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Reconstructing the Surface Structure of Li-Rich Cathode for High-Energy Lithium-Ion Batteries Jianming Fan, Guangshe Li, Baoyun Li, Dan Zhang, Dandan Chen, and Liping Li ACS Appl. Mater. Interfaces, Just Accepted Manuscript • DOI: 10.1021/acsami.9b02827 • Publication Date (Web): 09 May 2019 Downloaded from http://pubs.acs.org on May 9, 2019

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ACS Applied Materials & Interfaces

Reconstructing the Surface Structure of Li-Rich Cathode for High-Energy Lithium-Ion Batteries Jianming Fan,a, b Guangshe Li,a Baoyun Li,a Dan Zhang,a Dandan Chen,a Liping Li*a a State key laboratory of inorganic synthesis and preparative chemistry, Institute of Chemistry, Jilin University, Changchun 130012, People’s Republic of China b College of Chemistry and materials, Longyan University, Longyan 364012, People’s Republic of China * Email: [email protected]; [email protected] Abstract: Reconstructing a favorable surface layer could contribute to superior charge transfer and stabilize bulk structure, and thus achieve the excellent electrochemical performance of lithium- and manganese-rich oxides, but is still challenging. In this work, the surface structures of Li-rich oxides have been successfully reconstructed via a facile strategy utilizing hydrothermal glucose carbonization and subsequent reduction procedure. Surface microstructure and chemical state analysis reveal that the reconstructing process involves roughing for the surface connects with the extraction of lithium ions and the reduction of Mn ions as well as the formation of spinel phase due to the distortion of oxygen anion or the presence of oxygen deficiency. The reconstructed Co-free Li-rich oxide by using 0.025 g glucose exhibits superior electrochemical performance. Its maximum discharge capacities are 237 and 193 mAh/g respectively at 100 and 600 mA/g, and the corresponding capacity retention ratios are higher than 93% at 100th cycle. Furthermore, reconstructing the surface structure also enhances the discharge capacity and cycling performance of Co-contained Li-rich cathode. The findings in present work would offer hints for surface structure reconstruction of many oxides used in energy and other fields. Key Words: Li-rich, surface post-treatment, reconstruction, spinel phase, cycling performance

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1. Introduction On the list of commercial energy storage equipment for mobile devices, lithium-ion battery (LIB) is pretty high up for its reliable energy density and cycle life.1-2 However, to realize the power electric vehicles with a 300-plus mile range, it requires further improvements of the components in LIB, particularly cathode materials. In the past decades, the search interests for cathodes with high energy density, low cost and creditable safety have been immensely focused on lithium- and manganese-rich oxides (LMROs).3-4 It has been reported that both monoclinic C/2m and trigonal R3m symmetry always coexist in such materials, thus debates about solid solution or composite structure are far from settled down.5-7 Anyway, the complex combination of the two phases in LMROs can easily produce a distorted lithium-ion diffusion pathway, leading to detrimental rate-capability.8-9 Moreover, inevitable phase transformation in electrochemical cycling could also cause a rapid decay of capacity and voltage.10 If LMRO is used as a cathode material, two important issues related with particle surface must be solved, i.e. unfavorable lithium-ion channel and unstable surface structure.11 Therefore, much effort has devoted to optimize their surface properties.12-14 Duo to the positive effect on specific capacity and cycling performance, surface modification, including artificial coating layers and the post-treatments, is being intensively explored by lots of researchers still.15 Coating materials like carbon and reduced graphene oxide could relief the side reactions in batteries and activate the particle surfaces of cathodes.16-19 Analogously, particle surface layer could be modified by the treatments via applying (NH4)2SO4, HNO3 or other acid reagents.20-22 It should be noted that surface structure of those LMROs after the above modifications was always reconstructed to give a spinel structure as revealing by many experimental detections. Such reconstruction of surface layer would probably contribute to superior discharge capacity and initial coulombic efficiency. The advantages of introducing spinel phase in Li-rich cathodes have been confirmed in many ways, like building spinel/rock salt tunnels, directly synthesizing layered-spinel


