High-Performance Li(Li0.18Ni0.15Co0.15Mn0.52)O2@Li4M5O12

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Chemistry of Materials

High Performance Li(Li0.18Ni0.15Co0.15Mn0.52)O2@Li4M5O12 Heterostructured Cathode Material Coated with Lithium Borate Oxides Glass Layer

Xiaofei Biana, Qiang Fua, Hailong Qiua, Fei Dua, Yu Gaoa, Lijie Zhanga, Bo Zoub， Gang Chena,b, Yingjin Weia,*

a

Key Laboratory of Physics and Technology for Advanced Batteries (Ministry of

Education), College of Physics, Jilin University, Changchun 130012, P. R. China. b

State Key Laboratory of Superhard Materials, Jilin University, Changchun 130012,

P. R. China.

1

ACS Paragon Plus Environment


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ABSTRACT The continuous phase transformation to spinel LiMn2O4 seriously hinders the electrochemical properties of Li-excess layered oxides in lithium ion batteries. Herein, we prepared a heterostructured Li-excess layered cathode material consisting of a Li(Li0.18Ni0.15Co0.15Mn0.52)O2 active material in conjunction with a surface Li4M5O12 spinel and a Li2O-LiBO2-Li3BO3 glass coating layer. The material showed improved electrochemical kinetic properties with respect to its pristine counterpart because the Li2O-LiBO2-Li3BO3 glass layer not only improved the ionic conductivity of the material but also depressed the side reactions of the electrode with the electrolyte. In addition, the surface Li4M5O12 spinel constantly grew inwards the bulk of the material during long term charge-discharge cycling instead of the conventional LiMn2O4 transformation for the pristine Li(Li0.18Ni0.15Co0.15Mn0.52)O2. As a result, the heterostructured cathode material showed overall improved electrochemical performance. An initial discharge capacity of 258.8 mAh g-1 was obtained at the 0.2 C rate with remarkable capacity retention of 92.2 % after 100 cycles. Moreover, the material showed excellent rate capacity delivering a high discharge capacity of 130.4 mAh g-1and 100.4 mAh g-1 at the 10 C and 20 C rates, respectively. Differential scanning calorimetry showed that the exothermic temperature of the fully charged electrode was elevated to 324.2 oC with little thermal release of 232.5 J g-1 demonstrating good thermal safety of the material.

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1. INTRODUCTION Li-excess layered oxides, Li[LixM1-x]O2 or xLi2MnO3·(1-x)LiMO2 (M = Mn, Ni, Co), have been one of the most promising cathode materials for the next generation of lithium ion batteries due to their high discharge capacities of more than 250 mAh g-1 in the voltage window of 2.0-4.8 V vs. Li+/Li. It is well known that the high reversible capacity of Li-excess layered oxides is associated with an initial loss of oxygen from the material lattice above 4.5 V.1, 2 This voltage is too high for most LiPF6-based electrolytes to maintain thermodynamically stable. Thus the Li+ extraction/insertion in Li-excess layered oxides is always accompanied by complex interfacial phenomena. Side reactions and growth of a solid electrolyte interface (SEI) film on Li-excess layered oxides upon cycling have been verified by different groups.3-5 For example, Komaba et al. reported that the oxygen containing species of SEI would partly decompose during cycling thus contributing some reversible capacity for Li-excess layered oxides.3 But, the poor electronic conductivity of SEI film may hinder the transportation of electrons which could result in poor rate capability of the electrode. In addition, the loss of oxygen may lead to a phase transition from the layered structure to a defect-spinel structure at the surface region of the materials, resulting in a large initial irreversible capacity.6, 7 On the other hand, the bulk region of Li-excess layered oxides is also unstable during cycling. It has shown that the originally well integrated LiMO2 and Li2MnO3 layered structure slowly transforms to a nano-composite where spinel LiMn2O4-like regions are embedded in a layered LiMO2 framework.8-13 As a result, the average discharge voltage gradually decreases leading to reduction in the specific energy of the battery. Furthermore, continuous capacity fading is always observed as a result of Mn dissolution and Jahn-Teller distortion of the Mn3+ ions. In such a LiMn2O4/LiMO2 nano-composite, the Li ions have to 3



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crossover numerous phase boundaries during electrochemical processes and transport in different diffusion channels thus resulting in very low lithium diffusion coefficients. It has reported that the lithium diffusion coefficients of Li-excess layered oxides during oxygen loss process are only 10-19 cm2 s-1. Even though the lithium diffusion coefficients increase to 10-15 ~ 10-14 cm2 s-1 in the following discharge, they are still much lower than those of traditional layered cathode materials such as LiCoO2 and LiNi1/3Co1/3Mn1/3O2.14, 15 Surface coating has been an effective method to overcome the above problems of Li-excess layered oxides. The most commonly used coating materials are metal oxides such as TiO2, ZnO, Al2O3 and MgO,16-19 or metal fluorites such as CaF2 and AlF3.20,

