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A Mg-doped High-nickel Layered Oxide Cathode Enabling Safer, High-energy-density Li-ion Batteries Qiang Xie, Wangda Li, and Arumugam Manthiram Chem. Mater., Just Accepted Manuscript • DOI: 10.1021/acs.chemmater.8b03900 • Publication Date (Web): 22 Jan 2019 Downloaded from http://pubs.acs.org on January 27, 2019

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Chemistry of Materials

A Mg-doped High-nickel Layered Oxide Cathode Enabling Safer, High-energy-density Li-ion Batteries Qiang Xie, Wangda Li, and Arumugam Manthiram* Materials Science and Engineering Program & Texas Materials Institute, The University of Texas at Austin, Austin, TX 78712, USA ABSTRACT: High-nickel layered oxide cathodes with a Ni content of > 90% show substantial potential for next-generation lithiumion batteries (LIBs) due to their high capacity and lower cost. However, they are plagued by rapid capacity decay and poor thermal stability, which hamper their practical viability. We present here Li0.98Mg0.02Ni0.94Co0.06O2 (NC-Mg) with 2% Mg doping, aiming to provide a strategic guideline for solving the issues. The Mg2+ ions occupy the lithium layer and are proposed to act as pillar ions, which substantially enhances the structural reversibility and reduces the anisotropic lattice distortion upon cycling, thereby greatly improving the electrochemical and thermal stability of NC-Mg compared to the undoped LiNi0.94Co0.06O2 (NC). Specifically, NC-Mg delivers 214 mA h g-1 with a capacity retention of 80.1% after 500 cycles in pouch-type full cells, much higher than the retention of NC (56.3%). A discharge capacity of 158 mA h g-1 at 10C rate demonstrates its remarkable rate capability. Additionally, the Mg doping significantly elevates the exothermic peak temperature of NC-Mg to 211 oC, in sharp contrast to 177 oC for NC, highlighting the improved thermal stability of NC-Mg. Collectively, the superior performance of NC-Mg demonstrates a feasible alternative strategy for developing safer, high-energy-density LIBs.

1.

Introduction

and Li-ion conductive materials,21-23 increase the surface robustness of cathode materials to electrolyte corrosion. Additionally, tuning of electrolyte solutions proves effective in improving the cell performance by minimizing adverse interfacial interactions between cathode and electrolyte.24-26 Previous work on high-nickel cathodes is fruitful, but mainly centers around the cathodes with circa 80% Ni or less. Through a further increase in Ni content to 90% or higher, higher capacities (~ 230 mA h g-1) can be achieved to meet the increasing energy needs of LIBs, which, however, is accompanied by significantly increased cycling and thermal instabilities. The poor cycling stability in high-nickel cathodes and LiNiO2 primarily originates from their poor structural stability involving severe anisotropic lattice variations and aggressive surface reaction with the electrolyte at deep charge.7,27 Hence, stabilizing their lattice structure is vital to improve the electrochemical and even thermal performance of high-nickel cathodes, which can be achieved by doping with alien ions. As the lattice contraction in high-nickel cathodes is mainly caused by the contraction of the lithium layer space, introduction of pillar elements to the lithium layer can benefit the lattice stability and cycling performance of high-nickel oxides.6,28 Common doping elements such as Al, Ga, and Ti can improve the performance of high-nickel oxides, but they do not function as pillar ions in lithium layer due to their preference for the transition-metal layer.5,9,12 In undoped high-nickel oxides, small amount of Ni2+ ions residing in lithium layer (i.e., Li/Ni mixing) initially act as pillars in their structures. Unfortunately, the oxidation of Ni2+ to smaller Ni3+ or Ni4+ ions upon charging and the rapid increase of Li/Ni mixing during cycling seriously compromise the pillar effect of Ni2+, destroying active lithium-intercalation sites, deteriorating lithiation kinetics and resulting in

