Surface Structure Evolution of LiMn2O4 Cathode Material upon

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Surface structure evolution of LiMn2O4 cathode material upon charge/discharge Daichun Tang, Yang Sun, Zhenzhong Yang, Liubin Ben, Lin Gu, and Xuejie Huang Chem. Mater., Just Accepted Manuscript • DOI: 10.1021/cm501125e • Publication Date (Web): 09 May 2014 Downloaded from http://pubs.acs.org on May 29, 2014

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Chemistry of Materials

Surface structure evolution of LiMn2O4 cathode material upon charge/discharge Daichun Tang, Yang Sun, Zhenzhong Yang, Liubin Ben, Lin Gu*, and Xuejie Huang* Beijing National Laboratory for Condensed Matter Physics, Institute of Physics, Chinese Academy of Sciences, Beijing 100190, China. ABSTRACT Surface dissolution of manganese is a long standing issue hindering the practical application of spinel LiMn2O4 cathode material, while few studies concerning the crystal structure evolution at surface area have been reported. Combining X-ray photoelectron spectroscopy (XPS), electron energy loss spectroscopy (EELS), scanning transmission electron microscopy (STEM), and density functional theory (DFT) calculations, we investigate the chemical and structural evolutions on the surface of LiMn2O4 electrode upon cycling. We found that an unexpected Mn3O4 phase was present on the surface of LiMn2O4 via the application of an advanced electron microscopy. Since the Mn3O4 phase contains 1/3 soluble Mn2+ ions, formation of this phase contributes significantly to the Mn2+ dissolution in LiMn2O4 electrode upon cycling. It is further found that the Mn3O4 appears upon charge and disappears upon discharge, coincident with the valence change of Mn. Our results shed light on the importance of stabilizing the surface structure of cathode material, especially at charged state. The understanding of the manganese dissolution reaction that occurs in the LiMn2O4 can certainly be extended to other oxide cathodes. Key words: LiMn2O4, surface structure evolution, Mn3O4, EELS, STEM imaging

INTRODUCTION

ACS Paragon Plus Environment


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Lithium-ion batteries are of great significance in developing hybrid electric vehicles (HEV) and electric vehicles (EV), which are expected to alleviate the increasingly serious air pollution.1, 2 Spinel LiMn2O4 is a promising cathode material of Li-ion battery for electric vehicles, owing to its low cost, high power density, environmental friendliness, and good safety.3, 4 However, one major disadvantage of LiMn2O4 is the severe capacity fading during electrochemical cycling or extended storage time, especially at elevated temperature. Various capacity fading mechanisms have been proposed, e.g., electrochemical reaction with electrolyte at high voltage,5 instability of the two-phase structure in the charged state leading to a more stable single-phase structure,6,

7

phase transformation from cubic spinel to tetragonal

rock-salt structure due to nonequilibrium lithiation,8 loss of crystallinity during cycling,9, 10 and manganese dissolution.6, 11, 12 Among these, manganese dissolution is generally believed to be the main cause of capacity degradation,7, 13-15 whereas the origin of manganese dissolution has not been sufficiently understood and many proposed mechanisms are controversial.6, 7, 16, 17

Hunter proposed that the pristine LiMn2O4 could partially dissolve under acid

conditions through the following disproportionation reaction, 2LiMn2O4 → 2Li2O + 3MnO2 + MnO.16 As the Li+ and Mn2+ are soluble in acid solution, λ-MnO2 is relatively stable and is left after a complete reaction. This disproportionation reaction is usually adopted to explain the manganese dissolution phenomenon of LiMn2O4 electrode in Li-ion batteries.18-21 However, it should be noted that the lithium content of LixMn2O4 varies with the state of (dis)charge. Jang et al. first reported that manganese dissolution is noticeably high at charged states,6 which has been confirmed by latter studies.22-25 This finding indicates that the manganese dissolution is dependent on the state of charge, i.e., the lithium concentration of LixMn2O4 particle. Moreover, the manganese dissolution rate of LiMn2O4 electrode reaches the maximum at the state close to fully delithiation (λ-MnO2), contradicted to Hunter’s theory. The contradiction suggests that the manganese dissolution mechanism could be more complicated than the aforementioned disproportionation reaction. Besides,