21

These coating species were found to effectively depress the electrolyte

reduction and formation of SEI film. Besides these, carbonaceous materials,22, 23 and conductive polymers24, 25 have been used to increase the electronic conductivity of Li-excess layered oxides. Recently, metal phosphates including Li3PO4 and Li-Ni-PO4 have also been studied as coating materials for Li-excess layered oxides.26, 27 The inert chemical properties of metal phosphates were effective in preventing side reactions between the cathode and the electrolyte. However, most of these works trend to form an "isolative" layer on the particle surface which hardly affects the surface structure of the materials. Therefore, some shortcomings of Li-excess layered oxides such as gradual voltage decay and capacity fading were not substantially resolved. Alternately, recent proceedings of surface treatment have shown applausive effects on the electrochemical properties of Li-excess layered oxides. It has reported that acidic treatment of Li-excess layered oxide using HNO3 could cause proton exchange between H+ and Li+, resulting in Li vacancies and spinel transformation in the surface region of the material after a post-annealing process.28 This could eliminate the initial 4


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capacity loss but the strong acidic solution may destroy the lattice structure of the material thus damaging its cycle stability and rate capability. Surface treatment of Li-excess layered oxides by reducing agents has been proposed recently. Cho et al. showed

that

a

NiO

cubic

phase

was

formed

on

the

surface

of

0.4Li2MnO3·0.6LiNi1/3Co1/3Mn1/3O2 with hydrazine treatment.23 This NiO layer can act as a surface stabilizer; but too thick of such a surface layer would depress the rate capability of the material due to the poor conductivity of NiO. Song et al. modified the surface of Li(Li0.2Mn0.54Ni0.13Co0.13)O2 with carbon black.22 A phase transformation from Li2MnO3 to LiMn2O4 was observed at the surface region resulting from reduction of Mn4+ to Mn3+, as well as migration of Li to the tetrahedral 8a site and Mn to the octahedral 16d site. Similar phase transformation was also obtained by AlF3 coating as reported by Sun et al. in an earlier time.29 The surface LiMn2O4 layer with good electronic and ionic conductivities improved the rate capability of the material. However, the materials bulk still transformed to a LiMn2O4-like phase after several tens of cycles, resulting in continuous capacity fading and voltage decay. Taking account of their distinctive electrochemical properties, following concerns should be carefully addressed in order to overall improve the electrochemical properties of Li-excess layered oxides: i) the surface of the material should be chemically and electrochemically inert to avoid harmful side reactions with the electrolyte; ii) the material should have good electrochemical kinetic properties to allow effective charge transfer reactions and fast Li ion diffusion; iii) the LiMn2O4-like phase transformation should be effectively depressed to keep stable discharge voltage and good cycle stability. In this work, we design a novel heterostructured LLMO@Li4M5O12@LBO material to achieve these aims all together. 5



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The structure is consisted of a Li(Li0.18Ni0.15Co0.15Mn0.52)O2 (LLMO, M = Ni, Co, Mn) active material, a Li4M5O12 surface spinel and a Li2O-LiBO2-Li3BO3 (LBO) glass coating layer. Based on careful analysis, we found that the LiMn2O4-like transformation that occurred in traditional Li-excess layered oxidesis replaced by a Li4M5O12 transformation using this heterostructured material. In addition, the LBO surface layer not only acts as a fast Li ion transport media but also protects the active material from harmful attacks by the electrolyte. As a result, the material shows excellent electrochemical performance including high rate capability, slight voltage decay, stable capacity retention and good thermal safety. 2. EXPERIMENTAL SECTION The pristine Li(Li0.18Ni0.15Co0.15Mn0.52)O2 material was prepared by the sol-gel method using Li2CO3, Co(CH3COO)2, Ni(CH3COO)2 and Mn(CH3COO)2 (Sinopharm Chemical Reagent, analytical grade) as the starting materials. The starting materials were weighted in a stoichiometric ratio of Li: Co: Ni: Mn = 1.18: 0.15: 0.15: 0.52 and dissolved in de-ionized water with vigorous stirring. Then citric acid was added into the solution with a molar ratio of Li: citric acid = 1: 2. The pH value of the solution was adjusted at 7.3 by adding ammonia water dropwise. After stirring at 50 °C for 15 h, the viscous solution was dried at 120 °C in vacuum oven. The obtained substance was pre-treated at 450 °C for 5 h to eliminate the organic contents. Then, the precursor was pressed into pellet and annealed at 900 °C for 12 h followed by quenching in liquid nitrogen. For preparation of the LLMO@Li4M5O12@LBO material, 0.4 g of LLMO powder was mixed in 50 mL of 0.06 mol L-1 LiOHsolution with vigorous stirring for 3 h. Then, 10 mL of 0.1 mol L-1 H3BO3 dilute solution was dropwise added into the suspension, followed by stirring at 80 oC for 24 h to 6


