Mitigating Voltage Decay of Li-Rich Cathode Material via Increasing Ni

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Mitigating Voltage Decay of Li-Rich Cathode Material via Increasing Ni-Content for Lithium-Ion Batteries Ji-Lei Shi, Jie-Nan Zhang, Min He, Xu-Dong Zhang, Ya-Xia Yin, Hong Li, Yu-Guo Guo, Lin Gu, and Li-Jun Wan ACS Appl. Mater. Interfaces, Just Accepted Manuscript • DOI: 10.1021/acsami.6b06733 • Publication Date (Web): 20 Jul 2016 Downloaded from http://pubs.acs.org on July 21, 2016

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ACS Applied Materials & Interfaces

Mitigating Voltage Decay of Li-Rich Cathode Material via Increasing Ni-Content for Lithium-Ion Batteries Ji-Lei Shi,†,‡ Jie-Nan Zhang,# Min He,# Xu-Dong Zhang,†,‡ Ya-Xia Yin,† Hong Li,# Yu-Guo Guo,*,†,‡ Lin Gu,*# and Li-Jun Wan*,†,‡

†

Beijing National Laboratory for Molecular Sciences, CAS Key Laboratory of Molecular

Nanostructure and Nanotechnology, Institute of Chemistry, Chinese Academy of Sciences (CAS), Beijing 100190, (P. R. China). #

Beijing National Laboratory for Condensed Matter Physics, Institute of Physics, Chinese

Academy of Sciences (CAS), Beijing 100190, (P. R. China). ‡

University of Chinese Academy of Sciences, Beijing 100049, (P. R. China).

ACS Paragon Plus Environment

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ABSTRACT Li-rich layered materials have been considered as the most promising cathode materials for future high-energy-density lithium-ion batteries. However, they suffer from severe voltage decay upon cycling, which hinders their further commercialization. Here, we report a Li-rich layered material 0.5Li2MnO3·0.5LiNi0.8Co0.1Mn0.1O2 with high nickel content, which exhibits much slower voltage decay during long-term cycling compared to conventional Li-rich materials. The voltage decay after 200 cycles is 201 mV. Combining in situ X-ray diffraction (XRD), ex situ XRD, ex situ X-ray photoelectron spectroscopy (XPS) and scanning transmission electron microscopy (STEM), we demonstrate that nickel ions act as stabilizing ions to inhibit the JahnTeller effect of active Mn3+ ions, improving d-p hybridization and supporting the layered structure as a pillar. In addition, nickel ions can migrate between the transition-metal layer and the interlayer, thus avoiding the formation of spinel-like structures and consequently mitigating the voltage decay. Our results provide a simple and effective avenue for developing Li-rich layered materials with mitigated voltage decay and a long lifespan, thereby promoting their further application in lithium-ion batteries with high-energy-density. KEYWORDS Lithium batteries, cathode materials, high-capacity, Li-rich, voltage decay


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1. Introduction The state-of-the-art lithium-ion battery technology has conquered the portable electronics market; however the energy density still needs to be improved for its application in electric vehicles (EVs) and hybrid electric vehicles (HEVs).1-2 One way to satisfy the demand of the current and future energy storage devices is to improve the gravimetric and volumetric energy densities of commercialized lithium-ion battery materials;3-5 another option is to develop new high-capacity electrode materials.6-11 Li-rich materials, first introduced by J. Dahn and M. Thacheray, with a general formula of zLi2MnO3·1-zLiNi1-x-yCoxMnyO2 are considered attractive cathode materials; they exhibit a reversible capacity exceeding 250 mA h g-1, which is substantially greater than that of current commercialized cathode materials.12-19 Unfortunately, these Li-rich high-capacity materials suffer from poor electrode kinetics and serious voltage decay during cycling, which deteriorates their electrochemical performance, particularly the energy output of the battery.20-28 All these shortcomings inhibit the successful commercialization of these materials in high-energy-density lithium-ion batteries.29-34 To explore effective strategies for alleviating the continuous voltage decay of Li-rich materials, considerable efforts have been devoted to interpreting its origin.35-37 The voltage decay is generally acknowledged to originate from the continual structural transition from a layeredlayered phase to a spinel-like phase during the repeated lithiation and delithiation processes.38-39 This process is related to the trapping of transition-metal (TM) ions in the intermediary tetrahedral sites as they migrate from octahedral sites in the TM slab to octahedral sites in the Li slab through intermediary tetrahedral sites during the lithiation and delithiation processes.40-41 The migration of TM ions is an intrinsic and inevitable characteristic of lithium-rich layered materials, where repulsion forces between the two neighbouring oxygen layers are minimized to


