distortion with formation of the LiMn3O4-like structure at 900 °C for 10 h (Figure 6c). No temperature higher than 900 °C was investigated since additional effects such as lithium evaporation may occur. For LMO heat-treated in flowing argon atmosphere, the atomic-level structure shows surprisingly severe distortion, even after only a few minutes heat-treated at 600 °C. Furthermore, the surface structural evolution with increasing duration of heat-treatment is similar to that during cycling but occurs over a much larger region. The general structure of LMO after 10− 30 min of heat-treatment remains phase-pure spinel, with no evidence of secondary phases; see Supporting Figure S5. However, a typical STEM-HAADF image of an LMO sample after 10 min of heat-treatment (LMO-10m) in argon at 600 °C (Figure 7) shows both the surface and the subsurface regions (∼10 nm) to be distorted from the standard spinel structure. The enlarged regions (Figure 7b and c) and line profiles (Figure 7d1/d2 and e1/e2) clearly show additional contrast on the lithium tetrahedral sites, evidence of surface and subsurface LiMn3O4-like structures, similar to that in cycled LMO (Figure 2c1/c2) but appearing in a much larger region. The line profiles also suggest that the Mn ions in the lithium tetrahedral sites in the surface region may have higher concentration by exhibiting stronger associated intensity Mn(t)/M1 ratio (Figure 7e2) compared to that in the subsurface region (Figure 7d2). This suggests that more oxygen is lost from the region closer to the outermost surface. Further information about the LiMn3O4-like structure is obtained from FFTs of image, as in the inset in Figure 7a, which shows that all the reflections but {0 0 2}spinel can be fully indexed to LMO (Fd3m ̅ ) cubic symmetry. The {0 0 2}spinel reflections should be symmetrically absent but they are observed, possibly due to uneven distribution of Mn ions in the {0 0 2}spinel plane in the surface and subsurface regions, as explained by Phillips et al.30 Further heat-treatment of LMO at 600 °C in argon for 30 min (LMO-30m) led to severe surface structure distortion the layered-like structure again, observed by STEM-HAADF. The general crystal structure, as observed by XRD, however, was still phase-pure spinel with a slightly increased lattice parameter (Supporting Figure S5). The surface region with the layered-like structure in LMO-30m (Figure 8) is much larger than that observed in cycled LMO (Figure 3c). Also, a LiMn3O4-like region is observed on the surface of some particles (not shown), similar to that of cycled LMO (Figure 2c) and heat-treated LMO-10m (Figure 7). Enlarged regions (Figure 8b and c) and line profiles (Figure 8d1/d2) further confirm that the intensity of Mn2 peaks is nearly the same as that of Mn1 peaks. The contrast observed here also has an interlayer spacing of ∼4.7 Å.

Figure 8. (a) STEM-HAADF image of LMO-30m. (b, c) Enlarged images of the surface and subsurface regions, corresponding, respectively, to the red and blue boxes in panel (a). (d) Line profiles corresponding respectively to the red line in panel (b), blue line in panel (c), and purple line 1 in Figure 3d2. (Schematic lattice structures are overlaid in panels a−c.)

Information about the structural evolution from spinel LMO to the layered-like structure during heat-treatment was also obtained from FFTs of Figure 8a, as shown in Figure 9. If the reflections are indexed purely on a monoclinic basis, e.g., layered LiMnO2 with C2/m symmetry, some anomalous reflections are present in the FFTs (indicated by blue cycles, Figure 9c1)which can be indexed as, e.g., {1 0 2}layered, {101}layered, {1 1 0}layeredshould be absent due to symmetry extinction in the monoclinic C2/m symmetry (see simulated image, Figures 9b1). However, these anomalous reflections can also be indexed as weak {2 2 0}spinel, {1 1 1̅}spinel, and {0 0 2̅}spinel, etc. on a spinel basis (Figures 9c2), which is present in the cubic Fd3m ̅ symmetry (see simulated image, Figures 9b2), apart from the {2 0 0}spinel, whose origin has been explained above. The most plausible explanation for the spinel-associated FFTs obtained from the layered-like region observed by STEM is that the distorted surface region is not purely a layered-like structure, but is in an intermediate stage of transformation, containing both the spinel and the layered structures. Close examination of the image in Figure 8a indeed suggests that some residual Mn ions are present in the interlayer positions, as would be found in a weak spinel-like structure integrated with the layered-like structure; see Supporting Figure S6. Surface Electronic Structure of Heat-Treated LMO. Response of the surface electronic structure to heat-treatment was also investigated. XPS spectra for Mn 2p of the pristine and heat-treated LMO (LMO-10m and LMO-30m) are shown in Figure 10. As before, all the spectra are normalized to Mn 2p3/2 for better comparison. Fitted results (Figure 10c) suggest that after 10 min of heat-treatment (LMO-10m), the percentage of Mn4+ ions decreases to ∼37.17% and Mn3+ increases to ∼62.83%. On further heat-treatment, Mn3+ dominates the surface of LMO-30m with a value of ∼70.83%, while Mn4+ decreases further to ∼29.17% (Figure 10d). The observed XPS results suggest that oxygen is lost from the LMO, particularly from the surface, during heat-treatment, and compensated via reduction of Mn4+ to Mn3+. An analogous charge compensation mechanism likely occurs on the surface of electrochemically cycled LMO. The oxygen loss results in the presence of oxygen vacancies in the heat-treated LMO, which is higher for LMO35469