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evaporate the excessive water. The resulting moist powder was dried at 120 oC and then heat treated at 450 oC for 3 h in air thus obtaining the LLMO@Li4M5O12@LBO material. The crystal structure of the material was studied by X-ray diffraction (XRD) on a Rigaku AXS D8 diffractometer with Cu Kα radiation. The lattice parameterswere calculated by the Celref 3 Program. The morphology of the materials was studied bya JSM-6700F field emission scanning electron microscope (FESEM). High resolution transmission electron microscope (HRTEM) was performed ona FEI Tecnai G2 F20 S-TWIN. Raman scattering was collected on a ThermoScientific FT-Raman using Nd-line laser source (λ=532 nm). X-ray photoelectron spectroscopy (XPS) was performed on an ESCALAB spectrometer using Mg-Kα light source. All binding energies were calibrated using the C 1s peak at 284.6 eV. Fourier transform infrared spectroscopy (FTIR) was measured on a Thermo Scientific Nicolet 6700 spectrometer using the KBr disk method. The electronic and ionic conductivities of the samples were

studied

by

ac

conductivity

measurement

on

a

Bio-Logic

VSP

potentiostatic-galvanostatic system. Differential scanning calorimetry (DSC) of the fully charged electrode was carried out on a TA-Q2000 thermal analyser at a heating speed of 10 °C min-1 without removing the residual electrolyte. Electrochemical experiments of the materials were performed on 2032-type coin cells using metallic Li as the anode electrode. The cathode electrode was composed of 75 wt.% active material, 15 wt.% Super P conductive additive and 10 wt.% poly-vinylidenefluoride (PVDF) binder dissolved in N-methyl-2-pyrrolidone (NMP). The slurry mixture was coated on an Al current collector and dried overnight in vacuum oven. The electrode was cut into 8 × 8 mm2 for use. The cathode and anode electrodes were separated by Celgard 2320 membrane. The electrolyte was 1 mol·L-1 7



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lithium hexafluoro-phosphate (LiPF6) dissolved in ethylene carbonate (EC), dimethyl carbonate (DMC) and ethylmethyl carbonate (EMC) (EC: DMC: EMC = 1: 1: 8, by v/v ratio). Galvanostatic charge-discharge experiments were performed on a Land-2100 automatic battery tester. Electrochemical impedance spectroscopy (EIS) was studied on a Bio-Logic VSP multichannel potentiostaticgalvanostatic system by applying an ac voltage of 5 mV. Galvanostatic intermittent titration technique (GITT) was performed throughout the first discharge on the above VSP instrument. For each GITT step, the battery was discharged with a current flux of 30 mA g-1 for 0.5 h, followed by an open circuit stand for 4 h to reach the quasi-equilibrium state.

3. RESULTS AND DISCUSSION 3.1. Structural and Physical Properties The morphological features of the materials were studied by SEM (Supporting Information, Figure S1). The pristine LLMO exhibits sphere-like shape and clean surface. The particles of the material are not uniform in size but most of them are in the range of 200 ~ 500 nm. LLMO@Li4M5O12@LBO shows similar particle shape and size as those of LLMO but its particles are rougher than the pristine one. The crystal structures of LLMO and LLMO@Li4M5O12@LBO were studied by XRD as shown in Figure 1. LLMO shows a typical XRD pattern of Li-excess layered oxides. The main diffraction peaks (marked by “R” subscripts) can be indexed to the rhombohedral LiM'O2 (M' = Li0.18Ni0.15Co0.15Mn0.52) phase with R-3m symmetry. The weak peaks (marked by “M” subscripts) in the 2θ range of 20o ~ 25o are due to the LiMn6 superstructure of the monoclinic Li2MnO3 phase with C2/m symmetry. The XRD pattern of LLMO@Li4M5O12@LBO is almost the same as that of LLMO. But 8


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the (003)R peak slightly shifts from 18.70o to 18.74o. In addition, a shoulder peak appears on the left of the (101)R peak. These features suggest the formation of a spinel-like phase in LLMO as reported in literatures,22,

30

but are not conclusive

enough to enable unambiguous identification of the structure changes. The lattice parameters of the materials are calculated based on the R-3m symmetry, which are a = 2.8531 Å, c = 14.2420 Å for LLMO and a = 2.8538 Å, c = 14.2413 Å for LLMO@Li4M5O12@LBO, respectively.

Figure 1. XRD patterns of the LLMO and LLMO@Li4M5O12@LBO samples.

FTIR was used to study the phase compositions of the samples as shown in Figure 2. The pristine LLMO shows two strong absorption peaks at 540 and 613 cm-1 which are assigned to the asymmetrical stretching and bending vibrations of the M-O bond, respectively. The weak peaks at 872, 1438 and 1490 cm-1 are attributed to Li2CO3 and the peak at 1635 cm-1 is due to the adsorbed water. Li2CO3 is often found as an impurity in layered cathode materials but could not be detected by XRD because of its amorphous state and very little amount in the product. All these FTIR features are reserved well for LLMO@Li4M5O12@LBO. In addition, a series of small peaks are observed in the range of 850 ~ 1600 cm-1 which could be attributed to two 9



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different B-O local structures, i.e. the tetrahedral BO4 group of LiBO2 (850 ~ 1100 cm-1) and the trigonal BO3 group of Li3BO3 (1200 ~ 1600 cm-1).31 Thus, it is believed that a Li2O-LiBO2-Li3BO3 (LBO) glass phase is formed in the material. Note that the Li2O component of the LBO glass should originate from the decomposition of LiOH during post-annealing but shows no FTIR peaks in the experimental wavenumber range. In addition, this glass phase is not detected by XRD because it is only about 0.5 wt.% of the total sample weight assuming that the LiOH and H3BO3 agents were completely consumed in the preparation process.

