Enhanced Electrochemical Performance of Layered Lithium-Rich

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Enhanced electrochemical performance of layered lithium-rich cathode materials by constructing spinel-structure skin and ferric oxide islands Shi Chen, Yu Zheng, Yun Lu, Yuefeng Su, Liying Bao, Ning Li, Yitong Li, Jing Wang, Renjie Chen, and Feng Wu ACS Appl. Mater. Interfaces, Just Accepted Manuscript • DOI: 10.1021/acsami.6b14862 • Publication Date (Web): 20 Feb 2017 Downloaded from http://pubs.acs.org on February 22, 2017

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ACS Applied Materials & Interfaces

Enhanced electrochemical performance of layered lithium-rich cathode materials by constructing spinel-structure skin and ferric oxide islands Shi Chen, ※, †, ‡, § Yu Zheng, ※, †, § Yun Lu, *, †, § Yuefeng Su, *, †, ‡, § Liying Bao, †, § Ning Li, † Yitong Li, † Jing Wang, †, ‡, § Renjie Chen, †, ‡, § Feng Wu*, †, ‡, § † School of Material Science and Engineering, Beijing Key Laboratory of Environmental Science and Engineering, Beijing Institute of Technology, Beijing, 100081, China. ‡ Collaborative Innovation Center for Electric Vehicles in Beijing, Beijing, 100081, China. §National Development Center of High Technology Green Materials, Beijing, 100081, China. KEYWORDS: lithium-ion batteries, layered lithium-rich cathode material, spinel structure, ferric oxide, surface modification.

ABSTRACT：：Layered lithium-rich cathode materials have been considered as competitive candidates for advanced lithium-ion batteries, due to their merits in high capacity (more than 250 mAh·g-1), low cost and environmental benignity. However, they still surfer from poor rate capability and modest cycling performance. To address these issues, we have proposed and constructed a spinel-structure skin and ferric oxide islands on the surface of layered

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lithium-rich cathode materials through a facile wet chemical method. During the surface modification, Li ions in the surface area of pristine particles could be partially extracted by H+, along with the depositing process of ferric hydrogen. After calcination, the surface structure transformed to spinel structure and ferric hydrogen was oxidized to ferric oxide. The as-designed surface structure was verified by EDX, HRTEM, XPS and CV. The experimental results demonstrated that the rate performance and capacity retentions were significantly enhanced after such surface modification. The modified sample displayed a high discharge capacity of 166 mAh·g-1 at a current density of 1250 mA·g-1, and much more stable capacity retention of 84.0 % after 50 cycles at 0.1C rate in contrast to 60.6 % for pristine material. Our surface modification strategy, combining with the advantages of spinel structure and chemically inert ferric oxide nanoparticles, have been proved to be the key of realizing the layered lithium-rich cathodes with surface construction of fast ion diffusing capability as well as robust electrolyte corroding durability.

1.

INTRODUCTION Lithium-ion batteries are big success in modern society. They power myriad portable

electronic devices and play an increasingly important role in electric vehicles and energy storage systems1. Currently, the high-energy and long-life lithium-ion batteries are urgently demanded. The energy density of the commercialized cathode materials, such as LiCoO2, LiNixCoyMnzO2 and LiFePO4,

continuously increases through various modifications2.

However, their practical capacities are inferior to those of layered lithium-rich cathode materials, hence the later has drawn much attentions as ones of the most promising candidate cathode materials for the next generation Li-ion batteries3.