The huge success of lithium-ion batteries (LIBs) in portable electronics has created considerable enthusiasm in extending LIBs to the field of electric vehicles (EVs). However, such extension places demanding requirements for LIBs, encompassing high specific energy (e.g., ca. 350 W h Kg-1 or higher), high safety, low cost, and environmental benignity1,2 Current LIBs based on LiCoO2, LiNi1/3Co1/3Mn1/3O2 or LiNi0.8Co0.15Al0.05O2 are incapable of meeting these requirements in view of their practical capacity, cost, and environmental issues.3 Therefore, developing high-performance, cost-effective LIBs is of tremendous significance to the development of EVs. LiNiO2 was a focus with great efforts as a cathode for next-generation LIBs due to its higher reversible capacity (~ 250 mA h g-1), lower cost and reduced toxicity,4-6 but its application was seriously hindered by the difficulty in synthesizing stoichiometric LiNiO2, rapid capacity decay, and poor thermal stability.7,8 As a substitute for LiNiO2, high-nickel layered transitionmetal oxides (LiNixM1-xO2, x > 0.6 M = Co, Mn, Al, Ti, Ga, etc.) are regarded as promising cathode materials for next-generation LIBs.5 Utilization of a high Ni content ensures high capacity while the incorporation of other elements enhances the electrochemical and thermal performance.9-11 For example, the doping of Co, Mn, and Al in high-nickel oxides benefits the cycle and thermal stability.5,9,12,13 Ga-doping greatly enhances the cycle performance and over-charge resistance of pure LiNiO2.10 Other strategies have also been utilized to refine the performance of high nickel cathodes. Concentration-gradient design with high-Ni at the interior and low-Ni at the surface takes advantage of the lower surface reactivity of non-Ni elements.14 Surface coatings, including oxides,15-18 phosphates,19 fluride,20

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poor cycling performance and large voltage polarization. Mg2+ is a promising pillaring ion due to several merits. Experimental and calculation studies indicate that Mg2+ ions tend to preferentially enter the lithium layer particularly when the dopant content is less than 2% due to the similar size to Li+ (rLi+ = 0.76 Å, rMg2+ = 0.72 Å) and lower formation energy to enter Li sites.6,29 Unlike Ni2+, Mg2+ ions are proposed to work as a stable pillar due to the unchanged size and valence upon cycling. It is worth noting that Mg2+ doping may detrimentally affect the diffusion of lithium and lower the capacity, but the impact extent can be minimized by controlling the doping amount.30 Despite the potential adverse impact, the advantages of Mg2+ ions highlight its value in improving the performance of high-nickel cathodes. In view of this, we herein report a Mg-doped high-nickel cathode Li0.98Mg0.02Ni0.94Co0.06O2 (hereafter referred to as NCMg). The utilization of a higher content of Ni (94%) is able to deliver a higher discharge capacity and decreased cost compared to the currently used LiNi0.8Co0.15Al0.05O2. Meanwhile, the incorporation of 2% Mg2+ ions is anticipated to enhance the structure, cycling, and thermal stability of NC-Mg due to a stable pillaring effect in the lithium layer without significantly affecting lithium diffusion and capacity. The doped Mg2+ ions are predominately present in the lithium layer, in-line with our expectation. In order to evaluate the effect of the doped Mg2+ ions on the performance of NC-Mg compared to the undoped LiNi0.94Co0.06O2 (hereafter referred to as NC), diverse characterization methods, including ex-situ and in-situ X-ray diffraction (XRD), redox titration techniques, high resolution transmission electron microscopy (HR-TEM), and differential scanning calorimeter (DSC), are used to systematically investigate both samples. In-situ XRD results reveal that the presence of Mg2+ ions in the lithium layer of NC-Mg alleviates the lattice anisotropic distortion and structural collapse upon charging. Electrochemical and DSC tests indicate that NC-Mg outperforms NC significantly in both electrochemical and thermal aspects. Overall, the results attest to the feasibility of our strategy, laying the foundation for the practical viability of high-nickel oxides.