Figure 2. FTIR patterns of the LLMO and LLMO@Li4M5O12@LBO samples.

The microstructures of the materials were studied by HRTEM as shown in Figure 3. Well-ordered atomic columns could be clearly observed for the pristine LLMO. These atomic columns are separated by a distance of 4.8 Å which fits well with the (002) planes of the C2/m layered structure. The C2/m structure could be further verified by the fast Fourier transform (FFT) pattern. This C2/m structure is observed over the whole part of the material indicating that a pure layered phase is built up throughout the sample. LLMO@Li4M5O12@LBO shows a different 10


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microstructure from that of LLMO. At the outer surface of the particle, the non-crystalline phase with thickness about 2 nm is assigned to the LBO glass layer. The bulk of the material also shows the (002) planes of the C2/m layered structure. But, a newly-formed crystalline phase with thickness about 4 nm could be observed on the skin boundary of the particle which is intimately integrated witht the C2/m main phase. The selected regions marked by orange squares reveal a cubic spinel structure with Fd3m symmetry based on FFT analysis. Note that the zone axis applied for FFT analysis are [010]layered and [01-1]spinel, respectively.

Figure 3. HRTEM images of the (a)LLMO and (b)LLMO@Li4M5O12@LBO samples.

Raman scattering which is sensitive to the local structure of materials further confirms the spinel phase in LLMO@Li4M5O12@LBO. As shown in Figure 4, the pristine LLMO displays several Raman peaks at 243, 324, 368, 434, 486 and 599 cm-1, most of which are consistent with those of monoclinic Li2MnO3.32 According to Julien et al,33 the weak peak at 243 cm-1 is due to the LiO6 octahedral of layered LLMO, and the peaks above 300 cm-1 are due to the MO6 octahedral of LLMO. For the Raman pattern of LLMO@Li4M5O12@LBO, the LiO6 vibration becomes almost 11



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invisible indicating that the amount of six-fold coordinated lithium is reduced. On the other hand, the appearance of a new peak at 407 cm-1 indicates the formation of four-fold coordinated LiO4 which probably arises from a spinel-type structure.34 The two strongest peaks at 486 and 599 cm-1 remain stable for both samples. But a shoulder peak is observed at 651 cm-1 only for LLMO@Li4M5O12@LBO. This shoulder peak is a typical feature of spinel Li4Mn5O12.32 Note that the corresponding peak for another type of Li-Mn-O spinel phase, LiMn2O4, is always observed at a lower wavenumber of 625 cm-1.33 Therefore, it is believed that a spinel Li4M5O12-like phase is formed in LLMO@Li4M5O12@LBO. This is different from the previous literatures reporting that a LiMn2O4-like phase would be formed in Li-excess layered oxides after surface treatment with metal fluorites such as AlF3 or reducing agents such as active carbon.22, 29

Figure 4. Raman patterns of the LLMO and LLMO@Li4M5O12@LBO samples.

The XPS results presented in Figure 5 provide relevant information on the surface chemistry of the samples in a depth of several nanometers. The B 1s XPS of the LLMO@Li4M5O12@LBO sample is located at 191.7 eV which is due to the B-O 12


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bond of LBO. The Ni 2p1/2 (854.6 eV), Co 2p1/2 (780.3 eV) and Mn 2p1/2 (642.4 eV) binding energies of LLMO correspond well to those of Ni2+, Co3+ and Mn4+, respectively.35 The XPS signals for LLMO@Li4M5O12@LBO are much weaker than those of LLMO due to the LBO surface layer. All of the Ni, Co and Mn elements are detected by XPS indicating that the Li4M5O12-like phase is a Ni/Co/Mn-contained spinel rather than accurate Li4Mn5O12. XPS analysis shows that the relative atomic ratio of Ni: Co: Mn is 1: 1: 2.5. This indicates that the major transition metal in the surface Li4M5O12 heterostructure is Mn. The binding energies for Ni 2p1/2, Co 2p1/2 and Mn 2p1/2 show no changes with respect to those of the pristine LLMO. This indicates that the oxidation states of the transitional metals of the Li4M5O12 surface spinel are consistent with those of the material bulk. Therefore, it is suggested that only atomic migrations took place at the surface region during the acid treatment and post-annealing process.

Figure 5. B 1s, Ni 2p, Co 2pand Mn 2pXPS patterns of the LLMO and LLMO@Li4M5O12@LBO samples.