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In addition to their high specific capacity of 250 mAh·g-1, Li-rich layered oxides (LLOs) 47

, xLi2MnO3·(1-x)LiMO2 (M=transition metal), are low cost and environmental benign, due

to the merit of high abundance and low toxicity of manganese. However, they suffer from some intrinsic defects, such as low initial coulombic efficiency8-10, poor rate capability11-12, continuous decrease of average voltage13-14 and severe capacity fading15. Many strategies have been adopted to address these problems16-28.Among these methods, surface coating is the most commonly used and effective method to suppress the capacity degradation of layered lithium-rich cathode material29. Recently, a lot of work focused on enhancing the rate performance of layered Li-rich cathode materials with spinel structure26, 30-32. For instance, our group have proposed two reasonable designs, spinel/layered heterostructure30 and ultrathin spinel membrane33, to improve the rate performance of layered Li-rich cathode materials. The better rate performances are believed due to the high Li ion conductivity of spinel structure and their structure compatibility with layered lithium-metal oxides. However, spinel structure itself suffers from poor cycling performance associated with the dissolution of Mn. In order to suppress the dissolution of Mn, nano-size oxide particles have been reported to added in as HF scavenger34. HF has been blamed for dissolution of the transition metals and the surface corrosion of cathode materials35, which could be neutralized by the added HF scavenger. In light of above researches, here we propose a construction of spinel-structure skin and nano-size ferric oxide islands on LLOs, so as to exploit the advantages of both spinel structure and nano-size oxide particles. For simplicity, we label the sample with such surface construction as layer@spinel@oxide (LSO). In our approach, spinel structure provides fast lithium ion diffusion channel and nano-size oxide particles help to stable the surface structure. As a consequence, LSO demonstrates enhanced electrochemical performance in terms of rate capability and cycling performance.



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2.

EXPERIMENTAL SECTION

2.1. synthesis of pristine material Nanomaterials have larger specific surface area thereby the side reactions on which are more serious. Here, we chose to test the modification effects on nanomaterials. Pristine layered lithium-rich cathode materials, Li1.2Mn0.6Ni0.2O2, was synthesized by sol-gel method. The detailed processes are as follows. Firstly, stoichiometric amounts of Ni(CH3COO)2·4H2O, Mn(CH3COO)2·4H2O and CH3COOLi·2H2O were dissolved together in distilled water. A 5 % excess of lithium salt was used to offset lithium evaporative loss during the high temperature reaction. Citric acid as chelating agent was added dropwise into the salt solution after dissolution. The PH value of the result solution was adjusted to 7 using ammonia. Then, the solution was evaporated at 75 ℃ with continuous stirring until a gel was obtained. The obtained gel was dried in vacuum oven at 120 ℃ for 12 h. The dried precursors were preliminary calcined at 450 ℃ for 5 in air. Afterward, the obtained powders were pressed into pellets and calcined at 900 ℃ for 12h in air. After cooling, pristine material was obtained. In this work, all the chemicals were of analytical grade and used as-received without further purification. 2.2. Procedures of surface construction The surface modification strategy is composed of partly extraction of Li+ through Li+/H+ ion exchange, deposition of ferric hydroxide and calcination process. The scheme is illustrated in scheme 1. and detailed description is as follows: Scheme 1. Schematic illustration showing the strategy of the surface modification.


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Firstly, 0.1616g Fe(NO3)3·9H20 and 0.046g NH4H2PO4 were dissolved in 50 mL distilled water, respectively. Owing to the ionization effect of H2PO4-, the NH4H2PO4 solution is weakly acidic. Then, 2g as-prepared pristine powders were added into the ferric nitrate solution and ultrasonically dispersed for 30 minutes. Next, NH4H2PO4 solution was slowly dropped into the suspension. The result suspension solution was continuous stirred for 3h. During the process, Li ion in the surface of pristine material was extracted through Li+/H+ ion exchange. In the meantime, Fe(OH)3 gradually deposited on the surface due to its smaller Ksp(Fe(OH)3) than Ksp(FePO4). Thus, the surface component of pristine material transformed to . . . and Fe(OH)3 particles were attached like islands on the surface. The corresponding reactions are as follows:

↔ ↔ . . . ↔ . . . 3 ↔



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Thereafter, the solution was filtered and the obtained powders were dried in vacuum oven at 80 ℃ for 12 h. Finally, they were calcined at 450℃ for 5h, during which the surface atoms were rearranged to form a defect-spinel structure and Fe(OH)3 was oxidized to Fe2O3. After cooling down, the construction was completed.