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~ 5 mg cm-2, which was then dried, punched and stored in a vacuum oven at 120 oC overnight. 2032 coin cells were assembled in an argon-filled glovebox with the prepared positive electrodes, electrolyte, Celgard 2500 separator, and Li metal counter/reference electrode (MTI Corp). 1.0 M LiPF6 in ethylene carbonate (EC)/ethylmethyl carbonate (EMC) (3:7 by weight) with 2 wt.% vinylene carbonate (VC) was used as an electrolyte solution. Coin cells were cycled at C/3 rate between 2.8 and 4.4 V vs. Li/Li+ at 25 oC (1C = 180 mA g-1). Pouch-type full cells were prepared to investigate the long-term cycling stability of both samples. Each full cell consists of one prepared cathode and one mesocarbon microbead graphite anode with a N/P ratio of ~ 1.08. Full cells were cycled at C/10 rate (current density 18 mA g-1) for 3 times, followed by at C/2 rate for 500 times between 2.5 and 4.3 V at 25 oC with a constant current-constant voltage charging mode. Electrochemical impedance spectroscopy (EIS) tests of the full cells were performed on an impedance spectrometer (Solartron 1260A) in the frequency range of 1 MHz to 2 mHz with a 5 mV potential perturbation. Characterization: Chemical compositions of the precursor and freshly prepared samples were analyzed by inductively coupled plasma (ICP) analysis (Varian 715 ES). The content of residual lithium was determined by a chemical titration.1 Briefly, 1.0 g of the fresh sample was immersed into 15 mL of de-ionized water for 5 min to dissolve LiOH and Li2CO3. Afterwards, the solution was filtered and washed to collect a clear solution. The collected solution was titrated with 0.1 M HCl solution with methyl orange as indicator. The total amount of the HCl consumed at the titration endpoint was used to calculate the residual lithium content for each sample. Oxygen content values of NC and NC-Mg were determined by a redox titration, employing KMnO4. Specifically, ~ 35 mg of the fresh sample was added to 80 mL of 0.1 M oxalic acid and 1.0 M H2SO4 solution, which was then heated at 60 oC until the sample completely dissolved. The obtained solution was titrated with 5.0  10-3 M KMnO4 solution. The total amount of the KMnO4 solution consumed was used to calculate the number of electrons transferred during the reaction between oxalic acid and transition-metal ions. Based on the titration values, the oxygen content of the samples could be determined using electroneutrality principle. Sample morphology was explored with a scanning electron microscope (SEM) (FEI Quanta 650). Elemental mapping of the NC-Mg particle was examined by high-resolution transition electron microscopy (HRTEM) with energy disperse X-ray spectroscopy (EDS) (JEOL 2010F) and SEM. Focused ion beam (FIB) was used to prepare a TEM cross-section sample. The cross-section samples for SEM mapping were obtained by ion-milling the cycled samples from the pouch full cell. Structure and phase purity of the synthesized samples were studied with X-ray powder diffraction (XRD) (Rigaku Miniflex 600) (Cu Kα radiation) in the 2θ range of 10 – 80 º with a scan step of 0.02 º. Rietveld refinement was performed on Fullprof program.32 The in-situ XRD test was conducted on a Rigaku UltimaIV XRD instrument coupled with a Land cell testing system (CT2001A), in which beryllium metal acted as a cathode current collector and the cell was cycled at C/10 rate between 2.8 and 4.4 V vs Li/Li+ for the first formation cycle, followed by charging to the same state of charge (x = 0.8 in Li1xNi0.94Co0.06O2/Li0.98-xMg0.02Ni0.94Co0.06O2). The data collected in the second charge process was used to compare the structural

2. Experimental Section Materials Synthesis: The Ni0.94Co0.06(OH)2 precursor was synthesized by a co-precipitation method with the detailed process described previously.31 Briefly, 1.0 M aqueous solution of NiSO4·6H2O and CoSO4·7H2O, 2.0 M KOH solution, and concentrated NH4OH solution were continuously pumped into a batch tank reactor. The temperature, pH, and stirring speed were carefully controlled during the co-precipitation process. The synthesized precursor was washed with deionized water, filtered, and then dried at 100 oC overnight. LiNi0.94Co0.06O2 (NC) was obtained by calcining LiOH·H2O with Ni0.94Co0.06(OH)2 (molar ratio of Li : (Ni + Co) = 1.05 : 1) at 640 oC for 15 h under oxygen atmosphere. To obtain Li0.98Mg0.02Ni0.94Co0.06O2 (NCMg), LiOH·H2O, Ni0.94Co0.06(OH)2, and magnesium acetate (molar ratio of Li : (Ni + Co) : Mg = 1.05 : 0.98 : 0.02) were homogeneously mixed and calcined at 680 oC for 15 h under oxygen atmosphere. Electrode fabrication and electrochemical tests: The cathode slurries were prepared by mixing the active material, carbon black, and poly(vinylidene fluoride) (PVDF) (weight ratio = 8 : 1 : 1) in N-methyl-2-pyrrolidone (NMP). The well-mixed slurries were cast onto an Al foil with an active material loading of