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the reason why manganese dissolution rate increase with increased state of charge22, 24, 25

needs to be clarified. It is natural to speculate that the Mn dissolution rate is associated with the crystal

structure on the surface of LiMn2O4. However, most previous works on the issue of manganese dissolution have not addressed the accompanied crystal structure change of LixMn2O4 upon cycling.21, 26 It is usually considered that the λ-MnO2 host is stable with cycling,27, 28 with only Li ion extracted in charge and re-inserted after discharge. Nevertheless, the surface structure evolution, which is highly related to the manganese dissolution process but could be hardly probed by diffraction techniques, needs to be clarified. In this paper, X-ray photoelectron spectroscopy (XPS) and electron energy loss spectroscopy (EELS) are used to identify the valence change of Mn ions, and scanning transmission electron microscope (STEM) to directly image the atomic structure of LixMn2O4 particle, at different states of (dis)charge. It is found that the valence change of Mn ions at the surface region during (dis)charge differs from the case in the bulk region. The surface structure evolution accounting for this unexpected phenomenon is identified, and its possible association with manganese dissolution is discussed.

EXPERIMENTAL SECTION Synthesis. LiMn2O4 was prepared by the conventional solid-state reaction. The mixture of stoichiometric amounts of Li2CO3 (Alfa Aeasar) and EMD (Xiangtan Electrochemical Scientific Ltd.) was used as the starting material. The mixture was heated for 10 h at 900℃ in an oxygen atmosphere followed by slow cooling to room temperature. Electrochemical Characterization. For the fabrication of the electrode, LiMn2O4 was mixed with conductive acetylene black and polyvinylidene fluoride (PVdF) binder in a weight ratio of 8:1:1. The electrolyte used was EC (ethylene carbonate) and DMC (dimthyl carbonate) (1:1 by volume)/1M LiPF6. Lithium metal disc was used as the anode. Cells were assembled in an argon-filled glove box.



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Charge and discharge tests of the cells were carried out at room temperature. The experiment was performed on the charge and discharge of constant current within the operating voltage range of 4.3-3.0V with a current density of 0.2C. For evaluating the cathode materials, the cells were disassembled in an argon-filled glove box, and the cathode was rinsed with DMC solution for removing the electrolyte salt. X-ray photoelectron spectroscopy. The 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 the 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. The change in manganese valence state at different charged and discharged states were interpreted from the XPS data. TEM-EELS. Electron energy-loss spectroscopy (EELS) was performed using a Tecnai F20 transmission electron microscope at an acceleration voltage of 200kV. The EELS spectra of O-K and Mn-L2,3 edge were collected from the surface of LiMn2O4 samples at different charge/discharge states. The detailed information is shown in Supporting Information Figure S2. The spectral background was removed by the power law fitting. The Mn valence was analyzed according to the methods used in ref. 29. Microscopy. An aberration-corrected scanning transmission electron microscope JEM ARM200F (JEOL, Tokyo, Japan) equipped with two CEOS (CEOS, Heidelberg, Germany) probe aberration correctors was used to probe the structure at atomic scale. Z-contrast annular bright field (ABF) and high-angle annular dark field (HAADF) imaging was performed in thin specimen regions with a probe convergence angle of 25 mrad. The collection angle of them is 12-24 mrad and 70-250 mrad, respectively. For each sample, the representative images were selected for analysis. First-principles calculations. Spin-polarized calculations were performed using the Vienna ab initio simulation package (VASP)30,


31

within the projector

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augmented-wave approach.32 Generalized gradient approximation (GGA) in the parametrization of Perdew, Burke, and Ernzerhof (PBE)33 pseudopotential was used to describe the exchange−correlation potential. A Hubbard-type correction U was taken into account due to the strongly correlated nature of the Mn 3d electrons.34,

35

According to a previous work, the effective Un value was set to 5 eV.36 A kinetic energy cutoff of 520 eV was used in all calculations. Geometry optimizations were performed by using a conjugate gradient minimization until all the forces acting on ions were less than 0.01 eV/Å per atom. K-point mesh with a spacing of ca. 0.03 Å−1 was adopted. Although the actual spin sequence of Mn ions may be complex in various manganese oxides, herein we adopt the ferromagnetic spin configuration to simplify the problem in all calculations as the spin sequence has negligible impact on total energy.