2.3. Preparation of reference samples Preparing Fe2O3 nanoparticles: 1.616g Fe(NO3)3·9H2O was dissolved in 80 mL deionized water with continuous stirring. Then, 0.46g NH4H2PO4 was added in. The result solution was transparent. Thereafter, PH was adjusted with NH4OH and red brown precipitate appeared immediately. The precipitate was filtered, washed and dried at 80 ℃ before sintering at 450

℃ for 5 h. Here, the as-prepared ferric oxide was micro-size particles through the observation of SEM. After ball milling for 6 h, the micro-size particles became nanoparticles. Finally, 0.028g as-prepared Fe2O3 nanoparticles was mixed with 1.6g LLOs for 1h using agate mortar. Constructing surface spinel layer: 0.046g NH4H2PO4 was dissolved in 50 mL distilled water. Then, 2g as-prepared pristine powders were added in. The result suspension solution was continuous stirred for 3h. Thereafter, the particles were filtered, washed and dried at 80

℃ before sintering at 450 ℃ for 5 h. 2.4. Characterization methods Morphology and structure characterization: Power X-ray diffraction (XRD) was performed on Rigaku UlTIMA IV – 185 with Cu Kα radiation from 2θ=10°~ 90° at a scan rate of 8° per minute. The morphological studies were observed by FEI QUANTA 6000 scanning electron microscope (SEM) and JEOL JEM-1200EX transmission electron microscope. The X-ray photoelectron spectroscopy (XPS) was conducted on PHI Quantera II and analyzed by the software of XPSpeak. The energy of the spectra was calibrated by the


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binding energy of the hydrocarbon (C-H) at 284.6 eV. Raman spectroscopies were carried out on the Renishaw RM2000 with RL 633 nm laser.

Electrochemical performance tests: Electrochemical performances of all samples were tested with coin-type half cell (2025) at 30 ℃. The positive electrode was made of 80 wt % as-synthesized cathode materials, 10 wt % acetylene black and 10 wt % polyvinylidene difluoride. 1 M LiPF6 dissolved in mixed solvents of ethylene carbonate (EC) and dimethyl carbonate (DMC) with volume ratio 1:1 was used as electrolyte. Circular lithium metal was used as negative electrode and Celgard 2400 membrane was used as the separator. Galvaostatic charge/discharge test, rate performance and Galvanostatic Intermittent Titration Technique (GITT) were performed by Land battery test system (Wuhan, China) in the voltage range of 2.0~4.8V. The current density of 250 mA·g−1 was defined as 1 C rate during test. The cells were charged and discharged for one cycle at 0.1 C rate before they were used to measure the performance cycled at 1 C rate. During rate performance test, all charge current densities were 0.1 C rate. In GITT test, the cells were charged at 0.1C rate for an interval of 1 h followed by an open circuit stand for 2 h to allow the cell voltage to relax to the ready-state value. The procedure was repeated until the cell potential reached the set cutoff value. Cyclic voltammetry (CV) measurement and Electrochemical impedance spectra (EIS) were conducted on CHI660 electrochemical workstation. The potential window was 2-4.8V and the scan rates was 0.1 mV·s-1 in CV test. EIS measurement was performed at frequencies from 100 kHz to 0.01 Hz with an AC perturbation signal of 5 mV.

3.