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integrated materials and so on.23-26 Recently, Li-rich cathode with a phase-gradient outer layer with “layered-coexisting phase spinel” structure had been reported to exhibit a pronounced cycling stability and improved rate capability.27 Similarly, an ultrathin spinel layer has been demonstrated to appear on Li-rich cathode particles after Polyvinylpyrrolidone coating and high temperature heating, and thus to offer a Li-ion diffusion highway in the interface.28 In the Zhang’s work, spinel Li4Mn5O12-coated lithium-rich layered oxides that were synthesized by controlled KMnO4 oxidation on the oxide precursor exhibited improvements of capacity retention and kinetic property.

29

Therefore, constructing a spinel layer on LMRO

particles should be a practical strategy to level up the electrochemical performance.13 What’s more, developing more convenient modification strategies is vital for lowering the cost and large-scale production. Herein, we employed a feasible post-treatment strategy via a combination of hydrothermal carbonization and reduction process to introduce spinel phase on the surface of Co-free LMROs. Comparing with the pristine material, the treated sample exhibited superior electrochemical properties, especially cycling performance. The structural and elemental analysis by transmission electron microscopy and X-ray absorption spectrum characterizations confirmed that spinel phase is formed on the surface of oxide particles and contributes to the improved performance of Co-free LMRO cathodes. The hydrothermal post-treatment strategy also improved the discharge capacity and cycling performance of Li1.2Mn0.54Ni0.13Co0.13O2, indicating a comprehensive effect of hydrothermal post-treatment strategy to enhance the electrochemical cycling performance of LMROs.

2. Experimental Section 2.1. Sample preparation Pristine sample: Pristine Co-free Li-rich layered cathode material was synthesized by a sol-gel method combined with high-temperature sintering. Detail procedures could be

described

as

following:

2.9413

g

Mn(CH3COO)•4H2O,

0.9954

g

Ni(CH3COO)•4H2O, 2.4484 g Li(CH3COO)•2H2O and 4.2028 g critic acid were



dissolved into 200 mL H2O and then continually stirred to form a transparent solution. Subsequently, the pH value of above solution was adjusted to 8 ~ 9 by ammonia (analytic grade). The solution was kept under water bath at 80 oC to form a transparent gel. The obtained gel was dried at 180 oC. After being thoroughly grinded, the dried powder was sintered at 500 oC for 4 hours and then at 900 oC for 12 hours. The as-prepared powder was named as sample P with cation molar ratio of Li:Mn:Ni=1.03:0.56:0.2 as determined by ICP. Post-treated samples: 0.5 g of the sample P was first dispersed into 60 mL aqueous solution containing required amount of glucose. The mixture was then transferred to 100 mL Teflon reaction vessel, and heated at 200 oC in an oven for 3 hours. After the hydrothermal treatment, the obtained powders were dried at 80 oC to eliminate residual water. According to the usage of glucose amounts, i.e. 0, 0.025 and 0.05 g, three treated samples were named as H0, H1 and H2, respectively. 2.2. Sample characterization Powder X-ray diffraction (PXRD) patterns of samples was recorded on a Rigaku Miniflex apparatus using a Cu source (Cu Ka, λ=1.5418 Å). The obtained patterns were refined with General Structure Analysis System (GSAS) to determine structure parameters. Particle morphologies were observed by a Field-emission scanning electron microscopy (SEM) (JEOL, model JSM-6700). Chemical compositions were analyzed by inductively coupled plasma atomic emission spectrometry (ICP-AES) with a relative error of 2%. X-ray absorption spectrum (XAS) analysis for Mn and Ni K-edge was examined in transmission mode at beamline 1W2B with a Si (111) double-crystal monochromator of Beijing Synchrotron Radiation Facility (BSRF), and the energy was calibrated using Mn or Ni foil. Transmission electron microscope (TEM) observation was conducted on a Tecnai G2 F20 apparatus working at 200 kV. X-ray photoelectron spectroscopy (XPS) was examined on an ESCA-LAB MKII apparatus with a monochromatic Al Kα X-ray source. The adventitious carbon contamination with main C 1s peak ar 284.8 eV is used as a charge reference for XPS spectra. The thermogravimetric analysis (TG) and Differential Scanning Calorimeter (DSC) curves of as-prepared samples were examined a NETZSCH STA 449 F3