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changes of both the samples. Thermal stability was evaluated by differential scanning calorimeter (DSC) (Netzch STA 449). For DSC tests, coin cells with both samples were first cycled three times between 2.8 and 4.4 V at C/5 rate for activation. In order to provide more insight into the thermal stability, DSC tests were conducted on both the samples with two charged states, namely, charging both samples (i) to the same cut-off voltage of 4.3 V and (ii) to the same state of charge (x = 0.78 in Li1-xNi0.94Co0.06O2/Li0.98-xMg0.02Ni0.94Co0.06O2) at C/10 after the activation cycles. The charged cells were disassembled in an argon-filled glovebox to collect the positive electrode, followed by washing with dimethyl carbonate (DMC) and drying inside a glovebox. The sample was scraped from the dried electrodes and sealed in a stainless crucible with ~ 35 wt.% electrolyte. DSC measurement was performed from 30 to 350 oC at a heating rate of 1 oC min-1 under argon atmosphere.

Inductively couple plasma (ICP) analysis and redox titration were used to determine the element composition and the oxygen content in the cooked samples. ICP results for the synthesized precursor and lithiated samples in Table S1 of the Supporting Information show that the elemental ratio of Ni, Co, and Mg agrees well with the purported value. The lithium content for both samples slightly departs from 1.0 due to Li volatilization during synthesis and surface reactivity with air to form residual lithium (LiOH, LiHCO3 and Li2CO3) on the surface.1 The results of residual lithium and oxygen content for both samples are given in Table S2 of the Supporting Information. As shown, both samples have the same oxygen content, but NCMg has a lower content of residual lithium. To have a better understanding of the cationic distribution, XRD test and Rietveld refinement were performed. The XRD patterns of pristine NC and NC-Mg (Figure S1, Supporting Information) clearly show that both samples have a hexagonal α-NaFeO2 structure with R3̅m space group, with no impurity phases observed. The obvious splitting of (006)/(102) and (108)/(110) peaks implies both materials have well-ordered layered structure

3. Results and Discussion 3.1. Cation distribution and morphology

Figure 1. SEM images of (a) pristine LiNi0.94Co0.06O2 (NC) and (b) pristine Li0.98Mg0.02Ni0.94Co0.06O2 (NC-Mg). (c) STEM magnified image of NC-Mg primary particle in Figure S3 (the red squared region), (d) - (f) STEM EDS elemental maps of NC-Mg primary particle in Figure S3. (g) SEM cross-sectional image of the NC-Mg secondary particle, (h) - (j) SEM EDS elemental maps of the secondary particle of (g).

The Rietveld refinement results are summarized in Figure S2, Table S3, and S4 of the Supporting Information. The calculated and experimental XRD patterns agree fairly well, suggesting that the structural models simulate the sample structures well and that the refinement results are reliable. As Table S4 shows, almost all the doped Mg2+ ions are located in lithium sites, in agreement with the results from a LiNi1-yMgyO2 system and calculations.6,29 Systematic structural and magnetic studies and calculations in the reference unequivocally showed that Mg2+ ions tend to preferentially enter the lithium layer when lithium vacancies are available and the doping amount is ≤ 2% due to their size similarity with Li+ (rLi+ = 0.76 Å, rMg2+ = 0.72 Å, and rNi2+ = 0.69 Å) and lower formation energy to enter Li sites. The presence of Mg2+ ions in the lithium layer is proposed to func-

tion as structural pillars for the high-nickel cathode at delithiated states at high voltages during cell operation, mitigating large anisotropic lattice variations and substantially enhancing structural reversibility.33 Moreover, the comparison of cation occupancy in Table S3, S4 reveals that the Li/Ni mixing in pristine NC-Mg (2.3 %) is slightly higher than that in pristine NC (1.8 %), consistent with the XRD patterns (Figure S1) and the intensity ratio of (003)/(104) peaks, an indicator of the degree of Li/Ni mixing (Table S2).34 The slightly higher Li/Ni mixing in NC-Mg may be related to the presence of Mg2+ ions in the lithium layer. Despite a smaller atomic number (Z) relative to the transition metal ions, Mg2+ ions are still able to reflect certain amount of X-ray, leading to a lowered (003)/(104) ratio in NC-Mg, namely, a slightly increased Li/Ni mixing. Note that the slightly increased degree of Li/Ni mixing for the Mg-doped