The formation mechanisms of the heterostructuredLLMO@Li4M5O12@LBO material could be illustrated in Figure 6. When the LLMO powder was dispersed in the H3BO3/LiOH solution, a part of the Li ions at the surface region would be replaced by protons due to the Li/H exchange reaction. As a result, the dry material is 13



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composed of Li1-xHxM'O2 with Li3BO3 and LiOH deposited on the surface of the particles. It is suggested that the Li/H exchange reaction would only take place in the lithium layer of LLMO because the energy of removing one Li from the lithium layer is much lower than that from the transitional metal (TM) layer.36 During the post-annealing process, the protons were extracted from the lithium layer accompanied by releasing of O2- (2H+ + O2- = H2O) thus leaving vacancies in both the lithium layer and the oxygen layer. The [M'O6] octahedral of LLMO is edge-sharing with three [LiO6] octahedral. When the protons were extracted, the adjacent Li ions would migrate from the octahedral site to the tetrahedral site, face-sharing with the [M'O6] octahedral. In the meanwhile, the M' ions have a tendency to migrate to the octahedral vacancies, thus forming a Li4M5O12 spinel ordering at the surface region where the 8a tetrahedral site is occupied by Li ions and the 16d octahedral site is occupied by both Li ions and M ions (i.e. the M' ions). At the outer surface region, a part of the BO3 groups transform to BO4 during the post-annealing process. In the meanwhile, LiOH was decomposed to Li2O thus forming the Li2O-LiBO2-Li3BO3 glass layer.

Figure 6. Formation mechanisms of the heterostructuredLLMO@Li4M5O12@LBO.

3.2 Galvanostatic Charge-discharge Cycling 14


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Figure 7 shows the first charge-discharge curves of the samples at the 0.2 C rate (1C = 200 mA g-1). The first charge curve is composed of a slope region below 4.5 V, followed by a voltage plateau centered at 4.55 V. The slope region is attributed to Li+ extraction from the lithium layer and the voltage plateau is due to the concurrent Li+ extraction and oxygen loss. The specific capacities of the slope region are similar for both samples. But the voltage plateau of LLMO@Li4M5O12@LBO is a little bit shorter than that of LLMO. This could be due to the initial oxygen loss during the post-annealing process which is accompanied by the release of H+ (2H+ + O2- = H2O). The first discharge capacity of LLMO@Li4M5O12@LBO is 258.8 mAh g-1 which is larger than that of 246.7 mAh g-1 for the pristine LLMO. The excessive discharge capacity could be attributed to the additional capacity below 3.2 V which is due to the activation of the Mn4+/Mn3+ redox reaction. As a result, the initial columbic efficiency of the material increases from 75.4 % to 85.8 %.

Figure 7. First charge-discharge curves of the LLMO and LLMO@Li4M5O12@LBO samples at the 0.2 C rate.

Figure 8a shows the cycling performance of the samples at the 0.2 C rate. The discharge capacity of LLMO decreases from 246.7 mAh g-1 to 160.9 mAh g-1 after 15



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100 cycles, corresponding to capacity retention of 65.2 %. On the contrary, LLMO@Li4M5O12@LBO shows significantly improved cycle stability. A much larger discharge capacity of 238.6mAh g-1 can be obtained after 100 cycles, resulting in capacity retention of 92.2 %. The long term cycle stability of the samples is studied at the 10 C rate as shown in Figure 8b. LLMO shows poor capacity retention at this high current rate which rapidly decreases to 15.0 mAh g-1 after 200 cycles, while LLMO@Li4M5O12@LBO could still deliever 108.6 mAh g-1 after 300 cycles. One of the big shortcomings of Li-excess layered cathodes is continuous voltage decay during cycling resulting from the phase transformation from the layered structure to a spinel-like structure. Figure 8c shows that the average discharge voltage of LLMO decreases from 3.66 V to 2.82 V after 100 cycles, while that of LLMO@Li4M5O12@LBO only decreases to 3.12 V. As a synergetic effect of slower capacity fading and slighter voltage decay, LLMO@Li4M5O12@LBO shows much better energy reserve ability than LLMO (Supporting Information, Figure S2). The specific energy of LLMO after 100 cycles is 477.9 Wh kg-1 comparing to the 761.7 Wh kg-1 of LLMO@Li4M5O12@LBO. Figure 8d shows the rate performance of the samples with the current rate increasing from 0.2 C to 20 C. At each current rate, LLMO@Li4M5O12@LBO always shows a larger discharge capacity than that of LLMO. For example, the discharge capacities of LLMO@Li4M5O12@LBO at the 1 C and 10 C rates are 212.4 and 130.4 mAh g-1, respectively, corresponding to 82.1 % and 50.4 % of that at the 0.2 C rate. In contrast, the discharge capacities of pristine LLMO at the 1 C and 10 C rates are only 183.6 and 80.5 mAh g-1, respectively, corresponding to 74.4 % and 32.6 % of that at the 0.2 C rate. More encouragingly, LLMO@Li4M5O12@LBO could still deliver a discharge

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capacity of 100.4 mAh g-1 at the 20 C rate while the pristine LLMO shows almost no available capacity at this high current rate.

Figure 8. Cycling performance of the LLMO and LLMO@Li4M5O12@LBO samples at the (a) 0.2 C rate and (b) 10 C rate; (c) average discharge voltage of the samples at the 0.2 C rate; (d) rate dependent cycling performance of the samples.

