Stable heteroepitaxial interface of Li-rich layered oxide cathodes with enhanced lithium storage

Author’s Accepted Manuscript Stable Heteroepitaxial Interface of Li-rich Layered Oxide Cathodes with Enhanced Lithium Storage Zhengping Ding, Chunxiao Zhang, Sheng Xu, Jiatu Liu, Chaoping Liang, Libao Chen, Peng Wang, Douglas G. Ivey, Yida Deng, Weifeng Wei www.elsevier.com/locate/ensm

PII: DOI: Reference:

S2405-8297(18)30871-7 https://doi.org/10.1016/j.ensm.2018.12.004 ENSM586

To appear in: Energy Storage Materials Received date: 15 July 2018 Revised date: 4 December 2018 Accepted date: 5 December 2018 Cite this article as: Zhengping Ding, Chunxiao Zhang, Sheng Xu, Jiatu Liu, Chaoping Liang, Libao Chen, Peng Wang, Douglas G. Ivey, Yida Deng and Weifeng Wei, Stable Heteroepitaxial Interface of Li-rich Layered Oxide Cathodes with Enhanced Lithium Storage, Energy Storage Materials, https://doi.org/10.1016/j.ensm.2018.12.004 This is a PDF file of an unedited manuscript that has been accepted for publication. As a service to our customers we are providing this early version of the manuscript. The manuscript will undergo copyediting, typesetting, and review of the resulting galley proof before it is published in its final citable form. Please note that during the production process errors may be discovered which could affect the content, and all legal disclaimers that apply to the journal pertain.

Stable Heteroepitaxial Interface of Li-rich Layered Oxide Cathodes with Enhanced Lithium Storage Zhengping Ding,a,1 Chunxiao Zhang,a,1 Sheng Xu,b Jiatu Liu,a Chaoping Liang,a Libao Chen,a Peng Wang,*,b Douglas G. Ivey,c Yida Deng,d and Weifeng Wei*,a a

State Key Laboratory of Powder Metallurgy, Central South University, Changsha, Hunan

410083, People’s Republic of China b

National Laboratory of Solid State Microstructures, College of Engineering and Applied

Sciences and Collaborative Innovation Center of Advanced Microstructures, Nanjing University, Nanjing 210093, People’s Republic of China c

Department of Chemical and Materials Engineering, University of Alberta, Edmonton,

Alberta, Canada T6G 1H9 d

School of Materials Science and Engineering, Tianjin University, Jinnan District, Tianjin

300350, People’s Republic of China

; [email protected] E-mail: [email protected] * Corresponding authors: Weifeng Wei, Peng Wang

1

These authors contributed equally to this work. 1

Abstract Lithium and oxygen activities can have substantial influences on the kinetics of ion and electron transport and the structural integrity of Li-rich layered oxide (LLO) cathodes, since reversible oxygen redox is ascribed to the extra capacity beyond the theoretical capacity from transition metal redox at high voltages. Herein, we demonstrate a liquid-solid interfacial reaction to generate a heteroepitaxial interface with tunable Li/O activities on LLOs using molten boric acid. The experimental and theoretical analyses indicate that the atomic scale interface is comprised of a disordered rock salt structure containing substantial Li/O vacancies along the layered structure, associated with a segregation tendency of Ni and Co. The formation of this heteroepitaxial interface with Li/O vacancies improves the ionic/electronic conduction and electrochemical/structural stability, leading to a high discharge capacity of 283 mAh g-1 with initial Coulombic efficiency of 91.7% (0.1 C, 2.0-4.7 V vs. Li+/Li), excellent rate performance (246 and 159.7 mAh/g at 1 C and 10 C, respectively) and enhanced cyclic performance with a capacity retention of 92% after 100 cycles. The findings highlight the importance of a well-engineered interface for the design of high performance layered cathode materials for Li storage.