RESULTS AND DISCUSSION

3.1. Characterization of the morphology and structure SEM images and energy dispersive X-ray (EDX) results of LLOs and LSO are shown in Figure 1. Clearly, the as-prepared particles are irregular convex polyhedrons and primary



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particle sizes are between 100~200 nm. After surface modification, many additional nanosize particles emerge on the surface of original particles (some of which are marked with circle in Figure 1c). Although some of the additional nanoparticles agglomerate together, the particle size is much smaller than that of the primary particles. In order to identify the chemical composition of the nanoparticles, the ICP measurement has been carried out to determine the contents of Li, Mn, Ni, Fe and P (normalized to the Ni content), and the results are shown in Table 1. The chemical formula of LLOs is close to the designed Li1.2Ni0.2Mn0.6O2 chemical stoichiometry. After surface modification, the Li content of LSO decreases significantly, as expected with the consideration of the Li+/H+ ion exchange. In the process of modification, a tiny amount of FePO4 is inevitably produced due to the presence of phosphate, which is detected by ICP for LSO. However, the molar ratio of detected P to Fe is only 16.7%, which confirms that the main coating agent is iron oxide rather than phosphate. The small amount of P compared to other elements also explains the experiment results that the EDX spectrum shows no trace of element P and the signal of P 2p in the XPS spectra (Figure S1) for LSO is also negligible. Thus, combining the results of EDX, ICP and XPS, it is reasonable to ascribe the coating nanoparticles to iron oxide.


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Figure 1. SEM images and energy dispersive X-ray (EDX) spectroscopy: (a) SEM image and (b) EDX spectroscopy of LLOs;(c) SEM image, (d) EDX spectroscopy and (e-g) EDX mapping of LSO. Table 1. Chemical Compositions of LLOs and LSO. Sample

Li

Mn

Ni

Fe

P

LLOs

5.906

3.010

1

-

-

LSO

5.340

3.024

1

0.155

0.026

The crystal structures of both samples have been characterized by XRD as shown in Figure 2. The two collected patterns are nearly the same. The main patterns can be well indexed with —

α-NaFeO2 layered structure (R 3 m symmetry)21,

36-38

. The weak super-lattice reflections

emerged around 2θ=20~23° are corresponded to the Li2MnO3 component (C/2m symmetry), indicating the LiMn6 cation arrangement in the transition metal layers4, 39-40. The clear splits of (006), (102) and (108), (110) peaks signify well layered structure17, 41-42. The calculated lattice parameters are listed in Table S1. Both lattice parameters a and c of LSO are larger



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than that of LLOs, which indicates unit cell expansion in a and c directions. The increased aparameter is attributed to the decrease in average metal-metal distance due to the reduction of some Mn4+ during the process of heat treatment of surface modification43. The extraction of Li ions during the process of Li+/H+ ion exchange increases the electrostatic repulsion between the oxygen layers, which leads to the lattice expansion along c-axis. Both the ratios of I(003)/I(104) (integrated intensity ratio) are greater than 1.2 which suggest low cation disorder44. Meanwhile, the intensity of peaks displays slightly decreased caused by the process of Li+/H+ ion exchange after surface modification. From the enlarged view in Figure 2b and 2c, no signal from spinel phase is found, which may due to its tiny amount. In order to identify the phase of iron oxide, we prepared Fe2O3 nanoparticles under similar conditions. The XRD pattern of as-prepared Fe2O3 nanoparticles (seen in Figure S2) shows no sharp peaks, indicating the decomposition product of ferric hydroxide at 450 ℃ is amorphous.

Figure 2. Powder XRD patterns of the two samples and some Miller indices of the main peaks; (b,c) are enlarged view in (a). To further identify the microstructure, we have observed LLOs and LSO under highresolution transmission electron microscope (HRTEM) and the typical images are shown in Figure 3. As for LLOs, the measured interlayer spacing of ca. 0.472 nm in Figure 3b corresponds to (003) fringes of typical layered structure16, 19. The corresponding selected area