RESULTS AND DISCUSSION Rietveld refinement of the XRD patterns shows that the sample can be fully indexed to cubic spinel LiMn2O4 with a lattice parameter of 8.24 Å (space group: Fd-3m), as shown in Supporting Information Figure S1. The first and second charge/discharge profiles of the cell at a rate of 0.2C at room temperature are shown in Figure 1a. When the upper cut-off voltage is 4.3V, the first cycle charge and discharge capacity are 117.3 mAh g-1 and 111.9 mAh g-1, respectively. The second cycle is almost overlapped with the first one. The corresponding differential capacity plots shown in Figure 1b indicate two oxidation peaks located at about 4.04V and 4.16V, and two reduction peaks located at about 4.0V and 4.12V. The cycling performance of LiMn2O4 in Figure 1c shows that the capacity fades slowly during cycling, with a capacity retention of 89.8% after 100 cycles. All these electrochemical performances are in accordance with typical LiMn2O4.37



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85 0

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Figure 1. (a) Charge/discharge curves of the first and second cycles. (b) The corresponding differential capacity (dQ/dV) plots. (c) Cycling performance of LiMn2O4 in the voltage range of 3.0-4.3V at a rate of 0.2C.

XPS provides a surface-sensitive analysis on the chemical environment of specific elements. Figure 2a shows the XPS spectra for Mn 2p at different states of (dis)charge. All the spectra are normalized to Mn2p3/2 for better comparison. Two main peaks, which correspond to Mn 2p3/2 and 2p1/2 respectively, are observed in the spectra for the pristine LiMn2O4.38 After being charged to 4.3V, the Mn 2p3/2 shows


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asymmetric characteristics, which can be interpreted as multiplet splitting from Mn4+ state.38 The discharged LiMn2O4 exhibits similar spectral features with the pristine LiMn2O4. In order to obtain detailed information on the valence states of Mn at different states, curve fitting was conducted on the Mn 2p3/2 spectra (Figure 2b-d), and the details of fitting are shown in Supporting Information (Section S1). The fitting parameters are in agreement with that of Biesinger’s (Table S1-S4 in Supporting Information).38 The results indicate 48.7% Mn4+, 48.9% Mn3+ and 2.4% Mn2+ are present in the pristine LiMn2O4 (Figure 3b), which is in agreement with the valence distribution of Mn in typical LiMn2O4. After being charged to 4.3V (Figure 2c), Mn4+ is dominant in LixMn2O4 (70.9%). However, there is still about 22.1% Mn3+ and 7.0% Mn2+ present in the charged sample. Based on the charge capacity (117.3 mAh g-1), the lithium content is ca. 0.21 per formula unit after being charged to 4.3V, corresponding to 89.6% Mn4+ and 10.4% Mn3+ on average. That is to say, the Mn valence indicated by the XPS results is much lower than expected. As the XPS mainly reflect the surface information of the sample, this discrepancy suggests that the valence change of Mn ions near the surface differs from that in the bulk region. Figure 2d shows the XPS results of the sample discharged to 3.0V after being charged to 4.3V. The content of lithium was Li0.96Mn2O4 according to the discharge capacity of 111.9 mAh g-1. The valence distribution of Mn is close to the pristine LiMn2O4, and the slightly difference is most likely due to the disproportionation of minor Mn3+ into Mn2+ and Mn4+.