2

Graphical Abstract

Keywords: Li- and Mn-rich layered oxide cathodes; heteroepitaxial interface; liquid-solid interfacial reaction; lithium and oxygen vacancies; surface segregation

1. Introduction Currently, Li-ion battery (LIB) technology is primarily limited by the availability of highenergy density cathodes [1-4]. Traditional cathode materials, such as LiCoO2, LiNiO2, LiMn2O4 and LiFePO4, operate through reversible extraction and reinsertion of Li+ ions during the charge-discharge process, and are considered to balance redox reactions involving transition metal (TM) ions that offer the electrons for charge compensation

[5-10]

. As such, their

theoretical capacities should be constrained by the maximum number of electrons provided by the TM ions in the structure. Nevertheless, recent observations of extra capacities beyond the redox of TM ions in Li-rich layered oxides (LLOs), such as Li1.2Mn0.6Ni0.2O2 [8, 11-18] and Li2Ru0.5Sn0.5O2 [2], have raised an important argument that oxygen anions in the lattice may participate in the redox reaction. 3

To achieve ultrahigh specific capacity in LLOs, an activation process involving Li+ extraction and oxygen oxidation is required during the initial charge process. The roles of surface and bulk oxygen anions on high reversible capacity in LLOs have been described in many studies, even though they have not been verified [2, 11, 13, 19-26]. For instance, Delmas’s group proposed that surface oxygen anions are oxidized into O2 gas and lost from the lattice permanently, leaving some oxygen vacancies in the surface or sub-surface regions [21, 22]. In contrast, Tarascon’s group determined that the ultrahigh capacity is essentially ascribed to the reversible anionic oxygen redox processes in the bulk structure of LLOs [23, 24]. Meng’s group further demonstrated that oxygen vacancies generated during the activation process at high voltages facilitate the migration of TM ions and surface structural transformation and gradual structure destruction from the surface to the bulk

[27-29]

. The formation of a spinel-like

structure at the surface may facilitate the dissolution of TM ions, especially Mn ions, from the cathode surface and the dissolved TM ions migrate through the electrolyte and redeposit on the anode surface [30, 31]. This TM dissolution–migration–deposition (DMD) process is closely related to the voltage decay and capacity fading of the LLOs during long-term cycling[32, 33]. To utilize reversible oxygen redox activity and suppress the irreversible O2 gas evolution upon cycling, the formation of surface oxygen vacancies has been attempted in as-prepared LLOs before electrochemical cycling. For example, the generation of surface oxygenvacancies was done through chemical treatment with nitric acid or hydrazine hydrate, accompanied by heat treatment [34-37]. Although the capacity loss is lowered, the bulk-layered structure tends to transform into spinel or rock salt structures, which reduce the rate and cycling performance [31, 34-39]. Recently, Meng’s group demonstrated improved capacity and cycling stability through heat treatment with CO2 produced from NH4HCO3 decomposition to yield oxygen-vacancy LLOs [40]. In addition, enhanced capacity and slow voltage fading was reported by Aurbach’s group, who exposed Li-rich layered materials to NH3 at 400 oC to form an oxygen-deficient spinel-like phase and cause Li+ removal from the bulk material [41]. Based 4

on theoretical calculations, Ceder’s group verified that oxygen redox activity in Li-rich layered and disordered oxides initiates from specific Li–O–Li configurations that generate orphaned oxygen states, suggesting that both oxygen and Li activities play a significant role in oxygen redox processes [42]. Inspired by these considerations, a liquid-solid interfacial reaction (LSIR) between LLO and molten boric acid (H3BO3) is proposed to generate a heteroepitaxial interface with substantial Li/O vacancies. On the basis of aberration-corrected scanning transmission electron microscopy (STEM), electron energy loss spectroscopy (EELS), energy-dispersive X-ray spectroscopy (EDS) characterization and theoretical calculations, the atomic-scale heteroepitaxial interface is unambiguously determined to be a Ni/Co-concentrated, disordered rock salt-like structure. The distinct crystal and electronic structure leads not only to higher electronic and ionic conductivities, but also to the chemical/electrochemical stabilization of the layered bulk materials, demonstrating a high reversible capacity, superior rate capability and outstanding cyclability. The findings suggest that the present material with a wellengineered interface is a very promising candidate for the cathode material of LIBs. 2. Experimental Section 2.1 Materials synthesis Synthesis of Li-rich Layered oxide (LLO): A co-precipitation method was used to synthesize the carbonate precursor

[43, 44]