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electron diffraction (SAED) patterns in Figure 3c present the reflections of (003) plane and (110) plane expected from rhombohedral α-NaFeO2 structure19, 45-46. After surface modification, the bulk layered structure is preserved and the measured bulk interplanar spacing of ca. 0.42 nm in Figure 3e is indexed with the (020) plane of C/2m symmetry. Different from the PL particle, the fringes of LSO are not stretched to the surface and display distinct characteristics. In the surface area, lattice fringes with interplanar spacing of ca. 0.248 nm and 0.294 nm are observed, which could be indexed with (311) plane and (220) plane of the cubic spinel (Fd-3m) structure. While on the surface, there is a tiny additional nano-particle adhered to the surface of this particle in Figure 3d, as marked in green rectangle. The measured lattice fringes with interplanar spacing of ca. 0.223 nm and ca. 0.271 nm, are quite indexed with (113) plane and (104) plane of R-3c symmetry, which are expected from ferric oxide. Thereby, the microstructures observed by HRTEM clearly confirm the as-design of spinel-structure skin and nanoparticle island, which reveals our strategy is carried out successfully.

Figure 3. HRTEM images of both LLOs and LSO: (a) The TEM image of the PL particle; (b) the HRTEM image in (a); (c) the SAED of (a); (d) The TEM image of the LSO particle; (e, f) the HRTEM image in (d). XPS as a powerful tool has been also applied to characterize surface chemical properties of both samples, as shown in Figure 4. It is demonstrated that there are two peaks at ca. 723.7



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eV and ca. 710.8 eV and a satellite peak at ca. 719.2 eV in the Fe 2p spectra, which are assigned to Fe2O347. The Mn 2p3/2 spectra obtained from LLOs and LSO are compared in Figure 4b. The Mn 2p3/2 peak of LLOs is located at 641.8 eV. After surface modification, the Mn 2p3/2 peak of LSO become broader and a peak with lower binding energy can be clearly observed. This indicates the local environment change originating from the formation of spinel structure48. Thus, the XPS results further confirm the existences of ferric oxide and the formation of spinel structure.

Figure 4. XPS and Raman spectra of LLOs and LSO: (a) Fe 2p region for LSO; (b) Mn 2p region for LLOs and LSO; (c) Raman spectroscopies of both samples. To further verify the transformed spinel structure, Raman spectroscopy as another powerful tool of determining the short-range local structure has been applied. Both two samples share the typical characterizations of lithium-rich layered materials, while an additional shoulder band emerging around 670 nm−1 for LSO, as indicated by the arrow, suggests the existence of spinel Li1+xMn2O47, 33. 3.2. Characterization of electrochemical performance To explore the advantages of our approach, we have investigated the electrochemical performance as illustrated in Figure 5 and the specific values are summarized in Table S2. It is demonstrated both the rate performance and the capacity retentions of LSO are significantly enhanced compared to those of LLOs. As shown in Figure 5a, the two samples deliver close initial specific discharge capacity at 0.1C rate. However, when testing the high-


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rate performance, LSO shows much superior performance thanks to the surface spinel structure. The maximum discharge capacity curves at different rates are compared in Figure 5b. The discharge capacity falls below 200 mAh·g-1 for LLOs when discharged at 1 C rate, while that of LSO is ca. 220 mAh·g-1 and its discharge capacity is still beyond 200 mAh·g-1 even at 2C rate. When discharged at 5 C rate, the maximum discharge capacity can achieve as high as ca. 166 mAh·g-1 for LSO. When cycled at 0.1C rate in Figure 5c, LLOs suffers sever capacity fading due to the nanometre size effect and unprotected surface. Its discharge capacity fades rapidly from 275.9 mAh·g-1 to 167.2 mAh·g-1 after 50 cycles, with a capacity retention of 60.6 %. Owing to the protection of ferric oxide nanoparticles, the capacity retention of LSO is improved to 84.0 % at the same condition. In addition, when cycled at 1C rate in Figure 5d, LSO not only delivers a superior discharge capacity of 230.6 mAh·g-1, but also displays better capacity retention of 71.7% compared to the capacity retention of 60.6% for pristine material after 150 cycles. The Coulombic efficiency at 1C rate is slight higher than that at 0.1 C rate. When testing at 0.1C, the smaller charging and discharging current will lead the electrolyte to be exposed to high voltage region for a longer time in each cycle. It has been shown that alkyl carbonate based electrolytes undergo oxidation at potentials in excess of 4.2 to 4.4 V vs Li+ leading to H+ generation49. The longer residence time under high voltage, the more side reactions. Thus, LLOs shows better cycle performance at 1C-rate than 0.1C-rate. After surface coating, the Coulombic efficiency at 0.1C rate is slight improved. Consequently, we could conclude that our approach shows great effect on enhancing the rate capability and cycling performance of the layered lithium-rich cathode materials.