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Mn4+ 51.0% Mn3+ 45.5% Mn2+ 3.6%

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Figure 2. (a) Mn 2p spectra for LiMn2O4 at different states (pristine, charged to 4.3V, discharged to 3.0V); fitted spectra of (b) pristine LiMn2O4; (c) LiMn2O4 charged to 4.3V; (d) LiMn2O4 discharged to 3.0V after being charged to 4.3V.

Electron energy loss spectroscopy can be used to imprint the local electronic structure of specific ions. Accordingly, it was performed to probe the chemical states of manganese ions on the surface region of the pristine, charged and discharged LiMn2O4 particles (see Figure S2 in Supporting Information). Figure 3a shows the representative EELS spectra of O-K edge and Mn-L2,3 edge on the surface region at various (dis)charge states. For comparison, the Mn-L edges of the surface regions in the pristine, charged and discharged samples are normalized, as shown in Figure 3b. Results show that the intensity of pre-edge peak of O-K edge decreases significantly and the Mn-L3 edge exhibits a shift of ca. 1.8 eV towards lower energy after being charged to 4.3V, both of which indicate a decrease in Mn valence on the surface of the LiMn2O4 particle.39, 40 For the subsequent discharge, the pre-edge peak of O-K edge reappears and the Mn-L3 edge returns to the original position, implying that the Mn valence increases upon Li-ion insertion. The intensity ratio of L3/L2 is another indicator of Mn valence. For the pristine LiMn2O4 sample, the L3/L2 ratio was


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calculated to be ca. 2.1, suggesting a Mn valence slightly higher than +3.5; after being charged to 4.3V, the L3/L2 ratio in the surface of LiMn2O4 increased to 2.9, corresponding to a Mn valence between +2.0 and +2.5; after being discharged to 3.0V, the L3/L2 ratio was ca. 2.3 in the surface of LiMn2O4, which suggests a Mn valence between +3.0 and +3.5.29, 39 Note that the signal-to-noise ratio of the spectra was low, thus the obtained Mn valence might not be so accurate. However, the tendency of Mn valence changes during charge/discharge is qualitatively consistent with the indication from the shift of Mn-L3 peak. Above results suggest that the Mn valence on the particle surface decreases during charging process and increases when discharging, which agree well with the XPS results. However, this novel phenomenon is contrary to the common expectation on LiMn2O4, for which the valence of Mn ions should increase from +3.5 to +4.0 after charge and decrease from +4.0 to +3.5 after discharge.

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1.8 eV Mn-L3 Mn-L2 pristine charged to 4.3V

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Figure 3. Representative TEM-EELS spectra of O K-edge (a) and Mn L-edge (a and b) from the



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surface of pristine, charged and discharged LiMn2O4 samples. The pre-edge peaks of O-K edge show that it disappears on charge and appears on discharge, Mn L2,3 edges show a shifting to lower energy in the charged sample, as compared with the pristine and discharged samples.

Employing the aberration-corrected STEM technique41, 42 equipped with annular bright field (ABF) and high-angle annular dark field (HAADF) detectors, the change of crystallographic structure after (dis)charge could be identified directly at atomic scale.43 Figure 4a shows the crystal structure of spinel LiMn2O4 viewed along the [110] direction, which clearly exhibits the separated columns of Li, Mn and O atoms. As indicated by the arrows, two different Mn columns are assigned to Mn1 and Mn2, since the stacking density of Mn1 column is twice that of Mn2 column. The Mn diamond configuration is clearly seen in the HAADF image of LiMn2O4 (Figure 4b). This STEM-HAADF image is consistent with the crystal structure shown in Figure 4a and previously reported results.44 Close observation of the ABF image shown in Figure 4c, light element columns (e,g, O and Li) can also be detected without difficulty, and the corresponding line profile (Figure 4d) gives the relative position of Li, O, and Mn more conveniently. From the STEM images in Figure 4b and Figure 4c, it is confirmed that the atomic arrangement of the surface structure is identical to that of the bulk, indicating a homogeneous crystal structure from the bulk to the surface region of the pristine LiMn2O4 sample.