. Typically, a MnSO4·2H2O, NiSO4·6H2O and

CoSO4·7H2O aqueous solution (2 mol/L, Mn:Ni:Co molar ratio = 4:1:1), a Na2CO3 solution (2 mol/L) and a NH3·H2O solution (0.24 mol/L) were pumped into a 5 L continuously stirred tank reactor at the same time. The reaction temperature was held at 55 oC and the pH value was kept at 7.8. The carbonate powders were filtered and washed with distilled water several times and dried at 120℃for 24 h. The resulting powders were mixed with Li2CO3 (7 at% excess of Li salt). The mixed powders were heated at 500 oC for 5 h and then at 900 oC for 15 h in air to obtain the pristine Li-rich layered oxide which was labeled as LLO. 5

LSIR process based on molten boric acid: Low-temperature molten salt surface modification using H3BO3 as the molten salt was carried out as follows. LLO and H3BO3 with a mole ratio of 15:1 were dispersed in a 40 mL ethanol solution and stirred for 2 h. After the solution was vaporized at 80℃ , the mixed precursor was calcined at 260 oC for 12 h and then at 500 oC for 4 h under vacuum conditions to obtain the LSIR-treated LLO, referred to as BLLO. The B-LLO material was washed several times with water to remove the product on the surface of LLO during the molten salt reaction and dried at 120 oC for 12 h in a vacuum oven. The as-obtained product was labeled as H-B-LLO. 2.2 Materials characterization Fourier transform infrared spectra (FTIR) were recorded on a Nicolet 6700 spectrometer at room temperature. The chemical compositions of the materials were measured by inductively coupled plasma-atomic emission spectrometry (ICP-AES, PS-6, Baird). Powder X-ray diffraction patterns were obtained using a Bruker AXS D8 Advance X-ray Diffractometer (Germany) with a step size of 0.02o at a dwell time of 2 s. Full-pattern Rietveld refinement was performed using GSAS+EXPGUI software [45]. The morphology and microstructure were characterized with a field emission scanning electron microscope (Nova Nano SEM230, USA). The chemical states of the elements were evaluated by X-ray photoelectron spectroscopy (XPS) (ESCALAB 250Xi, Thermo Scientific, USA). All XPS spectra were calibrated using the C 1s peak at 284.5 eV. Atomic-scale structural and chemical information were obtained using a FEI Titan3 G2 60–300 TEM/STEM equipped with a double aberration corrector. High angle annular dark field (HAADF)-STEM imaging at atomic resolution, electron energy loss spectroscopy (EELS) and energy dispersive X-ray spectroscopy (EDS) mapping were performed at 300 kV. 2.3 Electrochemical characterization Electrochemical performance tests were performed on CR2025 type coin cells, which were assembled using Li metal as the counter electrode, Celgard 2400 as the separator and 1 6

M LiPF6 in EC-DMC (1:1 v/v) as the electrolyte, in an Ar-filled glove box. For the working electrode, active materials, acetylene black and polyvinylidene fluoride with a weight ratio of 8:1:1 were mixed with N-methyl-2-pyrrolidone (NMP) to form a slurry. Subsequently, the slurry was pasted onto Al foil and then dried at 120°C for 12 h in a vacuum oven. The resultant electrode was cut into a circular shape (diameter = 12 mm) as the as-prepared cathode. Galvanostatic charge-discharge measurements were performed with a battery testing system (LANHE CT2001A, P. R. China) in the potential range of 2.0–4.7 V. Electrochemical impedance spectra (EIS) and cyclic voltammograms (CV) were recorded with an electrochemical workstation (PARSTAT 4000A Potentiostat/Galvanostat, Princeton Applied Research). Impedance spectra were recorded in the frequency range of 100 kHz to 0.1 Hz with an applied alternating voltage of 5 mV. 2.4 Theoretical calculations Calculations were performed with the well-known Vienna ab initio simulation package (VASP) [46], which is based on density functional theory (DFT), with a plane-wave basis set and the projector-augmented wave (PAW) method

[47]

. The generalized gradient

approximation (GGA) functional of Perdew-Burke-Ernzerhof

[48]

was used to describe

exchange and correlation interactions. In order to correct the self-interaction error of TM dorbitals, a Hubbard U term (DFT+U) was included in the Hamiltonian [49, 50]. The U values 6.8, 5.9, and 5.2 eV, used in our previous work [51-53], were adopted for Ni, Co and Mn ions, respectively. A 500 eV cutoff energy was used for the plane-wave basis. For k-space integration, the temperature smearing method of Methfessel-Paxton tetrahedron method of Blöchl-Jepsen-Andersen

[55]