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Figure 5. Electrochemical performances of the two samples: (a) initial charging and discharging curves; (b) maximum discharge capacity curves at different rates; discharge capacity vs. cycle number when cycled at 0.1 C (c) and 1 C (d) rate. To further confirm the effect of spinel-structure skin and ferric oxide nanoparticles, we have prepared the sample of L@O (LLOs + Fe2O3 nanoparticles) and L@S (LLOs with surface spinel structure) for comparison. We tested their cycling stability at 1C rate and rate performance, as compared in Figure 6. All the other samples show improved electrochemical performance compared with LLOs. Specially, the samples with Fe2O3 nanoparticles exhibit superior cycling performance, while the samples with surface spinel structure deliver better rate performance. These results strongly demonstrate the roles of Fe2O3 nanoparticles as HF scavenger and spinel structure as fast lithium ion diffusion channel.


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Figure 6. The compared electrochemical performances: (a) the cycling stability at 1 C rate and (b) rate performance of the four samples; the rate performance was tested with a charging current of 0.1C.



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3.3. Discussion In order to achieve more detailed understanding of electrochemical behaviours, we have carried out further characterization. Firstly, the CV profiles of the two samples for the first cycle are displayed in Figure 7. The observed two oxidation peaks below 4.5V result from the oxidation of Ni2+/Ni3+ and Ni3+/Ni4+. The strong peak at ca. 4.5 V, corresponding to the long plateau around 4.5V in initial charging/discharging curves in Figure 5a, is attributed to the reaction involving the removal of lithium and oxygen from Li2MnO350-53. Carefully examining the profile, a weak oxidation peak around 3.2 V is found for LSO and it could be observed more clearly in the corresponding dQ/dV plots of the first charging curve (seen in Figure S3), which results from the existence of the spinel structure. The whole oxidation area for LSO is smaller than that for pristine material, indicating less charge capacity. Given the facts of the less Li content and the reduced charge capacities in both regions (below and at ~4.5V plateau, as seen in Figure 5(a)) for LSO, we attribute the decline of the 1st charge capacity to two reasons: (1) Some Li ion lost during the process of ion exchange, which indicates the reduction of available Li source; (2) Ferric oxide as electrochemically inert cathode material adds extra mass, thus lowering the specific charge capacity. During the discharging, apart from the two reduction peaks above 3.5 V corresponding to the reduction of Ni4+ for both two samples and the special reduction peak at ca. 2.9 V relating to spinel structure in the surface of LSO, there are peaks below 3.5 V that signifies the partly reduction of Mn4+ in the bulk5, 39, 54.


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Figure 7. CV profiles of the two samples at a scan of 0.1 mV·s-1 between the voltage range of 2~4.8V. Differential capacity (dQ/dV) plot as useful tool to observe the capacity change at various voltages has been applied to investigate reason for the different capacity retentions. Figure 8 shows the dQ/dV plots of the two samples for the 2nd, 18th, 34th and 50th cycle when cycling at a current density of 25 mA·g-1. It is demonstrated that the cathodic peak intensities above 3.5V decrease upon cycling while the cathodic peak around 3.3 V shifts low to ca. 2.9 V after 50 cycles. This phenomenon (marked with background color) signifies the phase transformation from layered to spinel14, 55-56. Compared to the continuously decreased intensity of the cathodic peaks below 3.3V for LLOs, LSO retains considerable discharge capacity originated from the reduction of Mn4+. Obviously, although LSO could not suppress the phase transformation from layered to spinel, the deterioration of transformed spinel structure is suppressed. We believe that is the primary reason for the better capacity retention for LSO.