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avoid collapse of the layered structure in its deep delithiation state. Hence, mitigating the voltage decay of Li-rich materials and developing high-performance cathode materials for high-energydensity lithium-ion batteries is difficult. J. R. Dahn et al. report a very impressing core-shell material and not only found this Mn-rich shell can protect the Ni-rich core from side reactions, but also found this material expresses a little voltage decay42. To promote the practical application of high-capacity Li-rich layered cathode materials， herein, we

report

a

cost-effective

Li-rich

layered

cathode

material

of

0.5Li2MnO3·0.5LiNi0.8Co0.1Mn0.1O2 (LL-811) with high-energy-density and high cycling stability. The novelty of this Li-rich material arises from its substantially slower voltage decay during prolonged charge-discharge cycling compared to previously report Li-rich materials. Even after 200 cycles at a rate of 0.2 C (1 C = 250 mA h g-1), its voltage decay is as low as 201 mV. The unique advantage of this LL-811 Li-rich cathode material may arise from the high nickel content in the layered (R3തm) phase. For nickel-rich layered lithium TM oxides some Ni2+ ions commonly occupy the Li+ sites in the lithium layer; this mixed occupancy is known as cation mixing. The cation mixing, to some extent, improves the structural stability by supporting the Li slabs and reducing repulsion of neighbouring oxygen layers during the delithiation process. Likewise, cation mixing could play a positive role in the Li-rich materials.2,

43-44

Moreover,

during lithiation process prior redox of Ni4+/Ni2+ keeps the average oxidation state of Mn above +3 and therefore reduces the Jahn-Teller effect of active Mn3+ ions and improves the structural stability.45-47 In addition, traditional TM manganese ions in Li-rich materials are prone to be trapped in the intermediary tetrahedral sites during repeated charge and discharge cycling. Large Ni2+ ions (0.69 Å), however, are difficult to be trapped in the intermediary tetrahedral sites, thereby alleviating the voltage decay in Li-rich layered materials.37, 48-49


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2. Experimental Section 2.1 Preparation of pristine materials The pristine Li-rich layered materials 0.5Li2MnO3·0.5LiNi0.8Co0.1Mn0.1O2 (LL-811) and 0.5Li2MnO3·0.5LiNi1/3Co1/3Mn1/3O2 (LL-111) were produced via a co-precipitation method. The appropriate amounts of NiSO4·6H2O, CoSO4·7H2O, and MnSO4·H2O were dissolved in distilled water to form a 2.0 mol/L TM solution. The TM solution, a 2.0 mol/L Na2CO3 solution and the desired amount of NH3·H2O were separately pumped into a continuously stirred tank reactor (capacity 5 L). The obtained co-precipitated particles were filtered, washed and dried at 80 °C for three days. Finally, the precipitates were mixed with Li2CO3 and calcined at 450 °C for 5 h and then 850 °C for 12 h in air. The product was then cooled to room temperature in the furnace to yield the pristine material. 2.2 Electrochemical tests Cathode electrodes were prepared by mixing active materials with carbon black and polyvinylidene fluoride (PVDF) in N-methylpyrrolidone (NMP) as the solvent in a weight ratio of 8:1:1 and then coating the slurry onto Al foil. After the coated foil was vacuum dried at 80 °C for 12 h, electrodes were punched from the foil and weighed. The range of active-material loading was 4-5 mg cm-2. Type 2032 coin cells were assembled in an argon-filled glove box using Li metal as the negative electrode, a Celgard separator and an electrolyte consisting of LiPF6 (1 M) dissolved in ethylene carbonate, dimethyl carbonate and diethyl carbonate mixed in a 1:1:1 volume ratio. Galvanostatic discharge/charge cycles were performed on a LAND system in the voltage range from 2 to 4.7 V (vs. Li+/Li). Cyclic voltammetry (CV) was carried out using a PARSTAT (VMC) at a scan rate of 0.05 mV s-1 in the potential range from 2 to 4.7 V (vs.