[54]

and the modified

were used for dynamical and static

calculations, respectively. For all the compounds studied, a k-point mesh was used to ensure a convergence of 1 meV per unit cell in real space and to obtain the same density per volume unit of the reciprocal lattice. Periodic boundary conditions were applied in the x and y directions and the supercell was allowed to fully relax. The energy criteria were 0.01 and 0.1 7

meV for electronic and ionic relaxations, respectively, whereas 0.01 meV was used for static, self-consistent calculations. Spin-polarized magnetic configurations were considered in all cases. 3. Results and discussion 2.1. Bulk structure of pristine and LSIR-treated LLOs The bulk structures of the materials before and after LSIR treatment were characterized by scanning electron microscopy (SEM), X-ray diffraction (XRD), Fourier-transformed infrared spectroscopy (FTIR) and X-ray photoelectron spectroscopy (XPS). Figure 1a shows a schematic illustration of the interfacial reaction and morphology evolution of the LLO materials before and after LSIR treatment. After the interfacial reaction process, the spherical bulk morphology remains the same as the pristine LLO, but grain growth of the primary nanoparticles is evident in LSIR-treated materials. XRD structural refinements show no significant difference for the pristine and LSIR-treated materials (Figure 1b, Figure S1 and Table S2-S5), implying that the molten H3BO3 treatment does not greatly influence the average crystal structures of the LLO materials. As shown in Table S2, there is 42.68 wt% of the Li2MnO3 (C2/m) phase in the pristine LLO material. After the liquid-solid interfacial reaction (LSIR) treatment, the amount of Li2MnO3 phase decreases to 34.75 wt% for the BLLO sample and 34.57 wt% for the H-B-LLO sample, indicating that the Li2MnO3 constituent was consumed during the LSIR process. Furthermore, it is worth noting that, before water washing, the primary nanograins were covered with water-soluble species such as Li3BO3. The corresponding FTIR, B 1s XPS and ICP-AES results (Figure 1c-1d and Table S1) confirm that water-soluble lithium borate formed on the LLO surfaces after the molten H3BO3 treatment, which indicates that molten H3BO3 treatment has altered the stoichiometry of the surface of the pristine particle. Moreover, as shown in the O 1s, Mn 2p, Co 2p and Ni 2p XPS spectra (Figure S2), slight shifts to lower binding energies for the Mn/Co/Ni cations, associated with higher binding energies for the O anions, suggest that the average valence of 8

the Mn/Co/Ni cations decreases or some oxygen anions are extracted from the lattice, creating oxygen vacancies on the LLO surfaces. These results infer that the LSIR process substantially affects the surface structure and chemistry of the LLO materials.

Figure 1. (a) Schematic illustration and morphology evolution during the interfacial reaction process between Li-rich layered oxides and molten H3BO3. LLO, B-LLO and H-B-LLO refer to the pristine and LSIR-treated materials without and with water washing, respectively. (b) XRD patterns from the three materials. (c) FTIR spectra from the three materials. B-LLO without water washing has two additional peaks (1438 and 1500 cm-1), corresponding to the B-O asymmetric and symmetric stretching modes. (d) XPS spectra of B 1s peak. It is apparent that B-LLO has a strong B peak, which indicates the formation of lithium borate. 2.2. Atomic-scale characterization of the heteroepitaxial interface To understand the local structure and chemistry of the LLO and H-B-LLO materials, highangle annular dark field (HAADF)-STEM images, nanobeam electron diffraction (NBED) patterns and EDS maps are compared in Figure 2. The contrast of the HAADF images is closely related to the atomic number (Z-contrast). The bright spots correspond to TM ions, whereas the Li and O ions are virtually undetectable for this imaging condition. On the basis of the HAADF images along the [010] zone axis projection (Figure 2a-2b and 2e-2f), it is noted that structural changes induced by the LSIR process occur primarily on the surface of 9

H-B-LLO, i.e., within first several atomic layers. When compared with the NBED pattern taken from pristine LLO (Figure 2c), a disordered rock salt-like structure can be identified on the surface of H-B-LLO (Figure 2e), demonstrating that TM ions substitute for some of the Li atom columns in the rock salt-like lattice, giving rise to the extra HAADF imaging contrast. The orientation relationship between the layered bulk and the rock salt-like surface can be defined as: (003)L || (1-11)R and [010]L || [01-1]R. High-resolution EDS results, shown in Figure 2d and 2g and Figure S3, confirm the surface segregation of Ni and Co ions at the expense of Mn ions, suggesting that the Li vacancies generated by the LSIR process are mainly occupied with adjacent Ni/Co ions in the lattice. Note that this conformal Ni/Coenriched interface is evident on every primary particle studied in STEM mode, indicating that the LSIR process allows improved overall coverage of bulk materials in contrast to traditional coating methods. The above observations clearly show that the LSIR treatment not only alters the surface stoichiometry, but also brings about surface phase transitions.