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Figure 8. The dQ/dV curves of LLOs (a) and LSO (b) for the 2nd, 18th, 34th and 50th cycle with a 25 mA·g-1 current density between 2 V and 4.8 V In order to explore the reason for the larger amount electrochemical reversible manganese of LSO, GITT test has been carried out to assess the over-potentials at various stages. As shown in Figure 9, during the initial charging, the over-potentials are limited until period 5 where the cells are charged to 4.5 V. This phenomenon implies that the process associated with lithium ion extraction from the LiMn0.5Ni0.5O2 component is highly reversible, while the increasing overpotentials starting from period 6 to 15, suggest that the activation process of the Li2MnO3 component is kinetically limited12, 46, 57.


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Figure 9. GITT curves of the two samples for the first cycle at a charge/discharge current density of 25 mA·g-1: (a) GITT curves; inset: the corresponding OCVs; (b) The calculated specific overpotential at each stage. During the first discharge process, the equivalent open circle voltages (OCVs) of the two samples are nearly the same, but the over-potential of pristine material at the end of discharging increases more rapidly. During period j, the voltage of pristine material reaches to the lower cutoff voltage quickly while there is another period k for LSO. Definitely, it is the larger overpotential below 3.5 V (marked with background pattern) during discharging that limits the amount of reduced Mn4+ for LLOs. Namely, the surface modification endows LSO a higher discharge capacity by reducing its over-potential and leveraging more Mn4+/Mn3+ couples. This reduced over-potentials could be also attributed to a better condition of the surface structure which is protected by ferric oxide from the eroding of HF. To further prove the protection effect, EIS technique has been introduced to assess the change of the interface resistance. Figure 10 shows the Nyquist plots of the two samples in the second and fiftieth charges when cycled at a current density of 500 mA·g-1. All of the Nyquist plots show two semicircles and one slop. Based on previous researches on the same type of layered cathode materials58-60, the semicircle in the high frequency is related to the lithium ion diffusion through the solid-electrolyte interface (SEI), while the second semicircle in the middle frequency is associated with the charge-transfer reaction between the surface film and the active cathode mass, and the diameter of semicircle represents the charge-transfer resistance (Rct). Accordingly, the fitted plots are shown in Figure S4 as well as discharge capacity vs. cycle number. The corresponding impedance parameters are listed in Table S3.


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Figure 10. The Nyquist plots of the two samples in the second and fiftieth charges. The data were acquired over 100 KHz ~ 0.01 Hz frequency range after the cells charging to 4.0 V at a current density of 25 mA·g-1. In the second cycles, the measured resistances for the two samples are similar and the Rct of LLOs is even smaller than that of LSO. However, the charge-transfer resistances of both samples increase sharply after 50 cycles, which implies the interface structure suffering from damage. Obviously, the Rct of LSO is much lower than that of the pristine material after 50 cycles. The protection of the ferric oxide nanoparticles from the eroding of HF in the electrolyte is believed to be responsible for the lower surface charge-transfer resistance of LSO. 4.

CONCLUSION In summary, we have proposed and constructed a spinel-structure skin and ferric oxide islands

on the surface of layered lithium-rich cathode materials through a facile wet chemical method. This symbiotic surface structure was verified by EDX, HRTEM, XPS and CV. Through this strategy, spinel structure provides fast lithium ion diffusion channel and ferric oxide nanoparticles help to stable the surface structure by neutralizing HF in the electrolyte. It is


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demonstrated that the rate performance and capacity retentions are significantly enhanced after such surface modification. The GITT and EIS results confirms the lower surface charge-transfer resistance in the cycling which should result from the better surface condition for LSO. Our modification strategy is able to modify the subsurface structure and surface environment of layered lithium-metal oxides particles, simultaneously. We anticipate this design should inspire the development of advance intercalation materials and energy materials.