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Li+/Li). After electrochemical testing, the cells were disassembled in a glove box and the electrodes were washed with dimethyl carbonate five times and dried in the glove box for one week. The electrodes were also protected from air exposure during the ex situ XRD and XPS tests. 2.3 Materials characterizations The morphology of the pristine materials was determined by scanning electron microscopy (SEM, JEOL 6701). Cross sections for SEM analysis were prepared using a focused-ion beam microscope (Helios NanoLab 600i). The components of the pristine materials were obtained using an electron-probe X-ray micro-analyser (EPMA-1720). XRD patterns of pristine materials were collected using a Bruker D8 Advance diffractometer equipped with a Cu-Kα radiation source (λ1 = 1.54060 Å, λ2 = 1.54439 Å) and with a LynxEye XE detector. The patterns of the pristine materials were refined on the basis of the Rietveld method using the TOPAS software. The magnetic measurements were tested using a SQUID magnetometer (MPMS). The temperature dependences of the magnetic susceptibility were measured in a magnetic field of 0.1T by a zero-field-cooling process. The magnetization plots were tested at 5K. The differential scanning calorimetry (METTLER) tested using a stainless-steel sealed pan with a gold-plated copper seal. The electrodes were charged to 4.7 V and disassembled in an Ar-filled box. The scan rate used is 10 °C min-1. In situ XRD analysis was carried out in a homemade Swageloktype stainless-steel cell. ABF and HAADF images were obtained with a cold field-emission gun and double-hexapole Cs correctors (CEOS GmbH, Heidelberg, Germany) and a scanning transmission electron microscope (JEOL, Tokyo, Japan) operated at 200 kV; the spatial resolution of the microscope was defined by the probe-forming objective lens to be better than 80 picometres. XPS measurements were performed on a Thermo Scientific ESCALab 250Xi


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using 200 W monochromated Al- Kα radiation. The electrodes carefully transferred to the XPS spectrometer under a homemade protector without any exposure to air. The XPS analysis was performed using a 500 µm X-ray spot, and the pressure in the analysis chamber was approximately 3 × 10-10 mbar. The hydrocarbon C1s line at 284.8 eV from adventitious carbon was used for energy referencing. 3. Results and Discussion The Li-rich layered materials LL-811 and 0.5Li2MnO3·0.5LiNi1/3Co1/3Mn1/3O2 (LL-111) were synthesized via a co-precipitation process. Electron-probe X-ray microanalysis (EPMA) was performed to determine the chemical composition of LL-811 and LL-111. The ratio of TMs in each sample was similar to the designed stoichiometry (Figure S1). The morphologies of LL-111 and LL-811 are demonstrated in Figure 1a and 1d, respectively, and the cross-section images of these materials are shown in Figure 1b and 1e. These images clearly show that the secondary particles of LL-811 and LL-111 consisted of close-packed primary particles, and that the structure of LL-811 was much denser than that of LL-111. LL-811 also exhibited a higher tap density (2.3 g cm-3) than LL-111 (2.1 g cm-3). To determine the surface and bulk structures of the materials, both LL-811 and LL-111 were analyzed by Raman spectroscopy and powder X-ray diffraction (XRD). Two major Raman active modes of LL-111were observed at 485 (Eg) and 595 cm-1 (A1g), which match well with the Eg and A1g modes of layered TM oxides LiMO2 structure (Figure S2).50-51 With increasing Ni content, the intensity of the Raman spectra of LL-811 decreased drastically, which is attributed to the reduction of the rhombohedral distortion. Furthermore, the Raman spectrum of LL-811 showed three major peaks at 485 (Eg), 550 (A1g) and 600 (A1g) cm-1. The two A1g modes of the