Figure 2. Chemical composition and local structure of the pristine LLO and H-B-LLO materials. (a, b) HADDF-STEM images, (c) corresponding NBED pattern and (d) EDS elemental maps taken from surface region of the pristine LLO. (e, f) HADDF-STEM images 10

and corresponding NBED patterns and (g) EDS elemental maps taken from surface region of the H-B-LLO material. These illustrate the heteroepitaxial interface between the layered bulk and rock salt-like structure, with enrichment of Ni and Co, at the surface. 2.3. STEM-EELS characterization of the heteroepitaxial interface To further assess the electronic structure evolution across the heteroepitaxial interface, atomic-resolution EELS line-scanning from the surface to the bulk of the H-B-LLO material was conducted (Figure 3). Figure 3a shows a HAADF-STEM image with the EELS line-scan range for the H-B-LLO sample. Representative EELS spectra from the Mn-L2,3 and O-K edges taken from three distinct regions from the surface to the bulk are highlighted in Figure 3b. All the spectra in Figure 3b are normalized to the Mn L3 peak. It is evident that the two Mn-L3 and Mn-L2 peaks show a significant shift to lower energies at the interface, when compared with those in the core region (Figure 3b and 3c). Moreover, the L3/L2 intensity ratio declines from the interface to the bulk of the H-B-LLO sample (Figure 3c), which is indicative of a lower Mn oxidation state at the interface. In addition, the significantly suppressed O-K edge peak intensity and the reduced intensity ratio of the pre-peak to the main peak for the O-K edge are evident in the EELS spectrum taken from the interface, which can be ascribed to oxygen vacancies formed at the interface [29, 40, 56]. The chemical distribution of O/Mn/Ni/Co from the surface to the bulk is plotted in Figure 3d. It is worth noting that, when compared with the Mn/Ni/Co ratio of 4:1:1 in the bulk, the Ni/Co-enriched interface possesses an Mn/Ni/Co ratio of about 2:1:1. The corresponding O content is less than 60% on average in the Ni/Co-enriched interface compared with the bulk value of 65%. Also, the lower O and Mn content (higher Ni and Co content) are observed even at ~10 nm away from the surface. In other words, concentration gradients of Li/O vacancies and TM ions are generated by the LSIR process. Thus, as confirmed by the STEM/EELS results, Li/O vacancies associated with Ni/Co segregation lead to a well-defined rock salt-type layer with a distinct

11

electronic and crystal structure, where the oxygen vacancies are coordinated by TM ions and some of the Li vacancies are filled with migrated Ni/Co ions.

Figure 3. (a) HAADF-STEM image of a particle from the H-B-LLO sample with the EELS line-scan shown. (b) EELS spectra of the Mn-L2,3 and O-K edges taken along the green line in (a), where the black dotted lines show the main peaks of the Mn-L2,3 and O-K edges. All the spectra are normalized to the Mn-L3 peak. (c) Mn L3 and Mn L2 energy profiles and L3/L2 intensity ratio as a function of position. (d) Relative atomic percentage of O, Mn, Co and Ni as a function of position based on the EELS spectra in (b). 2.4. Electrochemical performance of pristine and LSIR-treated LLOs

12

Figure 4a compares the initial charge-discharge curves taken from the pristine and LSIRtreated LLOs at 0.1 C (25 mA g-1) between 2.0 and 4.7 V versus Li+/Li. Consistent with other LLO materials, a long charge plateau at ~ 4.5V can be observed in all three materials, which generally is attributed to the electrochemical activation process of the Li2MnO3 component [5761]

. Interestingly, the trend in initial charge capacity is: LLO (325.4 mAh g-1) > H-B-LLO