3.3 Phase Transformation During Charge-discharge Cycling Figure 9a shows the representative discharge curves of the samples at the 0.2 C rate. The discharge curves of LLMO show an IR drop from 4.8 V to 4.6 V in the initial two cycles. Then the discharge voltage smoothly decreases to 3.0 V followed by anabrupt concentration polarization to 2.0 V. The IR drop gradually decreases to 4.2 V after 100 cycles indicating increased internal resistance of the cell. In the meanwhile, the discharge curve continuously declines along with the appearance of a new voltage plateau in the voltage window of 3.0-2.4 V. On the contrary, LLMO@Li4M5O12@LBO shows an additional voltage slope below 3.2 V in the initial two cycles implying different electrochemical processes occurring in the material. The discharge curve declines much slower than that of LLMO; and the IR drop shows almost no change in 100 cycles. In addition, the newly formed voltage plateau is 17



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observed in the voltage window of 3.3-2.7 V, which is about 0.3 V higher than that of LLMO. All these differences indicate improved electrochemical performance of LLMO@Li4M5O12@LBO with respect to that of the pristine LLMO.

Figure 9. (a) Representative discharge curves and (b) corresponding dQ/dV profiles of the LLMO and LLMO@Li4M5O12@LBO samples at the 0.2 C rate.

The changes in discharge curves can be clearly reflected from their dQ/dV profiles as shown in Figure 9b. LLMO shows P1, P2 and P3 peaks in the first discharge which are due to Li interaction into the C2/m layered structure, relating with the reduction of Co4+, Ni4+ and Mn4+, respectively.22 It is noticed that the intensities of P1, P2 and P3 peaks decrease with cycling. In the meanwhile, the P3 peak gradually shifts to 2.75 V which is finally replaced by P5 after 50 cycles. Afterwards, the P5 peak increases slowly in following cycles. These changes correspond to the phase transformation of LLMO from the C2/m layered structure to a LiMn2O4-like spinel structure.37 One of the bad consequences of this phase transformation is the continuous voltage decay as shown in Figure 8c. In addition, the gradual capacity fading of LLMO could be partly attributed to this phase transformation because of the dissolution of Mn and the Jahn-Teller distortion of [MnO6] octahedral which are majorly caused by the Mn3+ ions of LiMn2O4. For the LLMO@Li4M5O12@LBO sample, the P5 peak due to the LiMn2O4-like phase immediately appears in the first discharge. In addition, a new 18


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peak labeled as P4 is observed at 2.91 V which is characteristic of spinel Li4M5O12.38 One can see that the P5 peak gradually decreases from the second cycle and then completely vanishes after 30 cycles. In the meanwhile, the P4 peak becomes stronger and stronger with cycling. This indicates that the amount of LiMn2O4-like phase is gradually reduced with charge-discharge cycling while the Li4M5O12 phase keeps growing. As a cathode material, Li4M5O12 shows much better cycle stability than LiMn2O4 because the Mn4+ ions in Li4M5O12 results in minimal [MnO6] Jahn-Teller distortion

and

Mn

dissolution.

Therefore,

the

capacity

fading

of

LLMO@Li4M5O12@LBO becomes slower and slower as illustrated in Figure 8a. In addition, it is noticed that the P1, P2 and P3 peaks of LLMO@Li4M5O12@LBO after 100 cycles are much stronger than those of LLMO. This indicates that the C2/m layered structure could be well reserved in the heterostructured material. In order to confirm the phase transformation of the materials during charge-discharge cycling, the battery cell was disassembled after 140 cycles. The electrode composite was scraped off the electrode and completely washed by DMC to remove any absorbance.

Figure 10. TEM and HRTEM images of the (a, b, c) LLMO and (d, e, f) LLMO@Li4M5O12@LBO samples after 140 charge-discharge cycles.

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HRTEM analysis of multiple particles for the pristine LLMO shows that not only the particle surface (Figure 10b) but also the material bulk (Figure 10c) match well with the [010] zone axis of a lithiated LiMn2O4-like spinel phase. In addition, apparent defects (marked by white squares) can be seen between adjacent spinel regions which could be attributed to an intermediate structure from the C2/m layered structure to the Fd3m spinel structure. Different from that observed for LLMO, the surface region of the cycled LLMO@Li4M5O12@LBO could be definitely indexed to a lithiated Li4M5O12 spinel phase as confirmed by its distinctive (400) and (220) diffractions along the [001] zone axis (Figure 10e). However, many C2/m layered regions are still observed in the material bulk (Figure 10f), indicating that the C2/m layered structure could be well reserved in the LLMO@Li4M5O12@LBO sample which is consistent with the dQ/dV analysis. The Li4M5O12 transformation for the heterostructured LLMO@Li4M5O12@LBO is quite different from the LiMn2O4-like phase transformation observed for the pristine LLMO. This Li4M5O12 phase is demonstrated to be closely related with the excellent capacity retention and slight voltage decay of LLMO@Li4M5O12@LBO. Regarding the mechanisms of this unique phase transition, herein we propose a "swallowing" effect as illustrated in Figure 11.

20


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Figure 11. Illustration of the phase transformation of LLMO@Li4M5O12@LBO during charge-discharge cycling.