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Jupiter simultaneous thermal analyzer. 2.3. Electrochemical test All electrochemical performances of the samples were evaluated with CR2025-type coin cells using a Neware Battery Test System at room temperature, and the potential range was set to between 2.0 and 4.8 V. The cathode was mixtures of sample powders, polyvinylidene fluoride (PVDF) and carbon black in a weight ratio of 8:1:1 coated on an Aluminium foil. Excluding the mass of Al foil, each of the cathodes for testing was set at 3.5±0.5 mg, and the active material was around 2.8 mg. The commercial electrolyte comprised 1 M LiPF6 in a mixed solvent of ethylene carbonate (EC), ethyl methyl carbonate (EMC), and dimethyl carbonate (DMC) (1: 1: 1 in volume). Celgard 2500 polymer separator and lithium foil anode were used in the cell. The coin cells were assembled in an argon-filled glove box (both H2O and O2 less than 1 ppm). Electrochemical impedance spectroscopy (EIS) was recorded using an electrochemical workstation (CHI660C) in a frequency from 100 kHz to 0.01 Hz.

3. Results and Discussion 3.1 Structural and elemental characterizations The influence of hydrothermal treatment on the chemical composition is firstly examined by ICP. As listed Table S1, the treated samples H0-H3 exhibit different metal-ion molar ratio from sample P. Although the Ni/Mn ratio in four materials remains almost the same, the lithium molar ratio reduces distinctly after hydrothermal treatment. Moreover, lithium contents in the sample H1 and H2 are less than that of sample H0. The ICP results demonstrate that hydrothermal process extracted some of lithium from cathode particles, and the presence of glucose accelerated the loss of lithium. The loss of lithium could also alter the lattice and surface structure of sample.



Fig. 1 XRD patterns of sample P, H0, H1 and H2 with insets displaying the region between 20 and 25 degree. The mark “﹡” refers to the diffraction peaks of KCl as an internal standard for refinement. The structures of the samples are investigated by XRD. As shown in Fig. 1, four samples exhibit well distinguished diffraction peaks, indicating they are well-crystallized whether undergoing a hydrothermal treatment or not. The strong peaks in the patterns could be indexed to layered trigonal structure with space group of R3m, and the weak peaks in the 2θ range from 20 to 25 degree are attributed to monoclinic C/2m structure, due to the existence of super-lattice ordering of Li and Mn ions.30-32 Such integration pattern of layered and monoclinic structure is a special characteristic of LMRO cathode material. By comparing the peaks of the four samples, especially those between 20 and 25 degree, it is found that the XRD patterns are almost similar. The structural similarity for obtained samples is further confirmed by XRD refinements. Lattice parameters obtained by structural refinement using R3m model33 are displayed in the Table S2, and the refined patterns are shown in Figure S1. Lattice parameters of sample P and H0 are nearly the same, demonstrating that


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hydrothermal treatment in absence of glucose did not alter the bulk structure of Li-rich manganese-based oxide. Comparatively, in presence of some glucose, hydrothermal treatment resulted in a slight lattice expansion for sample H1 and H2, but c/a value kept the nearly same as sample P and H0. Obviously, the lattice expansion caused by hydrothermal treatment, i.e. the extraction of lithium, is different from that of ion doping. For the latter case, c/a usually becomes increase accompanying lattice expansion when introducing larger foreign ions.

10, 34

XRD

measurements also show that the treated samples only contain layered phase. Certainly, the detecting limit of XRD could not distinguish the trace impurity phase, especially the trace impurity phase located at particle surface or interface of material.