sample does not greatly affect its rate performance (see below). The morphologies of both samples were examined by SEM and are shown in Figure 1a and b. As shown, both samples display a spherical morphology with a diameter of ~ 12 μm. HR-TEM images of NC-Mg primary particle and the corresponding elemental mapping in Figure S3 and Figure 1c - f as well as the EDS elemental mapping for the secondary particle in Figure 1g - j clearly demonstrate the homogeneous distribution of Ni, Co, and Mg in the NC-Mg particles 3.2. Structural and phase evolution In order to examine the impact of Mg doping on the structural changes and phase evolution during charge and discharge, in-

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situ XRD characterization was conducted on both samples. Figure 2a presents the contour plot of the (003) to (113) Bragg reflection peaks for both samples and their corresponding voltage – state of charge (SOC) curves. As plotted in Figure 2a, both samples display a continuous hexagonal phase transition (H1 → H2 → H3) upon charging, as was similarly observed in LiNiO2.35 Specifically, the (003)H1 peak smoothly shifts toward a lower 2θ upon charging and evolves into (003)H2 completely at ~ 40% SOC (~ 4.1 V). Upon further charging, (003)H2 peak shifts backward to a higher 2θ, accompanied by the evolution of H2 to H3 phase.

Figure 2. In-situ XRD of LiNi0.94Co0.06O2 (NC) and Li0.98Mg0.02Ni0.94Co0.06O2 (NC-Mg) charged at C/10 rate to the same state of charge (SOC). (a) Contour plot of the diffraction patterns from (003) to (113) peaks and the corresponding voltage – SOC curves, (b) and (c) the corresponding evolution of the lattice parameters a and c.

metal-oxide upon charging.33 For NC-Mg, the presence of Mg2+ ions in the structure impedes both the lithium/vacancy ordering and electronic redistribution,33 suppressing the two-phase separation and leading to a pseudo single-phase transition. The evolution of phases is strongly correlated to the change in the c and a lattice parameters upon charging and the related results are shown in Figure 2b and c. The lattice parameter a of

While both samples show the phase transition from H2 to H3 at the end of charge, they behave differently. NC exhibits a twophase separation transition while NC-Mg experiences a pseudo single-phase transition, due to the presence of Mg2+ in the lithium layer. The two-phase separation for NC is caused by the lithium/vacancy ordering in lithium layers, which is sensitive to dopants, and the electronic redistributions in the transition


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both samples, characteristic of the Ni-Ni distance, decreases continuously upon charging due to the oxidation of larger Ni 2+ ions and smaller Ni3+ ions. The a-axis contraction for both samples is around 2%. In comparison, the lattice parameter c of both samples increases first and then decreases. The increase in c parameter at the early stage of charging is attributed to the increase of the interslab distance. With Li+ ions extracted from the structure, the electrostatic repulsion between the Ni(Co)O2 slabs intensifies, leading to an increase in the interslab distance. On the other hand, the abrupt decrease in the c parameter at highly charged states is caused by the significant decrease in the interslab distance due to two factors. The increased covalency of the Ni4+-O2- bonds in the Ni(Co)O2 slabs lowers the electrostatic repulsion between the Ni(Co)O2 slabs and the negligible Li content in the lithium layer removes the pillaring effect, contributing to an abrupt decrease in the interslab distance.33,36,37 The difference in the increase of the c parameter at the early stage of charging between the two samples is not conspicuous. However, there is a stark distinction between the shrinkage of the c parameter for the samples at the end of charge. The c parameter shrinkage in NC (5.6 %) is much higher than that in

NC-Mg (3.7 %) at the end of charge (80% SOC). The remarkably reduced contraction in NC-Mg is believed to be caused by the pillaring effect of Mg2+ ions present in the lithium layer, which prevents the drastic lattice contraction and local collapse and stabilizes the lattice structure at deeply charged states, well corresponding to the calculation results of Min et al.29 Besides, since Ni2+ ions are oxidized to smaller Ni4+ ions at high voltages, the larger Mg2+ ions can provide a much more stable pillaring effect than Ni4+ in the lithium layer.34 The impaired pillar effect and structural instability in NC causes the migration of more nickel ions from the transitional-metal layer to the Li layer during cycling, resulting in a rapid increase in Li/Ni mixing.34,38 In short, in-situ XRD results reveal that Mg doping in the lithium layer of NC-Mg smoothens the transition of H2 phase to H3 phase, enhances the structural reversibility with a decreased Li/Ni mixing during cycling (shown below), and inhibits the large anisotropic lattice distortions that cause local structure to collapse during repeated charge-discharge cycles, all of which are beneficial to improving the performance of high-nickel, low cobalt cathode materials.