(308.1 mAh g-1) > B-LLO (294.7 mAh g-1). The subsequent discharge capacity follows a different tendency: H-B-LLO (282.5 mAh g-1) > LLO (261.8 mAh g-1) > B-LLO (251.1 mAh g-1), leading to a dramatic increase in initial Coulombic efficiency from 80.4 to 91.7% (see Figure 4a and Table S6). The corresponding differential capacity profiles (dQ/dV vs. potential), shown in Figure 4b, confirm that the suppressed oxidation peaks at ~4.5 V are detected in both LSIR-treated electrodes, which is consistent with the shorter charge plateaus observed in Figure 4a. This trend indicates that some of the Li2MnO3 constituent has been consumed to form the heteroepitaxial interface during the LSIR process, which is in accordance with the Rietveld refinement results in Table S2. In the subsequent discharge process, H-B-LLO exhibits a higher shift in the discharge plateau (or reduction peaks) and delivers a higher discharge capacity. It is worth noting that B-LLO before water washing shows inferior kinetics when compared with H-B-LLO, which is ascribed to generated lithium borate compounds with poor electrical conductivity. The rate capability and cycling performance further demonstrate the benefits of the generated heteroepitaxial interface (Figure 4c and 4d). At all C-rates, the H-B-LLO electrodes display higher discharge capacities than the pristine and B-LLO electrodes, as shown in Figure 4c and Figure S6d-6f. For instance, the pristine LLO and H-B-LLO electrodes deliver a discharge capacity of 183 and 241 mAh g−1 at 1 C rate, respectively. Even at a high C rate of 10 C (2500 mA g−1), the H-B-LLO electrode still demonstrates a high capacity of 160 mAh g−1, when compared with the capacity (95 mAh g−1) for the pristine LLO electrode. The superior rate performance is ascribed to the improved electrochemical kinetics induced by the 13

well-designed interface. Electrochemical impedance spectroscopy (EIS) and cyclic voltammograms collected at different scan rates were employed to understand the excellent rate capability, and are shown in Figure S4, Figure S5, Table S7 and Table S8. It is apparent that the heteroepitaxial interface leads to enhancement of both the Li ion diffusion coefficient (DLi+) and electron conduction for the H-B-LLO material, indicative of a more favorable charge transfer process across the heteroepitaxial interface. More importantly, the capacity retention plots in Figure 4d confirm the substantial improvement in cycling stability achieved with the H-B-LLO material. However, from the charge-discharge curves for extended cycles (Figure S6), it is noted that all the materials demonstrate voltage decay after 100 cycles and the improved capacity retention rate for the H-B-LLO material comes primarily from the potential region below 3.5 V versus Li+/Li. These results reveal that the Ni/Co-enriched heteroepitaxial interface is beneficial in facilitating the charge-transfer processes (Li-ion and electron transport) and in suppressing Mn dissolution–migration–deposition (DMD) process, leading to the improved rate capability and superior cycling stability of the H-B-LLO material.

14

Figure 4. Electrochemical performance of the LLO, B-LLO and H-B-LLO materials. (a) The initial charge-discharge curves at a rate of 0.1 C (1 C = 250 mA/g). (b) Differential capacity versus potential profiles. (c) Rate capabilities of the LLO, B-LLO and H-B-LLO materials from 0.1C to 10 C. (d) Cycling performance at 1 C rate of the LLO, B-LLO and H-B-LLO materials for 200 cycles. All the electrochemical properties were evaluated from 2.0-4.7 V vs. Li+/Li. 2.5. Phase stability of the heteroepitaxial interface In order to understand the intrinsic mechanism behind the phase transition along the interface, first-principles calculations were used to construct the Li-TM-O phase diagram. The computational details of the calculations are given in the Supporting Information. TM ratios of Mn:Co:Ni = 6:1:1 and Mn:Co:Ni = 2:1:1 were chosen to mimic the phases in the bulk and on the surface of as-sintered particles, respectively. The atomic structures and the ternary phase diagram for the Li-TM-O system with Mn:Co:Ni = 2:1:1 and Mn:Co:Ni = 6:1:1 are 15

shown in Figure 5 and Figure S7, respectively. Several characteristics can be identified from the Li-TM-O phase diagrams. Firstly, the stability of Li2TMO3 is quite different for Mn:Co:Ni = 2:1:1 and 6:1:1. According to the phase diagram, decomposition of Li2TMO3 is as follows: 𝐿𝑖