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LL-811 peaks may ascribe to presence some nanocomposite structures in the entire material. The A1g mode appeared at lower frequencies compared to those in the Raman spectra of other rocksalt compounds, which is related to the decrease in the bond covalence inside the layers.52 To check closely the structural difference between LL-111 and LL-811, their powder X-ray diffraction (XRD) patterns are shown in Figure 1c and 1f and their individual Rietveld refinement are summarized in tables S1 and S2. The results show that the pristine material crystallized in the R3തm space group; several additional weak peaks located in the 20-25° range are attributed to super-ordering in the monoclinic phase of this Li-rich material.53-54 With increasing nickel content in the materials the lattice constants a and c increased from 2.84(2) and 14.18(2) Å (LL-111) to 2.86(2) and 14.20(2) Å (LL-811), respectively. Furthermore, the cation mixing increased from 1.77% (LL-111) to 7.7% (LL-811) according to the refined results. To gain further insight into the atomic structure of LL-111 and LL-811 pristine materials, we used annular-bright-field (ABF) imaging and high-angel annular dark field (HAADF) imaging (Figure 2). ABF imaging in aberration-corrected scanning transmission electron microscopy (STEM) is a powerful technique for simultaneously imaging both heavy and light elements.35, 38, 55

Figure 2a clearly shows a high degree of crystallization and a perfect layered structure of

pristine LL-111 along the [010] orientation. Different from the perfect layered structure of LL111, cation mixing is clearly observed in pristine LL-811 along the [110] orientation (Figure 2d). To more clearly observe the cation mixing phenomenon, we collected HAADF images of the pristine materials. In Figure 2b and 2c, little Ni/Li ion exchange is evident on the surface of LL111 no TM ions are observed in the Li slab in the bulk of the material unlike the images of LL111, extensive cation mixing is observed in LL-811 (Figure 2e and 2f), analogous to the XRD refined and magnetic measurement results (Figure. S3).


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The electrochemical performances of LL-111 and LL-811 were tested in lithium half-cells, which were cycled between 2.0 and 4.7 V. The initial discharge capacities at a current density of 0.05 C were approximately 250 mA h g-1 for LL-111 and LL-811, respectively (Figure. S4). Both the LL-111 and the LL-811 electrodes exhibited a typical Li-rich profile with a wide, flat voltage plateau at approximately 4.5 V and large initial irreversible capacity. The voltage plateau at approximately 4.5 V is related to the activation of entire material, and the large coulombic inefficiency is caused by irreversible loss of oxygen from the material during the entire material that is being activated. In the first charge process, LL-811 delivered a specific capacity of 120135 mA h g-1 at potentials below 4.45 V, which is attributed to the oxidation of Ni2+→Ni4+ and Co3+→Co4+ in 0.5LiNi0.8Co0.1Mn0.1O2.56 When the charge potential for the Li/LL-811 half-cell was sequentially increased from 4.45 to 4.7 V, more than 200 mA h g-1 was obtained at the voltage plateau of approximately 4.5 V, which means more lithium ions were extracted from the 0.5Li2MnO3 component. The charge profile further confirms that the chemical component of LL811 is 0.5Li2MnO3·0.5LiNi0.8Co0.1Mn0.1O2.40 Cycling stability was tested between 2.0 and 4.7 V at 0.2 C and at room temperature 25 °C after activation at 0.05 C for five cycles. Compared with LL-111 (60.7% capacity retention after 100 cycles), the LL-811 electrode exhibited excellent cycle stability, with 92% capacity retention after 100 cycles and 82% capacity retention after 200 cycles (Figure 3a). In addition, the rate performance further highlights the advantage of the high-nickel-content Li-rich material LL-811, especially under high current densities such as 0.5 C, 1 C and 2 C, as shown in Figure 3c. The thermal stability of layered oxides cathode materials and safety of lithium ion batteries are important concerns for practical applications. The differential scanning calorimetry profile of LL-111 and LL-811 electrode material charged to 4.7 V in the presence of electrolyte shown in