Figure 3. Electrochemical performance comparison of LiNi0.94Co0.06O2 (NC) and Li0.98Mg0.02Ni0.94Co0.06O2 (NC-Mg). (a) Cycling performance, (b) Coulombic efficiency, (c) and (d) charge-discharge profiles, (e) and (f) evolution of the average charge/discharge voltage (calculated by energy/capacity) over 150 cycles and (g) rate performance comparison in coin cell (vs Li/Li+). (h) Comparison of the long-term cycling performance of both samples in pouch-type full cells (vs graphite anode, 1C = 180 mA h g-1), i) corresponding average voltage for full cells.

the slightly increased Li/Ni mixing and the presence of Mg2+ ions in NC-Mg. The capacity of both samples reaches their pinnacle after around 12 cycles, 214 mA h g-1 for NC-Mg and 226 mA h g-1 for NC. In contrast to its decreased capacity, NC-Mg

3.3. Electrochemical performance Figure 3a depicts the cycle performance of both samples in half cells between 2.8 and 4.4 V. NC-Mg displays a decreased initial discharge capacity compared to NC, likely resulting from



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ions. Despite the presence in lithium layer, the doped Mg2+ ions do not significantly degrade the rate performance of NC-Mg, and it still delivers a discharge capacity of 158 mA h g-1 at 10C rate (corresponding to 73% of that at C/5 rate), demonstrating a good rate performance. The practical application of LIBs calls for the long-term stability of cathode materials paired with graphite anode in full cells. In order to evaluate the long-term stability of both samples, pouch full cells with both samples were assembled and cycled between 2.5 and 4.3 V at C/2 rate for 500 cycles. The results are presented in Figure 3h and i, averaged from two pouch cells for each sample. As Figure 3h and i show, the capacity and voltage retention after 500 cycles for NC-Mg is, respectively, 80.1% and 95.6%, much higher than 56.3% and 92.9% for NC. The excellent long-term cycling stability of NC-Mg in full cells is consistent with the results of half cells. Figure 4a - d compares the SEM cross-section images of NC and NC-Mg after 200 cycles in pouch full cells, and the crosssection images of Ni0.94Co0.06(OH)2 precursor and the freshly pristine NC and NC-Mg are also shown in Figure S5. The comparison of Figure 4a and c with Figure S5 shows that some microcracks are formed in the center part of the secondary particles of both cycled samples. Further examination on Figure 4a reveals that, apart from the cracks, lots of holes are also formed along the exterior of the secondary particle of the cycled NC as the dashed blue line indicated (more images see Figure S6), while those holes are not observed in the pristine materials and precursor (Figure S5). These dense holes in NC, probably caused by the repeated dramatic anisotropic lattice changes upon cycling as in-situ XRD results show, result in the separation of primary particles and even the delamination in secondary particles (Figure 4b), which adversely affects the electronic contact.31 Besides, the holes and delamination, along with the cracks, in NC provide more fresh surface where the unwanted electrode-electrolyte reactions happen. The decreased electronic contact and increased reactions could lead to the growth of impedance and polarization and the rapid capacity decaying, consistent with the electrochemical results above.