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Figure 9. Schematics showing the Mn sublattice in (a1) layered structure, e.g, Li2MnO3 (C2/m) viewed along the [0 1 0] direction and (a2) spinel LMO (Fd3̅m) viewed along the [1̅ 1 0] direction. (b1) Simulated electron diffraction patterns corresponding to the layered structure and (b2) the spinel structure (output from Crystal Maker/SingleCrystal software). FFTs acquired from the distorted surface region in Figure 8 indexed with respect to (c1) the layered structure and (c2) the spinel structure. Blue cycles in c1 indicate reflections should be absent from the layered structure. Blue boxes in c2 indicate reflections should be absent from the spinel structure.

Figure 10. (a) Mn 2P XPS spectra of pristine and heat-treated LMO. Fitted spectra of (b) pristine LMO, (c) LMO-10m, and (d) LMO-30m.

30m than that of LMO-10m. These oxygen vacancies could be partially refilled during TG experiments in oxygen environment (80% argon and 20% oxygen), leading to increased weight during heating to 600 °C (Supporting Figure S7). Identification of Bulk-Level Crystal Structures Associated with the Distorted Regions. These atomic-level investigations of surface structure suggest that the structural evolution during electrochemical cycling is closely related to that during heat-treatment with intentional induced loss of oxygen. Deeper understanding of the exact bulk-level structure associated with the distorted surface regions observed by STEM is possible by heat-treatment of LMO in flowing argon for longer duration. Eventually, after 300−600 min of heattreatment, the distortions in the surface extend to the whole bulk, enabling their identification by bulk-level characterization techniques such as XRD. As shown in Figure 11, XRD patterns

of LMO after 300 and 600 min of heat-treatment (LMO-300m and LMO-600m) show a significant number of additional diffraction peaks, beyond the main peaks associated with the spinel structure. These extra diffraction peaks are fully indexed to defect spinel LiMn3O4 (space group I41/amd, ICSD #164994) and layered Li2MnO3 (space group C2/m, ICSD #166861). The thermodynamically driven spinel LiMn2O4 to defect spinel LiMn3O4 transformation was reported previously but it occurred at higher temperature (>860 °C) in air.53 Note that the precise amount of lithium in LiMn3O4 cannot be obtained by XRDfurther neutron diffraction is required. Rietveld refinement (Supporting Figure S8) reveals that LMO300m (Figure S8a) is a phase-mixture of ∼93.83% spinel LMO, ∼0.57% defect spinel LiMn3O4, and ∼5.60% layered Li2MnO3, while it is ∼84.43%, ∼3.94%, and ∼11.63% for LMO-600m (Figure 8b). The refined parameters are shown in Supportings 35470

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Figure 11. XRD patterns and corresponding enlarged images of pristine LMO and heat-treated LMO (600 °C for 300 and 600 min). Blue and red asterisks indicate the reflections associated with the defect-spinel LiMn3O4 (I41/amd) and the layered Li2MnO3 (C2/m), respectively. Standard XRD patterns for LiMn3O4, Li2MnO3, and LiMn2O4 are placed at the bottom.