𝑇𝑀𝑂 (𝐶2/𝑚) ↔ 1⁄2 𝐿𝑖 𝑂 + 1⁄4 𝑂 + 𝐿𝑖

𝑇𝑀𝑂 (𝑙𝑎𝑦𝑒𝑟)

(1)

Li2TMO3 decomposes into Li2O and layered LiTMO2 and oxygen gas. Li2TMO3 is stable up to 1490 K for Mn:Co:Ni = 6:1:1, but only to 830 K for Mn:Co:Ni = 2:1:1. Secondly, the deficient layered LiTMO2 phase with Mn:Co:Ni = 2:1:1 will decompose to LiTM2O4 spinel and layered LiTMO2 from 830-890 K. 𝐿𝑖

𝑇𝑀𝑂 (𝑙𝑎𝑦𝑒𝑟) ↔ (𝑥)𝐿𝑖𝑇𝑀 𝑂 (𝑆𝑝𝑖𝑛𝑒𝑙) + (1 − 2𝑥)𝐿𝑖𝑇𝑀𝑂 (𝑙𝑎𝑦𝑒𝑟)

(2)

On the other hand, the deficient layered LiTMO2 phase with Mn:Co:Ni = 6:1:1 remains stable until the temperature is raised above 1800 K. Thirdly, the spinel phase with Mn:Co:Ni = 2:1:1 transforms into the disordered rock salt LiTM2O3 phase through reactions (3) and (4) at 890K. Note that the two reactions are favorable energetically. 𝐿𝑖𝑇𝑀 𝑂 ( 𝑝𝑖𝑛𝑒𝑙) ↔ (1 − 𝑥)𝐿𝑖𝑇𝑀 𝑂 (𝑟

𝑎𝑙 ) + (𝑥)(𝐿𝑖𝑇𝑀𝑂 (𝑙𝑎𝑦𝑒𝑟) + 𝑇𝑀𝑂(𝑟

𝑎𝑙 )) + 1⁄2 𝑂

(3) 𝐿𝑖𝑇𝑀𝑂 (𝑙𝑎𝑦𝑒𝑟) + 𝑇𝑀𝑂(𝑟

𝑎𝑙 ) ↔ 𝐿𝑖𝑇𝑀 𝑂 (𝑟

𝑎𝑙 )

(4)

Finally, the layered LiTMO2 changes into the disordered rock salt LiTM2O3 phase according to: 𝐿𝑖𝑇𝑀𝑂 (𝑙𝑎𝑦𝑒𝑟) ↔ 1⁄4 𝐿𝑖 𝑂 + 1⁄ 𝑂 + 1⁄2 𝐿𝑖𝑇𝑀 𝑂 (𝑟

𝑎𝑙 )

(5)

The final product, the disordered rock salt LiTM2O3 phase, becomes unstable above 1910 and 1890 K for Mn:Co:Ni = 6:1:1 and Mn:Co:Ni = 2:1:1, respectively.

16

The phase diagram provides a clue to understanding the phase transitions happening on the particle surface. After the LSIR process, in addition to Li and O, a small amount of Mn4+ may dissolve into acid and water to create a surface with a TM ratio near Mn:Co:Ni = 2:1:1. The increase in Co and Ni ions promotes the spinel phase transition. The empty TM and Li positions facilitate the formation of the disordered rock salt structure. The above phase reactions take place at 830 to 890 K, which is just a bit higher than the LSIR temperature 773 K. Since the interface is at the nanoscale, the fast kinetics may reduce the reaction temperature to the experimental value. It should be noted that further increases in Co and Ni ions would not promote the disordered rock salt phase transition. The ternary phase diagram for Mn:Co:Ni = 2:3:3 shows that neither the Li2TMO3 phase nor the disordered rock salt LiTM2O3 phase are stable (Figure S8). At this composition, the spinel LiTM2O4 phase decomposes into rock salt TMO and layered LiTMO2 phases, instead of disordered rock salt LiTM2O3. 2.6. Mechanism of the enhanced kinetics and stability of the H-B-LLO material. On the basis of above results, a mechanism is proposed to understand the multiple functions of the heteroepitaxial interface on the electrochemical kinetics and structural stability of the H-B-LLO materials when cycled in the electrolyte. Firstly, upon charging to the 4.5 V plateau, the generated oxygen vacancies near the heteroepitaxial interface may reduce the oxygen partial pressure and, in turn, suppress the formation of oxygen radicals and gaseous oxygen evolution from the surface during the electrochemical activation process