Fig. 2 Representative SEM images and particle-size distributions ranging from 0 ~ 0.7 um of sample (a-c) P, (d-f) H0, (g-i) H1 and (j-l) H2. SEM examination in Fig. 2 is aimed to reveal whether morphology or particle-size change through post-treatment processes. As seen in low-multiple SEM images, particles of all samples are well-dispersed and show irregular morphology. The



particle-size distributions corresponding to those images demonstrate that samples have the particle size in the range of 100 ~ 500 nm, and the size for most of particles is around 200 nm. Turning to magnified images, it can be found that the surfaces of the particles highlighted by yellow ellipse become rough after post-treatments, unlike the smooth surface of particles in sample P. This phenomenon occurred in all the three treated samples. Hydrothermal process roughens the surfaces of particles greatly, but does not alter their morphology and size distribution.

Fig. 3 (a, c) Normalized XANES spectra of Ni K-edge and Mn K-edge with the insets showing the 1s-4p peak energy of all samples; (b, d) k3-weighted Fourier transform magnitudes of Ni K-edge and Mn K-edge of the four samples. The results of XRD patterns and morphology examinations only show partial effects of the hydrothermal process on lattice structure and physical structure of particle surface, while specific electronic structural and the changes of surface elemental composition have not been well revealed. The normalized X-ray absorption near edge structure (XANES) and extended X-ray absorption fine structure (EXAFS) spectra at Ni K-edge and Mn K-edge are presented in Fig. 3 to illustrate the changes of local environment around transition metal ions and to determine the precise valences of bulk TM ions.11,

33

Almost overlapping Ni K-edge XANES spectra in Figure 3a


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demonstrate that the Ni ions in the hydrothermal treated samples have the same chemical environment as that in sample P. By comparing the spectra of our samples with those of Ni2+O and LiNi3+O2 and 1s-4p peak position in the inset of Fig. 3a, the valence state of Ni ions in all samples should be +2 dominantly.35 Different from the observation for Ni K-edge, the Mn K-edge XANES spectra show some changes for the hydrothermal treated samples. The two spectra of sample P and H0 completely overlap and their peak energies are located at 6560.8 eV, which directly means the manganese ions in the sample P and H0 have same valence state near to +4. The spectra of sample H1 and H2 shift to low energy, especially the absorption features at around 6550 and 6555 eV. The 1s-4p peak energies of H1 and H2 are also different from that of P. The lower peak energies of sample H1 and H2 indicate that the Mn ions in both of samples have lower valence state. According to previous works, the pre-edge part for Mn K-edge could be a vital criterion for the determination of Mn valence state.36-37 As showed in the inset of Fig. 3c (the bottom of right corner), no deviation is observed about the location of peaks among the pre-edge parts of the four samples. Therefore, the dominant valence state of Mn ions in the bulk of all the samples should be +4.

38

Nevertheless, especially in the sample H2, a very small

amount of Mn ions could be Mn3+. To reveal the valence states of TM ions on the surface of particles, the XPS spectra of Ni 2p and Mn 3s were shown in Fig. 4 and Table S3. Four Ni 2p spectra are obviously similar and the valence state of Ni ions are +2, which is evidenced by the satellite peaks, full width at half maximum (FWHM) and spin-orbital splitting value.

39-40

For the Mn 3s spectra, the splitting values of

doublet for sample H1 and H2 are 4.77 and 5.08 eV, respectively, larger than 4.5 eV observed for sample P and H0. Thus the valence state of surficial Mn ions for sample H1 and H2 is close to +3, and that for sample P and H0 are dominant +4. 14, 41-42 That is to say, adding glucose in the post-treatment results in a reduction of Mn ions on the particle surface, but Ni ions and bulk Mn ions remain unchanged.



Fig. 4 Ni 2p and Mn 3s of sample P, H0, H1 and H2 The changes of local structure around Mn and Ni ions can be known by comparing the k3-weighted Fourier transform magnitudes of Ni K-edge and Mn K-edge. As shown in Fig. 3 b and 3d, two dominant peaks below 4 Å of radical distance must be concerned. The peak located at lower radical distance corresponds to the scattering of nickel/manganese ions to their nearest-neighboring oxygen ions, and the one at higher distance correlates to the Ni/Mn-TM contributions. From Figure 3b and 3c, the same coordinated environment of Ni/Mn ions in the sample P and H0 can easily be confirmed by the overlapped curves. Moreover, there is little difference in the amplitude of Ni-O peaks, indicating that Ni ions in four samples are coordinated with six oxygen anions (showing a NiO6 coordination) even three samples undergoing a hydrothermal post-treatment. A distinct change is observed for the intensity of those Ni-TM peaks, which is also comparable with the amplitude changes of the Mn-TM peaks. The pronounced difference of peak intensity could attribute to the deteriorated cation-ordering of LiTM6 in the TM layers.