Computational simulations suggest that the layered structure is more energetically favorable than fully charged LMO with the spinel structure (λ-MnO2).54 Furthermore, theoretical studies have also shown that spinel LMO decomposes into layered Li2MnO3 and defect-spinel LiMn3O4.42 However, such transformations from spinel LMO via LiMn3O4 to layered Li2MnO3 require migration of Mn4+ ions from one octahedral site through an intermediate tetrahedral site in the initial spinel structure, which is unlikely in oxygen-stoichiometric LMO, due to the strong octahedral preference of Mn4+ ions.42 The structural transformation is made possible only if a small amount of oxygen is lost, resulting in reduction of Mn4+ ions to Mn3+ or lower for charge compensation, as observed by XPS for cycled or heat-treated LMO samples (Figure 5 and Figure 10). The Mn3+ ions prefer tetrahedral structures, and the mobility of Mn3+ ions is high within the spinel structure.55 Thus, Mn ions migrate to occupy lithium tetrahedral sites, resulting in the formation of John-Teller distorted defect-spinel (LixMny)8a(Mn1−y)16dO4‑δ (I41/amd), as observed by STEM (Figure 2b and Figure 7b). On further loss of oxygen, more Mn ions migrate to the lithium tetrahedral sites, leading to a disproportionate reaction: 2 Mn3+→ Mn2+ + Mn4+.55 The Mn2+ ions may remain at the lithium tetrahedral site, but Mn4+ ions prefer octahedral sites and may migrate back to the octahedral sites, resulting in formation of the layered Li2MnO3 (Figure 3c and Figure 8a). The formation and growth of the layered Li2MnO3 via the defect-spinel LiMn3O4 is also seen by the Mndeficient LiMn3O4 region observed in Figure 3b and FFTs of the layered region showing integrated layered and spinel structures in Figure 9. Thus, during the spinel-to-layered structural transformation, the concentration of Mn3+ ions induced by oxygen loss plays a dominant role in controlling the kinetics. In contrast, temperature has only a limited role during the structural transformation, since LMO still has standard spinel structure after heat-treatment up to 900 °C in either air or oxygen; see Figure 6. This may explain the controversial results obtained in previous thermodynamic studies of LMO, wherein different amounts of oxygen may have been lost due to preparation history, resulting in different degrees of observed structural transformation, even when the temperatures were similar.43−45

Table S1 and S2. The reversibility of the formation of LiMn3O4 and Li2MnO3 was also investigated by “reverse-processing” LMO-300m and LMO-600m in flowing O2 at 600 °C for several hours. The results show that the extra diffraction peaks disappear, suggesting that the layered-to-spinel structure transformation is fully reversible (Supporting Figure S9).

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GENERAL DISCUSSION Combined investigations of electrochemically cycled and thermally treated LMO reveal that the unusual layered-like distorted surface region observed via STEM is associated with a structural transformation from spinel LMO (Fd3̅m) to layered Li2MnO3 (c2/m), via the intermediate defect spinel LiMn3O4 (I41/amd), as shown in the schematic in Figure 12. Transformation between these structures involves only migration of Mn ions, without requiring oxygen rearrangement since all the structures share a cubic closed-packed oxygen anion sublattice.

Figure 12. Schematic of Mn ion migration in LMO surface with associated formation of the defect-spinel LiMn3O4 (I41/amd) and the layered Li2MnO3 (C2/m) structures. LMO is represented by yellow, defect-spinel LiMn3O4, and layered Li2MnO3 structures are indicated by green and orange, respectively. During normal-voltage cycling or heat-treatment at 600 °C for a limited time, a small amount of Mn ions migrate into the Li tetrahedral sites in the surface region to from the LiMn3O4 structure, accompanied by loss of a small amount of oxygen. If the cycling voltage is high or the duration of heat-treatment is increased, the surface structure shows severe distortion by migration of Mn ions from LiMn3O4 to form layered Li2MnO3, accompanied by a large loss of oxygen. The size of the regions, indicated by colors, is not drawn to scale. 35471

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initial cycles, due to local oxygen retention.30,68 This is consistent with our heat-treated LMO which shows reversible transformation between spinel to layered structures governed by the atmosphere (Figure S9). In our cycled samples, no such reversible transformation was observed, since the distortion occurs in the surface where oxygen is easily lost and consumed to oxidize the electrolyte. In addition, a spinel-to-layered transformation is commonly observed when Na is inserted into spinel λ-MnO2. Tarascon et al.69 and Yabuuchi et al.70 reported that sodium ions could be inserted into the chemically charged spinel phase (λ-MnO2), leading to a partial phase transition from spinel to monoclinic layered NayMnO2 phase. Also, Yang et al. reported that structural transformation from defect spinel Mn3O4 to layered birnessite (NaδMnOx·nH2O) took place in neutral electrolyte Na2SO4 during electrochemical cycling. Our work here clearly provides new structural information about the structural stability of cathode materials that are either lithiumrich or spinel-structured, relevant to Na intercalation.