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Mn

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Figure 4. Crystal structure, HAADF and ABF images taken along the [110] crystallographic direction of LiMn2O4. (a) Demonstration of diamond structure of LiMn2O4. (b) HAADF image of the surface region of pristine LiMn2O4 sample. (c) The corresponding ABF image of the HAADF image in b. (d) Line profile corresponding to the dark blue line in c. Blue spheres represent Mn, red for O, green for Li and MnO6 octahedrons are blue. Li occupies 8a and Mn occupies the 16d site. Two different Mn columns are assigned: Mn1, Mn2, as they have different stacking density.

Figure 5 shows the HAADF images of LiMn2O4 sample charged to 4.1V (charge capacity is about 58 mAh/g) and 4.3V，respectively. As the contrast exhibits a Z1.7 dependency for HAADF imaging,45 light elements, such as Li and O, are invisible in the HAADF images; on the contrary, the contrast of the heavy element Mn is enhanced. In the sample charged to 4.1V (Figure 5a), a phase transition exists progressively from the bulk to the edge, as indicated by the green dotted line. The bulk region (enlarged from the red box) is similar to the pristine structure, while the surface region (enlarged from the yellow box) is quite different from the bulk region. The later shows an enhanced intensity at 8a sites of spinel structure, suggesting a Mn3O4 phase viewed along its [111] direction (see Supporting Information Figure S3). To the best of our knowledge, this is the first time that a phase transition at the surface area is observed at atomic scale in an electrochemically delithiated LiMn 2O4 sample.



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The transformation from spinel LiMn2O4 to Mn3O4 is dependent on the states of charge, as can be seen in sample charged to 4.3V (Figure 5b), the Mn3O4 area on surface enlarged significantly; besides, obvious contrast on the 16c site position appears at the subsurface region, suggesting the presence of rocksalt phase between the λ-MnO2 and Mn3O4 phase. The presence of Mn3O4 phase rationalize the abnormal increase in Mn valence at surface region after charge process: the average Mn valence is 8/3 in Mn3O4, even lower than that of pristine LiMn2O4.

(a)

(b)

Figure 5. HAADF images taken along the [110] zone axis of LiMn2O4 spinel cathode charged to 4.1V (a) and 4.3V (b). Magnified views of selected regions are shown in the right panels, where the contrast corresponding to the Mn columns at 16d and 8a sites are indicated by blue and orange


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spheres, respectively. The boundary between the bulk and surface regions is marked by the green dashed line.

Figure 6 shows the HAADF images taken from LiMn2O4 spinel structure after being discharged to 4V (discharge capacity ~ 55 mAh/g) and 3.0V, respectively. When discharged to 4V (Figure 6a), the rocksalt phase at the subsurface region disappears, and the amount of Mn3O4 at the surface region are significantly decreased. After fully discharged to 3.0V (Figure 6b), the Mn3O4 phase almost disappears and the STEM images show an identical surface structure to that of pristine sample. The results are in good agreement with the valence change detected by XPS and EELS, and further evidence that the anomalous Mn valence change at surface region originates from the appearance/disappearance of Mn3O4 phase during charge/discharge. It should be noted that the tetragonal Li2Mn2O4 may be formed in the discharged LiMn2O4 electrodes under nonequilibrium conditions, particularly at high current density.8 Carefully checking all the STEM images taken from the discharged sample, however, we found no evidence of the presence of Li2Mn2O4 phase. This phase can be easily distinguished from the LiMn2O4 phase by the contrast in the center of the diamond configuration due to the occupation of Li ions in the 16c site.



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(a)

(b)

Figure 6. HAADF image of LiMn2O4 spinel cathode discharged to 4.0V (a) and 3.0V (b). Magnified views of selected regions are shown in the right panels, where the contrast corresponding to the Mn columns at 16d and 8a sites are indicated by blue and orange spheres, respectively. The boundary between the bulk and surface regions is marked by the green dashed line.