[62, 63]

. Fewer side

reactions between the highly active oxygen species and the electrolyte species implies that a thinner solid-electrolyte interphase (SEI) layer forms, facilitating Li+ diffusion across the electrode/electrolyte interface (see the electrochemical kinetics results in Figure S4, Figure S5, Table S7 and Table S8). In addition, the disordered rock salt structure has a high Li ionic conductivity as demonstrated by Ceder et al [3, 42]. Secondly, the surface enrichment of Ni/Co is essentially spontaneous during the LSIR process and the simultaneous substitution of empty 17

Li sites with Ni/Co ions leads to the formation of the disordered rock salt-type structure. The Ni/Co-concentrated rock salt-type overlayer has a distinct electronic and crystal structure, which gives rise to better electronic conductivity than the Mn-rich layered bulk material, as supported by the improved rate capability and detailed EIS results (Figure 4, Figure S4 and Table S7). Both of these physicochemical features jointly contribute to the improved electrochemical kinetics of the H-B-LLO material. Thirdly, the Mn DMD process mainly accounts for cycling degradation of Mn-based LLO materials. In contrast, the Ni/Co-enriched rock salt-type interface, which grows in-situ around entire LLO particles, could play an essential

role

in

protecting

against

corrosion

from

the

electrolyte,

providing

chemical/electrochemical stability to the high-activity Mn(III) species, as supported by the improved cyclability in Figure S5. More importantly, compared with common surface-coating layers that generally cause greater resistance and structural mismatch, the LSIR-induced Ni/Co-enriched rock salt-type structure is coherent with the layered bulk structure and electrochemically active in nature, leading to improved kinetics, less mechanical stress and better structural durability of the LLO electrode. These results suggest that this economical and scalable strategy is effective in controlling Li/oxygen activities, protecting the electrode from corrosion by organic electrolytes, suppressing voltage fade and improving long-term cycling stability.

18

Figure 5. (a) Possible crystal structure evolution in the Li-TM-O system with Mn:Co:Ni = 2:1:1. (b) Ternary phase diagrams of Li-TM-O system with Mn:Co:Ni = 2:1:1 at 830, 890, 1800 and 1890K.

4. Conclusions A Li-rich layered oxide (LLO) material, with a rock salt-like heteroepitaxial interface with tunable Li/O activities that possesses multiple benefits, has been fabricated through a liquidsolid interfacial reaction (LSIR) process based on molten boric acid. The nanoscale interface, with substantial Li/O vacancies and segregation of Ni and Co, can suppress the release of gas from the surface and enhance electronic and ionic conduction, leading to less first cycle irreversible capacity loss, a higher discharge capacity and a better rate capability in the LLOs. The Ni/Co concentrated interface may effectively reduce the contact between the Mn species and the electrolyte and suppress the dissolution–migration–deposition (DMD) process for Mn. The chemical/electrochemical stability of the LLOs is remarkably enhanced. As such, the H19

B-LLO material (LSIR-treated LLO after water washing) delivers a high discharge capacity of 283 mAh g-1 with initial Coulombic efficiency of 91.7% (0.1 C, 2.0-4.7 V vs. Li+/Li), excellent rate performance (246 and 159.7 mAh/g at 1 C and 10 C, respectively) and enhanced cyclic performance with a capacity retention of 92% after 100 cycles. The welltailored interface in transition metal-oxide systems may achieve excellent energy and power density in their applications as LIB cathode materials.

Acknowledgements Z.-P. D. and C.-X. Z. contributed equally to this work. The authors would like to acknowledge the financial support from the National Natural Science Foundation of China (51304248 and 11474147), the National Basic Research Program of China (Grant No. 2015CB654901), the Natural Science Foundation of Jiangsu Province (Grant No. BK20151383), the International Science and Technology Cooperation Program of China (2014DFE00200), the Innovation Program of Central South University (2016CXS003), the State Key Laboratory of Powder Metallurgy at Central South University and the Hunan Shenghua Technology Co., Ltd.

Appendix A. Supporting information Supporting Information is available from the Online Library or from the author.

References 20

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