First, it is believed that the LiMn2O4-like phase formed in the first discharge is originally from the LLMO phase therefore must be underneath the Li4M5O12 surface spinel. When the cell was discharged to 2.75 V, the Li ions would intercalate into the 16c site of the LiMn2O4-like phase via the lithiated Li4M5O12. But, a part of the Li(16d) ions of lithiated Li4M5O12 could easily migrate into the 16d octahedral site of LiMn2O4 due to a low energy barrier between these two 16d sites. This process looks like that a part of the LiMn2O4 phase is "swallowed" by Li4M5O12. In the following cycle, the vacant Li(16d) site of Li4M5O12 could be filled again by Li ions when the cell was discharged to 2.91 V. Therefore, the amount of the LiMn2O4-like phase is gradually reduced with charge-discharge cycling while the Li4M5O12 phase keeps growing. From the dQ/dV profiles it is seen that this "swallowing" process may continue for about 30 cycles. Afterwards, the material shows almost no capacity fading during long term cycling. It should be noted that even though the discharge 21



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voltage of Li4M5O12 is higher than that of LiMn2O4, it is still lower than that of the LLMO phase. Therefore, the discharge voltage of LLMO@Li4M5O12@LBO also decreases with charge-discharge cycling. However, the Li4M5O12 transformation of LLMO@Li4M5O12@LBO is slower than the LiMn2O4 transformation of LLMO. Therefore, LLMO@Li4M5O12@LBO shows slower voltage decay than that of LLMO (Figure 8c), resulting in excellent energy reserve ability. 3.4 Electrochemical Kinetic Properties EIS analysis was performed to deeply understand the improved electrochemical performance of LLMO@Li4M5O12@LBO. The Nquist plots of the electrodes after 20 and 50 cycles are displayed in Figure 12. The two semicircles in the high-to-medium frequency region are due to the SEI film and charge transfer process, respectively. The slope line in the low frequency region is due to Li+ diffusion in the electrode. The SEI resistance (Rf) for LLMO increases from 83.8 Ω to 147.4 Ω which indicates that the side reactions in the electrode are unstoppable during cycling. While the small and stable SEI resistance for LLMO@Li4M5O12@LBO indicates minimal side reactions occurring in the electrode. The surface morphologies of the materials after 140 cycles can be seen in Figure 10a and 10d. TEM shows that the surface of LLMO@Li4M5O12@LBO is much smoother and cleaner than that of the pristine LLMO. Under HRTEM, a thick SEI film (10 ~ 15 nm) could be seen on the particle surface of LLMO owing to aggressive side reactions on the electrode occurring during each charge to high voltages. Accumulation of such a thick SEI film is one of the major factors leading to the large cell polarizations, poor capacity retention and low rate capability of the LLMO material. In contrast, a very thin film (2 ~ 4 nm) can be seen for LLMO@Li4M5O12@LBO which is only a little bit thicker than its original 22


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LBO layer. Comparing to the highly oxidative transitional metal ions of the charged LLMO electrode (especially the Co4+ ions), the LBO surface layer of the charged LLMO@Li4M5O12@LBO is inert in electrolyte which effectively depresses the side reactions of the electrode thus resulting in a thin SEI film.

Figure 12. Nquist plots of the LLMO and LLMO@Li4M5O12@LBO samples after (a) 20 cycles and (b)50 cycles.

LLMO@Li4M5O12@LBO shows a charge transfer resistance (Rct) of 215.5 Ω which is much smaller than that of LLMO (1136.5 Ω). The improved charge transfer of LLMO@Li4M5O12@LBO is closely related to its heterostructure. First, the LBO glass layer is a good Li ion conductor which could improve the ionic conductivity of the material. AC conductivity measurement (Supporting Information, Figure S3) shows that the ionic conductivity of LLMO@Li4M5O12@LBO (1.18 × 10-4 S cm-1) is 18 times larger than that of the pristine LLMO (6.44 × 10-6 S cm-1). In addition, the electronic conductivity of LLMO@Li4M5O12@LBO (1.39 × 10-7 S cm-1) is also larger than that of LLMO (1.46 × 10-8 S cm-1). The inferior electronic conductivity of Li-excess layered oxides has been attributed to the insulate property of the 23



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Li2MnO3component.39 In contrast to the pristine LLMO, the Li4M5O12 phase of LLMO@Li4M5O12@LBO reduces the amount of Li2MnO3 at least at the surface region thus resulting in higher electronic conductivity. It is believed that

the larger

ionic conductivity and electronic conductivity of LLMO@Li4M5O12@LBO facilitate the charge transfer reactions at the electrode/electrolyte interface. Further, the lithium diffusion coefficients (DLi) of the samples during the first discharge are determined by GITT (Supporting Information，Figure S4). As shown in Figure 13, the lithium diffusion coefficients of LLMO are in the range of 10-17~ 10-15 cm2 s-1 which are similar to those of previous report.14 The lithium diffusion coefficients of LLMO@Li4M5O12@LBO, on the contrary, are overall larger than those of LLMO. The improved lithium diffusion coefficients can be partly attributed to the surface spinel phase of which the 3D diffusion pathway is much favorable for Li ion diffusion. In addition, the LBO glass layer with high Li ion conductivity can lower the ionic conductive barrier between the active material and the electrolyte.