ASSOCIATED CONTENT Supporting Information. The following files are available free of charge. XPS spectra of P 2p region for LSO, XRD pattern of Fe2O3 nanoparticle, the lattice parameters of the two samples, the detailed values of electrochemical performances for the two samples, the dQ/dV plots of the two samples corresponding to their initial charge/discharge curves, the fitted Nyquist curves and the impedance parameters of two samples in the second and fiftieth charges at the voltage of 4.0 V. ( PDF).

AUTHOR INFORMATION Corresponding Author * (Y.L) E-mail: [email protected]. Tel: 86-010-68918099. * (Y.S) E-mail: [email protected]. * (F.W) E-mail: [email protected].


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Author Contributions

※S.C. and Y.Z. contributed equally to this work. Notes The authors declare no competing financial interest. ACKNOWLEDGMENT This work was funded by the National Natural Science Foundation of China (51472032, 21573017), National Key Research and Development Program (2016YFB0100301), Program for New Century Excellent Talents in University (NCET-13–0044) and the Special Fund of Beijing Co-construction Project. REFERENCES (1) Kim, T.-H.; Park, J.-S.; Chang, S. K.; Choi, S.; Ryu, J. H.; Song, H.-K. The Current Move of Lithium ion Batteries Towards the Next Phase. Adv. Energy Mater. 2012, 2, 860-872. (2) Liu, W.; Oh, P.; Liu, X.; Lee, M. J.; Cho, W.; Chae, S.; Kim, Y.; Cho, J. Nickel-Rich Layered Lithium Transition-Metal Oxide for High-Energy Lithium-ion Batteries. Angew. Chem., Int. Ed. Engl. 2015, 54, 4440-4457. (3) Hy, S.; Liu, H.; Zhang, M.; Qian, D.; Hwang, B.-J.; Meng, Y. S. Performance and Design Considerations for Lithium Excess Layered Oxide Positive Electrode Materials for Lithium ion Batteries. Energy Environ. Sci. 2016, 9, 1931-1954. (4) Thackeray, M. M.; Kang, S.-H.; Johnson, C. S.; Vaughey, J. T.; Benedek, R.; Hackney, S. A. Li2MnO3-Stabilized LiMO2 (M = Mn, Ni, Co) Electrodes for Lithium-ion Batteries. J. Mater. Chem. 2007, 17, 3112-3125. (5) Johnson, C. S.; Kim, J. S.; Lefief, C.; Li, N.; Vaughey, J. T.; Thackeray, M. M. The Significance of the Li2MnO3 Component in ‘Composite’ xLi2MnO3·(1−x)LiMn0.5Ni0.5O2 Electrodes. Electrochem. Commun. 2004, 6, 1085-1091. (6) Johnson, C. S.; Li, N. C.; Lefief, C.; Vaughey, J. T.; Thackeray, M. M. Synthesis, Characterization and Electrochemistry of Lithium Battery Electrodes: xLi2MnO3·(1−x)LiMn 0.333Ni0.333Co0.333O2 (0 ≤ x ≤ 0.7). Chem. Mater. 2008, 20, 6095-6106. (7) Yu, H.; Zhou, H. High-Energy Cathode Materials (Li2MnO3-LiMO2) for Lithium-ion Batteries. J. Phys. Chem. Lett. 2013, 4, 1268-1280. (8) Lu, Z.; Dahn, J. R. Understanding the Anomalous Capacity of Li/Li[NixLi(1/3-2x/3)Mn(2/3x/3)]O2 Cells Using in situ X-Ray Diffraction and Electrochemical Studies. J. Electrochem. Soc. 2002, 149, A815-A822.


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Enhanced Electrochemical Performance of Layered Lithium-Rich

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