Many cathode materials lose a small amount of oxygen during cycling, which is caused by overcharging the electrode56 leading to oxidation of the O 2p band and consequent release of oxygen ions from the lattice.42 The amount of oxygen lost depends on the overlap between O 2p and transition metal d band as well as cycling voltage. The oxygen loss in spinel during cycling is generally quite limited,52 beyond the detection limit of most techniques, due to the relatively stable structure and, theoretically, no overlap between O 2p and Mn2+/Mn3+, Mn3+/ Mn4+.5,57 Layered cathode materials and lithium-rich layered cathode materials lose a significant amount of oxygen, which is directly detected.52,58 The loss of oxygen has also been observed to increase as cycling voltage increases. Note that recent studies show that an overcharge below 4.5 V may only induce holes in the lattice.52 Furthermore, even if there is still a small amount of Li remaining inactive in the LMO cathode material during cycling, overcharge and resulting oxygen loss still occur at the surface, because lithium content is lower or the number of lithium vacancies at the exterior of the particle is high.58−62 The surface region where the cathode is in close contact with the electrolyte exhibits significant loss of oxygen, possibly because the energy required for release of oxygen is lower than in the bulk, or because of interaction with the electrolyte.30,56,63,64 The oxygen loss during cycling is chargecompensated by reduction of Mn4+ ions to Mn3+ or lower, as confirmed by surface XPS in Figure 5. Such charge compensation was further confirmed by observations of LMO heat-treated in argon atmosphere, where the amount of oxygen loss was semiquantitatively measured by TG (Figure S7) and XRD investigations (Figure S8). During electrochemical cycling of LMO, the distorted surface regions do not grow significantly, since a continuous interaction with the electrolyte leads to Mn dissolution from the surface structure in the electrolyte.29 It is generally understood that the amount of Mn dissolved in the electrolyte increases along with charging voltage.65 This agrees with increased distortion observed here in the surface region with the increase of cycling voltage (Figure 4). The formation of layered Li2MnO3 at the expense of LiMn3O4 results in Mn-deficient LiMn3O4 regions during electrochemical cycling (Figures 3b), and Mn atoms in the regions may become weakly bonded, leading to delamination of the surface atomic layers into the electrolyte during cycling.66 Thus, further bulk-level structural information on the distorted region is difficult to obtain, even after prolonged electrochemical cycling with fast capacity degradation (Figure 1) induced by the locally distorted regions. However, heat-treatment of LMO in argon induces oxygen loss that can be gradually increased, resulting in slow formation and growth of the distorted surface structures which can be extended to the whole bulk (Figure S8). The oxygen loss during cycling is consumed to oxidize the electrolyte, whereas it is taken away by heat-treatment by flowing argon.52,67 The heat-treatment of LMO cathode materials with a controlled rate of oxygen loss is effective in reconstructing the surface structures, which is applicable for many other cathode materials with interesting cycling performance. Spinel-to-layered structural transformation during electrochemical cycling is not commonly observed, despite its significant influence on the cycling performance, due to its isolation in nanoscale regions. Recent studies of layered lithium-rich cathode materials by Phillips et al. and Shimoda et al., both using STEM, observed a possibly reversible spinelto-layered transformation in the interior of the cathode in the