The formation of Mn3O4 phase after delithiation of LiMn2O4 are supported by our density functional theory (DFT) based first-principles calculations. The delithiated state, i.e., λ-MnO2, is not energetically favorable compared to other polymorphs of MnO2 (Table S5). The metastable λ-MnO2 may decompose into more stable phases. Here we evaluate the thermodynamic driving force for the possible decomposition reactions of λ-MnO2, as follows: λ-MnO2 → (1/2)Mn2O3 + (1/4)O2

(1)


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λ-MnO2 → (1/3)Mn3O4 + (1/3)O2

(2)

λ-MnO2 → MnO + (1/2)O2

(3)

The formation Gibbs free energy per O2 for Eq. (1-3) can be written as Eq. (4-6), respectively. ∆G(Mn2O3) = ∆E + P∆V – T∆S ≈ 2E(Mn2O3) – 4E(λ-MnO2) + Ecorr(O2) – TS(O2)

(4)

∆G(Mn3O4) ≈ E(Mn3O4) – 3E(λ-MnO2) + Ecorr(O2) – TS(O2)

(5)

∆G(MnO) ≈ 2E(MnO) – 2E(λ-MnO2) + Ecorr(O2) – TS(O2)

(6)

where E(λ-MnO2), E(Mn2O3), E(Mn3O4), and E(MnO) refer to the calculated energies of λ-MnO2, Mn2O3, Mn3O4, and MnO, respectively. The effects of the P∆V and entropy in solids were omitted because they are negligible. Ecorr(O2) is the calculated energy of an isolated O2 molecule, taking account of the 1.36 eV per O2 correction to compensate for the GGA error in O2 (This correction also includes the P∆V resulting from O2 gas).46 The entropy of O2 gas is obtained from the JANAF thermochemistry tables.47 At room temperature, the calculated formation Gibbs free energy for reactions 4-6 are -506 meV, -10 meV, and 1162 meV, respectively, indicating that the reactions 4 and 5 are thermodynamically more favorable. Above calculation results suggest that the metastable λ-MnO2 is likely to transform into more stable Mn2O3 or Mn3O4, accompanied with a loss of oxygen. Although the experimentally observed Mn3O4 was localized on the particle surface, its thickness (up to 5nm, in Figure 5b) rationalizes the bulk phase approximation adopted in our calculations.



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Pristine

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Half charged

Fully charged LiMn2O4

Fully discharged λ-MnO2

Half discharged

Mn3O4

Figure 7. Schematic of phase evolution of a LiMn2O4 particle upon charge/discharge. The lithium concentration of LixMn2O4 is represented by the transparency of green colour where fully transparent green colour indicates λ-MnO2 and fully opaque green colour indicates LiMn2O4. The Mn3O4 phase is indicated by the orange colour. During the charge process, the Mn3O4 phase forms at the surface region and its amount reaches a maximum after being charged to 4.3V. The surface Mn3O4 phase decreases gradually on the subsequent discharge process, and it is almost negligible at the end of discharge.

Figure 7 summarizes the above experimental and computational results. It depicts the appearance of the Mn3O4 phase at the surface region upon charge, and then the gradual disapperance of the Mn3O4 phase during the subsequent discharge. The surface Mn3O4 phase seems to be reversible. It reaches a maximum amount at the end of the charge, and decreases to a negligible amount after the subsequent discharge. In contrast, previous studies showed that λ-MnO2 was relatively stable, based on differential scanning calorimetry (DSC) and X-ray diffraction measurements.27,

28

Likewise, our XRD measurements on the sample charged to 4.3V indicates no peaks corresponding to the Mn3O4 phase (Supporting Information Figure S4). A speculative explanation is that the structure transformation from λ-MnO2 to Mn3O4 merely occurs on the particle surface, and therefore it could hardly be detected by those global experimental measurements. Only by utilizing surface sensitive or high spatial experimental methods, as has been done in this work, can the surface structural information be accessed. As far as we know, this is the first time that a clear picture