outperforms NC in capacity retention after 150 cycles, 90% for NC-Mg and 77% for NC. Figure 3b shows the Coulombic efficiency versus cycle number. As seen, although the first cycle Coulombic efficiency of NC-Mg is slightly lower than that of NC, NC-Mg maintains higher Coulombic efficiency than NC during subsequent cycles. Figure 3c and d compare the evolution of charge-discharge curves for both samples. The first voltage curves for both samples are similar and the plateaus, characteristic of the H1, H2, and H3 phase transitions, are obvious in both curves. After 150 cycles, however, the voltage plateaus for NC become indistinguishable while those for NC-Mg are still distinguishable, implying that NC may suffer from serious irreversible structural changes upon cycling. The improved cycling performance of NC-Mg can be attributed to its superior structural stability compared to NC. Specifically, in-situ XRD results clearly indicate that NC suffers from severe lattice anisotropic contraction upon charging. The severe anisotropic contraction during cycling can induce the formation and propagation of micro-cracks which damage the material integrity, exposing fresh surfaces that produce further unwanted electrode/electrolyte reactions, as well as the accumulation of inactive species, causing rapid capacity decay and impedance growth in NC.27,39-43 In addition, the stabilizing pillar effect of Mg2+ enhances the structural reversibility of NCMg cathode and inhibits Li/Ni mixing during cycling, thereby considerably improving the cycling stability of NC-Mg over that of the undoped sample. Figure 3e and f display the average charge-discharge voltages for both samples. The difference between the charge and discharge voltages increases slightly from 0.06 to 0.11 V after 150 cycles for NC-Mg. In contrast, the difference for NC increases from 0.11 to 0.25 V, indicative of a larger impedance growth during cycling. Figure 3g compares the rate performance of both samples from C/5 to 10C rate in coin cells cycled between 2.8 and 4.4 V at 25 oC. Both samples present a similar capacity retention from C/5 to 5C rate, but the capacity retention of NCMg is slightly lower than that of NC at 7C and 10C rates, likely due to the initially higher Li/Ni mixing and the impact of Mg2+

Figure 4. SEM cross-section images of (a-b) LiNi0.94Co0.06O2 (NC) and (c-d) Li0.98Mg0.02Ni0.94Co0.06O2 (NC-Mg) after 200 cycles in pouch full cell. The XRD patterns comparison of pristine and cycled samples for (e) NC and (f) NC-Mg.


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XRD patterns comparison of both samples before and after 200 cycles in pouch full cell (Figure 4e and f) indicates that both samples do not show significant crystal structure changes before and after cycling, and the layered structure for both samples is well retained. However, the Li/Ni mixing for both samples increases as evidenced by the Rietveld refinements (Figure S2 and Table S2). Indeed, the Li/Ni mixing in NC and NC-Mg increases, respectively, from 1.8% to 7.8% and 2.3% to 4.0% after 200 cycles. The higher Li/Ni mixing for NC is due to the impaired pillaring effect and structural instability of NC upon cycling,34,38 The Li/Ni mixing result agrees well with the in-situ XRD and electrochemical results discussed above and also confirms the calculation results by Min et al. that Mg doping can inhibit the migration of Ni to Li sites.29

Figure 5. Nyquist impedance plots of the pouch-type full cells with both the samples cycled between 2.5 and 4.3 V at C/2 rate. (a) LiNi0.94Co0.06O2 (NC) and (b) Li0.98Mg0.02Ni0.94Co0.06O2 (NC-Mg).

Figure 5 depicts the impedance spectra of both samples in pouch full cells. The data were obtained at 1 st, 100th, and 150th cycles after three formation cycles. The first semi-circle in the high-medium frequency range is identified as surface film resistances (Rsf), resulting from multiple RC circuits including solid-electrolyte interphase (SEI) at the anode, the cathode surface film, and separator resistances, etc, while the second semicircle at medium-low frequency corresponds to the chargetransfer resistance (Rct) of cathode.13,44,45 The diameter of the semi-circle reflects the value of each resistance. As Figure 5 shows, the diameter of the second semi-circle of NC is almost twice that of NC-Mg at the 1st cycle, implying that NC has a larger resistance than NC-Mg even if only cycled a few times (185  cm2 versus 360  cm2). After 150 cycles, the Rct of NC increases from 360  cm2 to 430  cm2 but the increase in the Rct of NC-Mg is negligible. The difference in the evolution of resistance for both samples agrees fairly well with the electrochemical results, and conforms to the in-situ XRD results, SEM cross-section images as well as the XRD refinement results. The suppressed unwanted electrode-electrolyte reactions due to