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CONCLUSION Spinel LMO is not only an important parent material of many spinel-related cathode materials but also the well-known degraded structure assumed by many lithium-rich cathode materials during electrochemical cycling. Understanding the structural stability of LMO during electrochemical cycling is of vital importance for improving the cycling performance of both spinel and layered lithium-rich cathode materials. Electrochemically cycled (3−4.9 V) LMO shows a stable bulk-level spinel structure but distorted surface regions (5−6 nm) with an unusual layered-like structure. The distorted regions do not grow significantly on cycling because of dissolution/delamination during repeated reactions with the electrolyte. Such distorted surface regions are closely related to oxygen loss surface XPS shows reduction of Mn4+ ions to Mn3+ for charge compensation. The surface’s structural distortion and the associated small oxygen loss are fundamental to the degradation of electrochemical cycling performance of LMO, due to increased dissolution of Mn ions from the surface structure into the electrolyte, as well as increased impedance due to Mn ions blocking lithium migration pathways. The exact structures associated with the distorted surface regions and underlying thermodynamic/kinetic origins are investigated by heat-treatment of LMO between 600 and 900 °C in various atmospheres. The structural transformation of LMO is strongly dependent on the amount of Mn3+ ions generated, which mainly occurs during heat-treatment in argon, even at 600 °C for a only a few minutes. In contrast, LMO heat-treated between 600 and 900 °C in air or oxygen does not show strong evidence of surface structure distortion. After only 10 min of heat-treatment in argon, larger surface regions showing LiMn3O4-like structure are observed, while both the LiMn3O4-like and the layered-like structures are observed after 30 min of heat-treatment in argon. Further identification of the exact crystal structures associated with the distorted surface regions is realized by heat-treatment of LMO for extremely long duration. After 300 and 600 min heat-treatment in argon, LMO clearly shows bulk-level XRD reflections associated with LiMn3O4 (I41/amd) and layered Li2MnO3 (C2/m). Our work here explains that the complicated origin of the unusual layered LMO surface structure during high voltage cycling is associated with migration of Mn3+ ions generated by oxygen loss for charge compensation and a disproportionation reaction of Mn3+ ions. Mn3+ ions prefer tetrahedral sites and are very mobile within the spinel structure. 35472

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(12) Liu, J.; Manthiram, A. Understanding the Improvement in the Electrochemic al Properties of Surfac e Modified 5 V LiMn1.42Ni0.42Co0.16O4 Spinel Cathodes in Lithium-ion Cells. Chem. Mater. 2009, 21 (8), 1695−1707. (13) Zhong, Q.; Bonakdarpour, A.; Zhang, M.; Gao, Y.; Dahn, J. R. Synthesis and Electrochemistry of LiNixMn2‑xO4. J. Electrochem. Soc. 1997, 144 (1), 205−213. (14) Kim, J. H.; Myung, S. T.; Yoon, C. S.; Kang, S. G.; Sun, Y. K. Comparative Study of LiNi0.5Mn1.5O4‑δ and LiNi0.5Mn1.5O4 Cathodes Having Two Crystallographic Structures: Fd-3m and P4332. Chem. Mater. 2004, 16 (5), 906−914. (15) Cabana, J.; Casas-Cabanas, M.; Omenya, F. O.; Chernova, N. A.; Zeng, D.; Whittingham, M. S.; Grey, C. P. Composition-Structure Relationships in the Li-Ion Battery Electrode Material LiNi0.5Mn1.5O4. Chem. Mater. 2012, 24 (15), 2952−2964. (16) Park, S. B.; Shin, H. C.; Lee, W.-G.; Cho, W. I.; Jang, H. Improvement of Capacity Fading Resistance of LiMn2O4 by Amphoteric Oxides. J. Power Sources 2008, 180 (1), 597−601. (17) Dai, Y.; Cai, L.; White, R. E. Capacity Fade Model for Spinel LiMn2O4 Electrode. J. Electrochem. Soc. 2013, 160 (1), A182−A190. (18) Yoon, T.; Park, S.; Mun, J.; Ryu, J. H.; Choi, W.; Kang, Y. S.; Park, J. H.; Oh, S. M. Failure Mechanisms of LiNi0.5Mn1.5O4 Electrode at Elevated Temperature. J. Power Sources 2012, 215, 312−316. (19) Manthiram, A.; Chemelewski, K.; Lee, E.-S. A Perspective on the High-Voltage LiMn1.5Ni0.5O4 Spinel Cathode for Lithium-Ion Batteries. Energy Environ. Sci. 2014, 7 (4), 1339−1350. (20) Shao-Horn, Y.; Croguennec, L.; Delmas, C.; Nelson, E. C.; O’Keefe, M. A. Atomic Resolution of Lithium Ions in LiCoO2. Nat. Mater. 2003, 2 (7), 464−467. (21) Sharifi-Asl, S.; Soto, F. A.; Nie, A.; Yuan, Y.; Asayesh-Ardakani, H.; Foroozan, T.; Yurkiv, V.; Song, B.; Mashayek, F.; Klie, R. F.; Amine, K.; Lu, J.; Balbuena, P. B.; Shahbazian-Yassar, R. FacetDependent Thermal Instability in LiCoO2. Nano Lett. 2017, 17 (4), 2165−2171. (22) Yan, P. F.; Zheng, J. M.; Lv, D. P.; Wei, Y.; Zheng, J. X.; Wang, Z. G.; Kuppan, S.; Yu, J. G.; Luo, L. L.; Edwards, D.; Olszta, M.; Amine, K.; Liu, J.; Xiao, J.; Pan, F.; Chen, G. Y.; Zhang, J. G.; Wang, C. M. Atomic-Resolution Visualization of Distinctive Chemical Mixing Behavior of Ni, Co, and Mn with Li in Layered Lithium TransitionMetal Oxide Cathode Materials. Chem. Mater. 2015, 27 (15), 5393− 5401. (23) Hu, E.; Lyu, Y.; Xin, H. L.; Liu, J.; Han, L.; Bak, S.-M.; Bai, J.; Yu, X.; Li, H.; Yang, X.-Q. Explore the Effects of Microstructural Defects on Voltage Fade of Li- and Mn-Rich Cathodes. Nano Lett. 2016, 16 (10), 5999−6007. (24) Huang, R.; Ikuhara, Y. H.; Mizoguchi, T.; Findlay, S. D.; Kuwabara, A.; Fisher, C. A. J.; Moriwake, H.; Oki, H.; Hirayama, T.; Ikuhara, Y. Oxygen-Vacancy Ordering at Surfaces of Lithium Manganese(III,IV) Oxide Spinel Nanoparticles. Angew. Chem., Int. Ed. 2011, 50 (13), 3053−3057. (25) Amos, C. D.; Roldan, M. A.; Varela, M.; Goodenough, J. B.; Ferreira, P. J. Revealing the Reconstructed Surface of LiMn2O4. Nano Lett. 2016, 16 (5), 2899−2906. (26) Gu, L.; Xiao, D. D.; Hu, Y. S.; Li, H.; Ikuhara, Y. Atomic-Scale Structure Evolution in a Quasi-Equilibrated Electrochemical Process of Electrode Materials for Rechargeable Batteries. Adv. Mater. 2015, 27 (13), 2134−2149. (27) Lu, X.; Sun, Y.; Jian, Z.; He, X.; Gu, L.; Hu, Y.-S.; Li, H.; Wang, Z.; Chen, W.; Duan, X.; Chen, L.; Maier, J.; Tsukimoto, S.; Ikuhara, Y. New Insight into the Atomic Structure of Electrochemically Delithiated O3-Li(1‑x)CoO2 (0 ≤ x ≤ 0.5). Nano Lett. 2012, 12 (12), 6192−6197. (28) Tang, D.; Ben, L.; Sun, Y.; Chen, B.; Yang, Z.; Gu, L.; Huang, X. Electrochemical Behavior and Surface Structural Change of LiMn2O4 Charged to 5.1 V. J. Mater. Chem. A 2014, 2 (35), 14519−14527. (29) Tang, D. C.; Sun, Y.; Yang, Z. Z.; Ben, L. B.; Gu, L.; Huang, X. J. Surface Structure Evolution of LiMn2O4 Cathode Material upon Charge/Discharge. Chem. Mater. 2014, 26 (11), 3535−3543.