Figure 13. Lithium diffusion coefficients of the LLMO and LLMO@Li4M5O12@LBO samples during the first discharge process.

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We also studied the electrochemical performance of the LLMO@Li4M5O12@LBO samples with different LBO amounts (Supporting Information, Figure S5). It is seen that the present material with 0.5 wt.% LBO shows the best electrochemical performance in all samples. Detailed analysis of different LLMO@Li4M5O12@LBO samples is beyond the scope of this work. However, the above preliminary study shows that a small amount of LBO is enough to significantly improve the electrochemical performance of LLMO. Obviously, the thickness of the surface Li4M5O12 heterostructure also plays important roles on the electrochemical properties of LLMO. The thickness of Li4M5O12 could be controlled by the concentration (pH value) of H3BO3 or using different amounts of H3BO3 solution to soak the material. This problem will also be studied in detail in our future work.

3.5 Thermal Safety Properties The thermal stability of LLMO and LLMO@Li4M5O12@LBO was studied by DSC after the first charge (4.8V) without removing the electrolyte. As shown in Figure 14, the pristine LLMO shows a big exothermic peak at 203.0 oC and another small exothermic peak at 267.9 oC, resulting in total thermal release of 832.4 J g-1. In contrast, LLMO@Li4M5O12@LBO only shows a mild exothermic peak. Not only the onset temperature increases to 324.2 oC but also the total thermal release is significantly reduced to 232.5 J g-1. The up-cutoff voltage of 4.8 V used for the materials is too high for the electrolyte to maintain thermodynamically stable. At such high voltage, the highly oxidative transition metal ions (especially the Co4+ ions) of the fully charged LLMO could accelerate decomposition of the electrolyte which generates several gases such as CH3 and CO2 accompanied by aggressive thermal 25



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release. In addition, the fully charged LLMO is unstable at high temperature because oxygen could easily escape from the crystal lattice causing thermal release and decomposition or structural transformation of the material. In comparison tothe pristine LLMO, the LBO surface layer ofLLMO@Li4M5O12@LBO has high chemical stability which could isolate the electrochemically active LLMO@Li4M5O12 part from violent side reactions with the electrolyte. On the other hand, the Li4M5O12@LBO part can block the oxygen loss from the LLMO lattice. As a result, the gas generation and thermal release are significantly depressed resulting in excellent thermal stability.

Figure 14. DSC curves of the LLMO and LLMO@Li4M5O12@LBO samples after the first charge.

4. CONCLUSION In conclusion, we have successfully prepared a new heterostructured Li-excess layered cathode material for lithium ion batteries. The main part of the material is a sub-micron sized Li(Li0.18Ni0.15Co0.15Mn0.52)O2 active material which possesses a high capacity of more than 250 mAh g-1. This Li-excess layered oxide is heterostructured with a Li4M5O12 surface spinel. This Li4M5O12 surface spinel improves the electronic conductivity of the material. More importantly, it could induce the Li4M5O12 phase 26


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transformation of the cathode material which is different from the LiMn2O4-like phase transformation of traditional Li-excess layered oxides. This Li4M5O12 phase transformation is very important for the cathode material to maintain high working voltage and good cycle stability due to its stable Mn4+ ions and high redox potential. The outer surface of the material is coated by a thin Li2O-LiBO2-Li3BO3 glass layer. This glass layer is a good Li ion conductor which significantly improves the ionic conductivity of the material. In addition, it is chemically and electrochemically inert in electrolyte which not only depresses the formation of SEI film but also increases the thermal safety of the fully charged electrode. The above advantages significantly improve the electrochemical performance of Li-excess layered cathode materials which showed not only superior rate capability but also good capacity retention, slight voltage decay and high thermal safety. The heterostructure designed in this work could promote the practical applications of Li-excess layered cathode materials in lithium ion batteries.

ASSOCIATED CONTENT Supporting Information SEM images of the LLMO and LLMO@Li4M5O12@LBO samples (Figure S1), Specific discharge energy of the LLMO and LLMO@Li4M5O12@LBO samples at the 0.2 C rate (Figure S2), AC impedance spectroscopy of the LLMO and LLMO@Li4M5O12@LBO samples (Figure S3), Electronic conductivity and ionic conductivity of the LLMO and LLMO@Li4M5O12@LBO samples (Table S1), GITT analysis of the LLMO and LLMO@Li4M5O12@LBO samples during the first discharge (Figure S4), Cycling performance and rate performance of the LLMO@Li4M5O12@LBO samples with different LBO amounts (Figure S5). The 27



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supporting information of available free of charge on the ACS publication website at DOI: ……….

AUTHOR INFORMATION Corrsponding Author * Yingjin Wei, Email: [email protected] Notes The authors declare no competing financial interest.

ACKONWLEDGEMENTS This work was supported by the Ministry of Science and Technology of China (No. 2015CB251103), National Natural Science Foundation of China (No. 51472104, 21473075, 51272088), the Defence Industrial Technology Development Program (No. B1420133045), and the Graduate Innovation Fund of Jilin University (No. 2015046).

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Table of Contents (TOC) graphic

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High-Performance Li(Li0.18Ni0.15Co0.15Mn0.52)O2@Li4M5O12

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