43

It should be noted that Mn-O peak

exhibit a reduced amplitude for sample H1 and H2, different from Ni-O peak. Such a reduction in amplitude is likely related with the oxygen loss and/or structural disorder around Mn ions.


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Fig. 5 Representative TEM images, high-resolution TEM images and average fast Fourier transformation (FFT) patterns of (a, b) sample P and (c-e) H1; (f) simulated FFT patterns corresponding to pattern (e) of both monoclinic phase along [101] zone axis and spinel phase along [-110] zone axis. The XAS results suggest the possibility of structural change in the post-treated samples of H1 and H2. Transmission electron microscope (TEM) is applied to examine local structure of sample P and H1 more explicitly. The overview images support the SEM analyses, i.e. particles in two samples of P and H1 have the similar morphology. The non-uniform contrast is observed in the high-resolution images of the particles for sample H1, further confirming that these particles have a roughened surface. The FFT patterns of sample P in Fig. 5b is assigned to the [00-1] zone axis of the monoclinic phase with crystal facets like (110), (040), (-130) and (-220) being marked, as previous report. 44 Comparatively, the pattern in Fig. 5e is composed of



two zone axes, corresponding to two different phases. As showed in Fig. 5f, one of zone axes is the [101] of C/2m phase, and the other is the [-110] of spinel phase. 21, 24, 45

Therefore, the hydrothermal treatment with adding glucose triggered a structural

transformation from layered into spinel phase. Along with the results of XAFS and XPS characterizations, the region of structural change is probably localized on surfaces of particles. Namely, the cathode particle surface has been reconstructed by the hydrothermal process, whereas the bulk structure of treated samples did not change. As showed in the TEM images, no conductive carbon produced by glucose carbonization was found on particle surface. Free of conductive carbon residue are also confirmed by the TG-DSC curves and C1s XPS spectra in Figure S2 and S3. Same as sample P, sample H0, H1 and H1 did not show any mass loss in the temperature range of 25 to 850 oC. For C1s spectra, the signals of C-C at 284.8 eV and C-O-C at 286.4 eV bonding belong to adventitious carbon contamination. The similar C-C and C-O-C signals for the four samples suggest that there is no conductive carbon resulted from the glucose carbonization on particle surface of H1 and H2.46 Nevertheless, the Li2CO3 component of for sample P at 289.5 eV is much more apparent than that for e other samples. It means that hydrothermal process can effectively eliminate Li2CO3 component on the surface of particles. 3.2 Electrochemical tests


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Fig. 6 (a) Normalized capacity vs. voltage curves at 1st cycle and 10th cycle, and (b-d) cycling performance data of sample P, H0, H1 and H2 respectively at 200, 400 and 600 mA g-1 with corresponding capacity retention ratios attached. The normalized discharge curves and electrochemical performances of the four samples at different current densities are compared in Fig. 6 to specify the relationship between their microstructure change and property. Four samples exhibit different discharge curves both at the 1st and 10th cycle (in Fig. 6a). An excess plateau at about 2.8 V is observed in the curves of the treated samples. This plateau is usually ascribed to the feature of rock-salt component or spinel phase. 47 It is noted that at 10th cycle, this featured plateau is absent for sample H0. The result of ICP has confirmed the extraction of lithium ions during the hydrothermal process, and adding glucose can lead to more lithium loss from the oxides. Thus, the appearance of the plateau at 2.8 V in the curves of both sample H1 and H2 seems reasonable. Although the initial capacities of the treated samples are not preferable, most of those cells exhibited an increased specific capacity after several cycles, especially the cells of sample H1 (Fig. 6b). The maximum capacities of sample H1 are ca. 211, 197 and 193 mAh/g at current densities of 200, 400 and 600 mA/g, respectively. Correspondingly, sample P exhibited 200, 170 and 172 mAh/g, lower than those of sample H1. With the progress of charge and discharge cycling, the difference of electrochemical behaviors for four samples becomes more and more obvious. As indicated by the