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Figure. S5. The result suggests that the thermal stability is improved with increase of nickel content in Li-rich Mn-excess cathode materials. The most outstanding performance feature of the high-nickel-content Li-rich material LL-811 was its low voltage decay after extensive cycling (table S3). Figure 3b, 3d, and 3e show the average voltage versus cycle number and the long-term cycling charge-discharge profiles for LL111 and LL-811, respectively. The voltage deteriorated approximately 700 mV between the first and the 100th cycles for LL-111, whereas the voltage decayed only 201 mV from the first cycle to the 200th cycles for LL-811. The same phenomenon was also observed from the cycle voltammetry (CV) and dQ/dV plots for LL-111 and LL-811 (Figure S6). In agreement with the results of previous studies, LL-111, a typical Li-rich material, suffered from voltage decay and, after long-term cycling, the main contribution to its capacity occurred in a low-voltage region. This result confirms that the typical Li-rich material suffers from structural rearrangement and that a spinel-like phase forms during repeated charge-discharge processes. As a result, for typical Li-rich materials, the specific energy output of the battery decreased upon cycling because of the poor cycle stability (low specific capacity retention) and serious average voltage decay (specific energy = specific capacity × average voltage) of the LL-111 cathode material. In contrast to LL111, the charge-discharge profile of LL-811 remained almost unchanged even after 200 cycles and exhibited a charge-discharge profile typical of layered TM oxide materials. The specific capacity retention of LL-811 after 100 cycles was 92%, and the specific energy retention of LL811 after 100 cycles was 89% (Figure S7). In addition, another important factor that affects the average voltage is the increases impedance with cycling which not only decreases the average discharge voltage, but also increases the charge voltage. So, the voltage fade caused by impedance growth can be reflected in the difference between charge and discharge average


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voltage as a function of cycling (Figure 3e).57 The charge voltage increases approximately 200 mV between the second and the 100th cycles for LL-111 and 100 mV from the second cycle to the 200th cycles for LL-811. This certified that large part of voltage decay is caused by the growth of impedance rather than the layer to spinel-like structure evolution for LL-811 electrodes. These impressive cell performances of LL-811 indicate that high nickel content could overcome the intrinsic voltage decay exhibited by Li-rich material. These results suggest that Lirich materials with high nickel content, such as LL-811, demonstrate excellent structural stability during prolonged cycling, which enables the practical application of Li-rich materials in highenergy-density lithium-ion batteries. Given that the peer works have revealed the origin of the voltage decay in typical Li-rich materials, we here further disclose the underlying reason why the high-nickel-content Li-rich material LL-811 can mitigate serious voltage decay. To investigate the structural evolution of LL-811 electrodes, we performed in situ XRD at various stages of the first delithiation and lithiation processes (Figure 4a). Except for three strong diffraction peaks (38.5, 44.7 and 65.1°) originating from the Al current collector, all of the peaks can be indexed as a layered structure with R 3ത m symmetry. During the delithiation and lithiation processes the diffraction peaks slightly change their positions with the exception the (113) peaks. To reveal the phase transition clearly, plots of the intensities of the (003) and (104) peaks as a function of time are presented in Figure 4b because the intensities of these peaks reflect the TM ion migration and because (003)/(104) intensity ratio reflects the degree of cation mixing. Namely, the intensity of the (003) peak decrease when TM ions migrate out of TM slab in this structure. In the slope region of 0 h < time < 8 h, with the cell charged from 3.2 to 4.45 V, the intensities of the (003) and (104) peaks decreased slightly, which indicates that a few TM ions


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migrated out of the TM slab to support the structure. This result agrees well with those of previous reports that indicated the structure of the Li-rich materials is robust when they are charged at potentials less than 4.5 V. In the wide, flat voltage plateau region of 8 h < time < 20 h, where the voltage is approximately 4.5 V, the (003) peak intensity rapidly decreased, corresponding to TM ions migrating out of TM slab to support the structure. The (104) peak intensity decreased in the voltage plateau region, suggesting that the migrated out TM ions have a propensity to occupy the intermediary tetrahedral sites. At the end of the charge region, as the cell was charged from 4.5 to 4.7 V, the intensities of both the (003) and(104) peaks increased, indicating that some TM ions occupying the intermediary tetrahedral sites migrated into the TM slab to stabilize the structure and occupied the vacancies generated by the extraction of lithium ions. Meanwhile, some TM ions migrate into the Li slab and occupy the lithium sites. This process might suppress the structural collapse of the material at a deeply charged state. During the lithiation process, the increased intensity of the (003) peak and decreased intensity of the (104) peak suggested the reverse migration of the TM ions. This in situ result agrees with the (003)/(104) peak intensity ratios obtained from the ex situ XRD results (Figure S8). Moreover, the migration of TM ions during the delithiation and lithiation processes was also confirmed by the change of the crystal parameters, as shown in Figure. 4c. During charging from 2 V to 4.5 V the parameters a and b decreased, likely as a consequence of TM ions with a high oxidation state shrinking the M-O bonds in MO6 octahedral. Meanwhile, the increase in lattice parameter of c was caused by the increased repulsion force between the two neighbouring oxygen layers when Li ions were extracted from the Li layers. During the plateau region at approximately 4.5 V, the TM ions at higher oxidation states could not be further oxidized; Thus, a and b remained constant and the decrease in the c parameter indicated that the migration of TM