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depicting the surface structural evolution of LiMn2O4 electrode has been obtained, which suggests that this traditional cathode material for Li-ion batteries is not as stable as expected, particularly in the surface region. The structural stability is crucial for practical electrode materials, where the Li-ions are reversibly accommodated in a host during charge/discharge. Normally, the cathode material is prone to structural degradation at charged (delithiated) state, which is highly oxidative and thus is unstable.48 Significant structural instability has been well recognized in Li-rich layered cathode material,49-55 which can be written as xLi2MnO3·(1-x)LiMO2 (M = transition metal). According to these studies, a layered to spinel phase transition is usually observed upon cycling. The structural instability of charged Li-rich material is not unexpected, as the extent of delithiation could readily exceed the theoretical limit, e.g., in Li2MnO3, the Mn4+ can not be further oxidized and therefore the theoretical capacity derived from Mn redox is zero. The over-delithiation will inevitably lead to the oxidation of oxygen ions. In most cases, the oxidized oxygen ions make the Li-rich cathode prone to oxygen loss and the consequent structural change,51, 56-59 even though recent studies suggested that the anionic O2-/O22- redox could be reversible to some extent.54,

60

For conventional

layered oxide materials, i.e., the ratio of lithium to transition metal ions is 1:1, surface structural degradation are also found in overcharged electrode,61, 62 although it is not so significant as that in Li-rich materials. In theory, the electron removal upon charging could be entirely compensated by oxidizing the transition metal redox in these materials, whereas the electrode decomposition is still possible, considering catalytic effect of organic electrolyte.63 Moreover, Hwang et al. argued that the extent of delithiation varies as a function of the distance from the electrode/electrolyte interface, and therefore, the phase transitions may occur more readily at the surface because there is a lower lithium content or higher number of lithium vacancies at the exterior of particle.62 They also found that transition metal ions at the particle surface are reduced when in the overcharged state (x = 0.1),62 similar to the case of LiMn2O4 electrode in this study. Presumably, an analogous mechanism may work for LiMn2O4



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cathode: at the charged state, the Li ions on the particle surface are almost depleted, even though ca. 0.2 Li per formula unit still remains in the whole electrode; the surface lithium deficiency results in an unstable surface structure that is vulnerable to lattice degradation. A lot of studies have been devoted to investigate the thermal stability of cathode materials at charged state, e.g., heating the delithiated electrode to a high temperature and then characterizing the consequent structural transformation.28, 64-74 These results could reflect the possible structural evolution pathway of a delithiated material. Generally, a surface phase transition from NaAlO2-type rhombohedral structure (R-3m) to spinel-type structure (Fd-3m), and finally to rocksalt-type structure (Fm-3m), is observed for layered cathode materials.65-73 Some of the studies demonstrated that the 8a tetrahedral sites in the spinel phase are partly occupied by transition metal ions.68, 69 Furthermore, the M3O4-type spinel, where the 8a sites are fully occupied by transition metal ions, was also found as an intermediate phase between

spinel

and

rocksalt

phase,

by

heating

an

overcharged

LixNi1/3Co1/3Mn1/3O2.70-73 In addition, Hu et al. investigated the decomposition pathway of spinel LixNi0.5Mn1.5O4 electrode and found that the decomposed product is NiMn2O4-like spinel, with 8a sites mostly occupied by Mn ions.74 All these findings indicate that the Mn3O4 could be formed as a result of degradation of LiMn2O4 electrode at delithiated state. In a practical Li-ion battery, the structural transformation may be promoted by the oxidizing organic electrolyte, although the temperature is not that high. The thickness of the Mn3O4 layer shows a gradual decrease over discharge, and consequently the surface structure is rather close to that of the pristine sample after being discharged to 3.0V. Although it seems that the Mn3O4 phase changes back to the LiMn2O4 phase during discharge, this reaction is less possible. As the ratio of oxygen to manganese is lower in Mn3O4 than that in LixMn2O4, a considerable amount of oxygen loss is expected during the formation of the Mn3O4 phase upon charge. It is unlikely that the lost oxygen could be restored on discharge.