cycling profiles, the discharge capacity of sample P decreased rapidly. H0 and H2 also showed a reduced capacity, but their dropping is not as much as that of sample P. Sample H1 gave a stable discharge capacity at current densities of 200, 400 and 600 mA/g. After 100 cycles under those current densities, the capacities of sample H1 remained 99%, 97% and 98% of its maximum values, which pronouncedly exceeds the performance of the others. Thus, the cells of sample H1 possessed the highest specific capacities of 212, 191 and 191 mAh/g at 100th cycle. The preferable electrochemical performance of sample H0 than that of sample P is also remarkable though both samples have identical structure and the valence states and surround of TM ions. H0 sample underwent a hydrothermal treatment. The elimination of unfavorable surficial residue and the roughening of particle surface could be beneficial to the intercalation/de-intercalation of lithium ions in the discharge/charge electrochemical process. However, the hydrothermal process with pure water is unhelpful for cycling stability. Adding glucose in the hydrothermal treatment process could improve the cycling performance of lithium-rich cathode by the formation of spinel component, but over amount of glucose could not enhanced the cycling performance probably due to the extraction of more lithium ions and excessive surface roughening. The effects of hydrothermal treatment on the electrochemical behaviors are also examined by EIS in frequency region from 100 kHz to 10 mHz. The obtained results are compared in Figure S4 and Table S4. Due to the same equivalent circuit model for all the half-cells, the values of Re, Rsf and Rct could illustrate the surface and bulk properties of the four samples.48 Evidently, the half-cell with sample H1 as cathode has the lowest Rsf and Rct values at any test cycle, which means the superior electrode kinetics and stable property of this cathode.


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Fig. 7 Specific capacity (blue) and energy density (red) vs. cycle number data of sample (a) P and (b) H1 at 100 mA/g with retention ratios marked at 100th and 200th cycle. The initial charge-discharge curves of sample P and H1 at low current densities are presented in Figure S5 of the Supporting Information. The lower initial charge capacity of sample H1 is closely related with the shortening of plateau above 4.5 V due to the presence of deficiency of lithium and oxygen ions. On the other hand, the coulombic efficiencies of sample H1is higher than that of sample P. Furthermore, Fig. 7 exhibits the prolonged electrochemical performance at a current density of 100 mA/g. Like the performances under higher current densities, several cycles were required for the two samples to achieve their maximum capacities. The maximum capacity of sample P is 209 mAh/g, and that of sample H1 is 237 mAh/g. As showed in Figure S6, the average maximum capacity of sample H1 is 236 ± 3 mAh/g, and the standard deviation of capacity value after 5 cycles is small. Meanwhile, the energy densities of sample P and H1 at the first cycle are 717 and 705 Wh/Kg, and their maximum energy densities are 751 and 827 Wh/Kg, respectively. Therefore, the performance of sample H1 at 100 mA/g is also better than that of sample P. After 100 cycles, the capacity retention ratio relative to the maximum specific capacity of sample H1 is 98.5%, higher than 93.0% for sample P. At the 200th cycle, the capacity retention ratios of sample P and H1 are 83.8% and 82.3%, respectively. Namely, the



discharge specific capacity of sample H1 at the 200th cycle remains 195 mAh/g, and the capacity of sample P is 175 mAh/g only. Moreover, as presented in the cycling profiles, the retention ratio of energy density for sample H1 is superior to that for sample P at any cycle. As shown in Figure S7, the average voltage decreased with the cycles. At the 200th cycle, sample H1 had a higher average voltage comparing to sample P even though sample H1 showed a low initial voltage. Therefore, the hydrothermal treatment in presence of appropriate amount of glucose could be helpful to enhance the specific capacity, energy density and cycling performance. 3.3 The mechanism of the post-treatment with glucose