mitigated particle cracking and separation with less lattice distortion at deeply charged state and the reduced Li/Ni mixing during cycling explain the much lower increase of impedance in NC-Mg upon cycling. 3.4. Thermal performance Safety is of paramount importance to high-performance LIBs, which strongly depends on the thermal stability of cathode materials. Differential scanning calorimeter (DSC) provides an effective approach to evaluate the thermal stability of cathode materials. Figure 6 plots the DSC profiles of both samples charged to 4.3V in the presence of electrolyte. The exothermic peak temperature and total released heat is, respectively, 177 oC and 1780 J g-1 for NC and 211 oC and 1640 J g-1 for NCMg. The exothermic peak temperature of NC-Mg (211 oC) is slightly higher than the reported temperature for LiNi0.8Co0.15Mn0.05O2 (205 oC) and LiNi0.85Co0.11Al0.04O2 (201 o C) and close to that of LiNi0.8Co0.1Mn0.1O2 (215 oC) at the same cut-off voltage.46,47 Besides, the DSC results of both the samples charged to the same state of charge/delithiation are also compared in Figure S7 and Table S5. All of these results further indicate that NC-Mg has a better thermal stability than NC, which could be attributed to the following reasons. The increased exothermic peak temperature and the reduced heat release from the burning of oxygen with the electrolyte reflect the enhanced resistance of the charged NC-Mg to oxygen release at elevated temperatures compared to the charged NC, as suggested by Ceder et al.48-50 The highly delithiated layered NC (LixNi0.94Co0.06O2 (x < 0.5)) will experience structural transformation to form spinel/rocksalt products (LiM2O4, M3O4, MO, M = Ni/Co) accompanied by the oxygen release at elevated temperatures.48 However, such a process strongly requires the migration of Ni/Co ions from the octahedral sites of the TM layer to the octahedral sites of the Li layer at elevated temperatures.48 Recent investigations showed that Mg-doping can stabilize the oxygen atom and inhibit Ni migration in the highly delithiated high-Ni cathodes,29 which also agrees well with the Li/Ni mixing results of the cycled samples in our study. Therefore, for the highly charged NC-Mg, a higher temperature is required to initiate the phase transition and oxygen release compared to the charged NC. Meanwhile, when the highly delithiated NC-Mg experiences phase transformation at higher temperature, it could form MgO besides the products as decomposed by the charged NC and consequently reduce the release of oxygen to burn electrolyte, which may be responsible for the reduced heat release in NC-Mg. Overall, the DSC results demonstrate the improved thermal stability of NC-Mg, despite a higher nickel content than LiNi0.8Co0.1Mn0.1O2.

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Figure 6. DSC profiles of LiNi0.94Co0.06O2 (NC, Orange) and Li0.98Mg0.02Ni0.94Co0.06O2 (NC-Mg, Green) charged to 4.3 V.

4. Conclusion A Mg-doped cathode Li0.98Mg0.02Ni0.94Co0.06O2 (NC-Mg) with a high nickel content of 94% offers a high capacity of 214 mA h g-1. The doped Mg2+ ions in NC-Mg greatly enhance the structural reversibility and alleviate the large anisotropic lattice distortions occurring at deep delithiation due to the proposed pillar effect of Mg2+ ions in the lithium layer. The improved structural stability significantly benefits the electrochemical and thermal performance of NC-Mg compared to the undoped LiNi0.94Co0.06O2 (NC). Specifically, NC-Mg shows a capacity retention of 80.1 % after 500 cycles in pouch full cells with graphite anode, much higher than the 56.3% for NC. DSC tests also highlight the improved thermal stability of NC-Mg over that of NC. The exothermic peak temperature of NC-Mg is 211 o C while it is only 177 oC for NC. Altogether, the results demonstrate the remarkable advantage of introducing an appropriate dopant like Mg2+ ion into the lithium layer to improve the performance of high-nickel oxides. We believe that the strategy presented here combined with other approaches such as electrolyte tuning and surface coating can facilitate the viability of high-nickel cathodes (Ni > 90%) in safe, high-energy-density, long-life lithium-ion batteries.

ASSOCIATED CONTENT Supporting Information. The supporting information includes XRD patterns, XRD Rietveld refinement results, ICP results, SEM cross-section images, and DSC results. This material is available free of charge on the ACS Publication website.

AUTHOR INFORMATION Corresponding Author * (A. M.) E-mail: [email protected]

ACKNOWLEDGMENTS This work was supported by the Assistant Secretary for Energy Efficiency and Renewable Energy, Office of Vehicle Technologies of the U.S. Department of Energy through the Advanced Battery Materials Research (BMR) Program (Battery500 Consortium) award number DE-EE0007762.

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