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ions decreased the repulsion force. At the end of the charge, the sharp increase in c indicated that some TM ions returned to the TM slab to occupy a Li vacancy and maintain the layered structure when Li+ ions were extracted from the TM slab of the material. To assess the structural stability of LL-811 during prolonged cycling, an ex situ XRD pattern was collected after 50 cycles and was subsequently compared with one of LL-111 (Figure S9). The XRD pattern of the LL-811 electrode significantly differed from that of the LL-111 electrode after 50 cycles. All the layered peaks were clearly indexed in the XRD pattern of the LL-811 electrode, whereas only the To examine which types of TM ions participate in the migration process during delithiation and lithiation, we analysed LL-811 at different delithiation and lithiation states using X-ray photoelectron spectroscopy (XPS) (Figure 5). The evolution of the O 1s core spectra during the delithiation and lithiation process agrees well with previously reported spectra of Li-rich materials. The appearance of the 530.5 eV O 1s component at the plateau region of LL-811 charged to 250 mAh g-1 (Figure 5c) and to a higher voltage of 4.7 V (Figure 5d) was characteristic of the formation of peroxo-like O22- species.58-59 The complicated evolution of the Ni 2p3/2 core spectra during the delithiation and lithiation processes was caused by the multiple splitting of the energy levels of nickel oxides, verifying that nickel ions were the main components participating in the migration process. In the pristine material, the main photo peaks at 853-857 eV correspond to Ni2+ and Ni3+ and the peak at approximately 860 eV is commonly assigned as a satellite peak of nickel oxides. During delithiation the new nickel peak centred at 851 eV is observed; this peak is mainly caused by the nickel ions at the intermediary tetrahedral sites. During lithiation, the intensity of this peak decreased, indicating that some nickel ions migrated back to the TM slab. The most significant change in the Ni 2p3/2 core spectra was observed at the satellite peak. Upon the delithiation and


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lithiation, the intensity of this satellite peak gradually increased, suggesting that the nickel electronic state changed. The increase in the satellite peak intensity was triggered by the electrons of the Fermi sea filling 4s holes and leaving 3d9 orbitals empty. The empty 3d9 orbitals improved the O2p hybridization with TM 3d states, which improved the structural stability during the extended oxygen-activation plateau60. The authors of previous studies of Li-rich materials have emphasized that one reason for the structural instability and the voltage fade in these materials is the Jahn-Teller effect of active Mn3+ and the irreversible migration of manganese ions.61-62 In the high-nickel-content Li-rich material prior redox of Ni4+/Ni2+ kept the average oxidation state of Mn above +3. Therefore the Jahn-Teller effect of active Mn3+ ions was dimished and the structure was stabilized, as inferred from the unchanged Mn 2p3/2 core spectra. Given the in situ XRD and ex situ XPS results, we believe that Ni played a positive role in improving the structural stability and mitigating voltage decay in the LL-811 material (Scheme in Figure 6). The higher cation mixing, which was caused by the high nickel content, not only serve as a pillars to stabilize the structure during charging voltage below 4.5 V but also resulted in the migration of ions into the TM layers to occupy the Li vacancies formed by the delithiation of the Li-rich phase. Furthermore, during subsequent charging in the LL-811 material nickel ions instead of manganese ions migrated out from TM layers to support the structure. Unlike manganese ions, which tend to be trapped in the intermediary tetrahedral sites, nickel ions reversibly migrated between the TM layer and the interlayer during the delithiation/lithiation process, suppressing the formation of a spinel-like phase and mitigating voltage decay. The main reasons behind this difference may arise from the relatively large ionic radius of Ni2+ ions (0.69 Å), which makes them difficult to be trapped in the intermediary tetrahedral sites. Finally, the