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A more reasonable explanation is that the Mn3O4 phase decomposes during the discharge process (from 4.3V to 3.0V), and thus leading to the exposure of the LiMn2O4 phase underneath. The decomposition mechanism of Mn3O4 still remains unknown, but two reaction paths could be proposed. One possibility is that the surface Mn3O4 phase completely decomposed into the electrolyte. The other is that the Mn2+ ions situated in the tetrahedral sites of Mn3O4 dissolved into the electrolyte, and the remaining structure, being similar to λ-MnO2, converted to LiMn2O4 after Li insertion. In either case, the decrease of the surface Mn3O4 phase with discharge is not a reverse process of its formation during charge. According to the earlier studies, Mn3O4 is an anode material that is cycled between 0.1 to 3.0V.75-77 The discharge process follows the conversion reaction of Mn3O4 + 8Li+ + 8e- → 3Mn + 8Li2O, with the voltage plateau at ~0.4V. Nevertheless our results prove that Mn3O4 is involved in the discharge process at a voltage higher than 3.0V vs. Li+/Li. It may be due to the fact that the Mn3O4 phase on the particle surface is highly reactive. The presence of the Mn3O4 phase on the surface of LiMn2O4 particle may be closely related to the manganese dissolution upon cycling. Previous studies have shown that the rate of manganese dissolution is notably high at the charged state.6, 22-25 Based on aforementioned observations, the Mn3O4 phase is formed upon charging at the surface area, and its amount reaches a maximum at the end of the charge. It suggests that a positive correlation exists between the amount of the Mn3O4 phase and the rate of manganese dissolution. It is also worth noting that 1/3 of Mn ions in Mn3O4 are divalent, which has been proved to be the only soluble Mn ion in organic electrolyte.78 Therefore, we may reasonably infer that the presence of the Mn3O4 phase at the surface region contributes to the manganese dissolution and even plays a major role in this harmful reaction. In order to alleviate the manganese dissolution, appropriate strategies, e.g., cation substitution,24, 79 surface coating,80, 81 are needed to stabilize the surface structure and prevent the undesirable surface reconstruction.



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CONCLUSION In summary, we have combined surface sensitive (XPS) and high spatial resolution (EELS and STEM) experimental tools to investigate the chemical and structural evolution on the surface of LiMn2O4 electrode during electrochemical charge/discharge process. Contrary to the case in bulk LiMn2O4, the valence of surface Mn ions decreases on charging and increases on discharging. According to the atomically resolved STEM images, this anomalous phenomenon is attributed to the appearance and disappearance of Mn3O4 phase at charged and discharged states, respectively. Upon charging, LiMn2O4 electrode is delithiated and becomes unstable, resulting in oxygen release and Mn rearrangement at the surface area, which leads to the formation of Mn3O4 containing more Mn2+ ions. As Mn2+ ion is soluble, the surface structural transformation from delithiated LiMn2O4 to Mn3O4 may play a key role in the manganese dissolution reaction and the resultant loss in active material. Our work shed light on the importance of understanding the surface structural evolution of cathode material. It further suggests that modification and stabilization of the surface crystal structure, especially at charged state, are vital to develop satisfied cathode material for Li-ion batteries.

ASSOCIATED CONTENT Supporting Information Notes on XPS fitting and the parameters used for fitting, Rietveld refinement of XRD pattern, STEM images and 2D maps of the TEM-EELS line scan, simulated ADF image of Mn3O4, and the calculated energy of MnO2 with various structures.

AUTHOR INFORMATION Corresponding Authors * E-mail: [email protected] (X. Huang). * E-mail: [email protected] (L. Gu).

ACKNOWLEDGMENTs


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This work was supported by the National Basic Research Program of China (Grant No. 2013CB934002) and the "Strategic Priority Research Program" of the Chinese Academy of Sciences (Grant No. XDA01020304).

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Table of contents Li8a

Mn8a

Bulk

Mn16dO6 octahedron

Surface

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Pristine

Charged to 4.3V


Discharged to 3V

Surface Structure Evolution of LiMn2O4 Cathode Material upon

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