Manganese ions initially migrate to lithium tetrahedral sites to form defect-spinel LiMn3O4. With continuing migration, disproportionation occurs (2 Mn3+→ Mn2+ + Mn4+), driving Mn4+ ions back to the octahedral site to form layered Li2MnO3. Our work here also provides a convenient method to control the surface oxygen loss and consequently reconstruct the atomic-level surface structure, by heat-treatment of cathodes in argon.

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ASSOCIATED CONTENT

S Supporting Information *

The Supporting Information is available free of charge on the ACS Publications website at DOI: 10.1021/acsami.7b11303. XRD patterns of pristine and cycled LMO, STEM images and FFTs of cycled LMO (bulk region), XRD patterns and Rietveld refinement of heat-treated LMO, TG results of heat-treated LMO (PDF)

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AUTHOR INFORMATION

Corresponding Author

*E-mail: [email protected] (X. J. Huang). ORCID

Yue Gong: 0000-0002-5764-3117 Lin Gu: 0000-0002-7504-031X Xuejie Huang: 0000-0001-5900-678X Notes

The authors declare no competing financial interest.

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ACKNOWLEDGMENTS 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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REFERENCES

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ACS Applied Materials & Interfaces

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DOI: 10.1021/acsami.7b11303 ACS Appl. Mater. Interfaces 2017, 9, 35463−35475

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DOI: 10.1021/acsami.7b11303 ACS Appl. Mater. Interfaces 2017, 9, 35463−35475