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high-nickel-content Li-rich material improved the O 2p hybridization with TM 3d states and inhibited the Jahn-Teller effect of active Mn3+ ions, which improved the stability of the material’s structure. 4. Conclusions In summary, cation migration, which reduces the repulsion force and suppresses structural collapse, is critical to support the layered structure of lithium-rich materials, especially at deep delithiation states. In typical Li-rich materials, such as LL-111, cation migration is not a completely reversible process; thus, numerous TM ions, primarily manganese ions, are trapped in the intermediary tetrahedral sites, resulting in the formation of a spinel-like phase and in voltage fade. However, in the high-nickel-content, Li-rich cathode material LL-811, electroactive nickel not only acts as a stabilizing agent to improve the material’s cycling stability but can also migrate between the TM layer and the interlayer to mitigate voltage decay. The demonstration of a co-precipitation-synthesized Li-rich cathode material with high nickel content will advance understanding of the voltage decay and the practical applications of high-capacity layered Li-rich cathode materials. ASSOCIATED CONTENT Supporting Information Supporting Information Available: EPMA analysis of pristine materials, Raman results of pristine materials, First charge-discharge profiles, CV and dQ/dV plots of LL-111 and LL-811, Specific capacity and specific energy as function of cycle number of LL-811 electrode, Ex situ XRD results of LL-811during the first charge discharge process, Ex situ XRD results of LL-111


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and LL-811 electrodes after 50 cycles, Rietveld Refinement results of LL-111and LL-811 pristine material, Summarized the voltage decay of advanced Li-rich cathode materials have been reported. AUTHOR INFORMATION Corresponding Author * Tel/Fax: (+86)-10-82617069, E-mail: [email protected] (Y.G.G.); [email protected] (L.G.) [email protected] (L.J. W.) Notes The authors declare no competing financial interest. ACKNOWLEDGMENT This work was supported by the Ministry of Science and Technology of the People’s Republic of China (Grants No. 2016YFA0202500, 2012CB932900), the National Natural Science Foundation of China (grant numbers 51225204, 21303222, and 21127901), and the “Strategic Priority Research Program” of the Chinese Academy of Sciences (grant number XDA09010000). REFERENCES (1).

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Figure 1. (a) Surface SEM image and (b) cross-sectional SEM image of LL-111. (c) XRD pattern with Rietveld refinement of LL-111 (d) Surface SEM image and (e) cross-sectional SEM image of LL-811. (f) XRD pattern with Rietveld refinement of LL-811.


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. Figure 2. ABF and HAADF STEM images of pristine materials. (a) ABF image of LL-111 along the [010] zone axis. (b) Edge and bulk HAADF image of LL-111 along the [110] zone axis (c) Magnified bulk HAADF image of LL-111 along the [010] zone axis. (d) ABF image of LL811 along the [110] zone axis. (e) HAADF image of LL-811 along the [110] zone axis. (f) Enlarged HAADF image of LL-811 along the [110] zone axis, corresponding to the dashed square in figure (e).


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Figure 3. (a) Discharge capacity as function of cycle numbers between 2.0-4.7 V for LL-111 and LL-811 electrodes. (b) Charge-discharge profiles between 2.0-4.7 V for 5-100 cycles at 0.2 C and the arrow indicates the voltage fade of traditional Li-rich LL-111 electrode. (c) Rate performances of LL-111 and LL-811 electrodes. (d) Charge-discharge profiles between 2.0-4.7 V for 5-200 cycles at 0.2 C and the arrow indicates the voltage fade of high nickel content Lirich LL-811 electrode. (e) Average voltage as function of cycle numbers between 2-4.7 V for LL-111 and LL-811 electrodes.


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Figure 4. (a) In situ XRD pattern of LL-811 during the first charge discharge process (b) Intensity evolution of the (003) and (104) planes of LL-811 during the first charge-discharge process. (c) Lattice parameter changes of LL-811 during the first charge-discharge process.


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Figure 5. (a) Pristine material of LL-811. (b) LL-811 electrode charged to 4.4 V. (c) LL-811 electrode charged to 250 mA h g-1. (d) LL-811 electrode charged to 4.7 V. (e) LL-811 electrode discharged to 2.0 V.


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Figure 6 Schematic of Ni2+ migration in high-nickel-content nickel Li-rich material during the first charge-discharge process.


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Mitigating Voltage Decay of Li-Rich Cathode Material via Increasing Ni

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