Fig. 8 Schematic diagram of the mechanism of the hydrothermal post-treatment with adding glucose The above characterizations are sufficient to reveal the effect of the hydrothermal post-treatment with adding glucose. The possible mechanism of this post-treatment process is briefly illustrated in Fig. 8. When aqueous mixture containing LMRO particles and glucose solution is heating up, LMRO particles undergo two main steps to reconstruct their surface. Firstly, oxide particles would be roughened and lithium ions in the near-surface layer could exchange with protons in the solution, similar to the influence of acid treatment.21-22 Meanwhile, an important chemical reaction, i.e. the carbonization of glucose to produce carbon species, is triggered under hydrothermal conditions accompanying other reactions involving glucose, as reported in previous works. 49-50 In the second step, carbon species around particles would react with oxide and reduce Mn ions to produce spinel phase. It should be emphasized that the Mn ions in the Mn-rich region located at the surface or near-surface of particles are easily reduced. The oxygen loss and/or the reduction of Mn4+ ions also results in a distortion of Mn-O coordinated polyhedrons. Therefore, the reconstruction of LMRO


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surface structure during the hydrothermal post-treatment mainly includes an extraction of lithium ions to rough the particle surface, a hydrothermal carbonization of glucose to produce carbon species that react with oxide to reduce Mn ions and promote the formation of spinel phase on the surface or near-surface. Notably, the charge compensation for losing lithium ions and the reduction of Mn ions in the post-treatment can be balanced by the intercalation of protons and the appearance of oxygen vacancies. 3.4 Electrochemical performance of other hydrothermal post-treated LMROs

Fig. 9 Discharge specific capacity vs. cycle number profiles of sample LMR333 and LMR333-H at current densities of (a) 100, (b) 200, (c) 400 and (d) 600 mA/g. Hydrothermal post-treatment method in present work can be applied for other LMROs

to

enhance

their

electrochemical

performance.

Stoichiometric

Li1.2Mn0.54Ni0.13Co0.13O2 material (the detail preparation is given in the experimental section of Supporting Information) was also hydrothermally treated via the process described above with addition of 0.025 g glucose. The pristine and treated oxides are named as LMR333 and LMR333-H, respectively. The electrochemical performances of the two samples were separately examined at current densities of 100, 200, 400 and



600 mA/g in the voltage region of 2.0 ~ 4.8 V. Similar to the Co-free cathodes, the cells reached the maximum discharge specific capacities after several cycles’ activation, as shown in Fig.9. The maximum capacities of sample LMR333-H are 244, 237, 224 and 222 mAh/g at the current densities of 100, 200, 400 and 600 mA/g, respectively, while the maximum values of sample LMR333 are 243, 223, 201 and 203mAh/g. Moreover, the capacity retentions of sample LMR333-H at 100th cycle are higher than that of LMR333. In more detail, sample LMR333-H remains 213, 198, 185 and 181 mAh/g with retention ratios of 86.8 %, 83.2 %, 83.0 % and 81.3%, respectively. The electrochemical performance results observed for LMR333-H suggests that hydrothermal post-treatment strategy is a widely applicable method to improve the surface microstructure of lithium-rich manganese-based oxides, and enhance their discharge specific capacity and cycling stability.

4. Conclusions In this work, the specific capacity and cycling performance of two LMROs were distinctly improved by reconstructing the surface microstructure via two steps of the hydrothermal glucose carbonization and subsequent reduction procedure. Especially, the Co-free sample H1 treated with 0.025 g glucose exhibits a discharge capacity of 237 mAh/g at 100 mA/g, and the capacity retention ratio at this tested current density are higher than 93 % after 100th cycle. The particle surface of samples after reconstructing had been found to be roughened, and the presence of spinel phase contributes to the enhanced electrochemical performance by improving the charge transfer across the interface. The formation process of spinel phase attributes to the reduction of surficial Mn ions and the distortion of oxygen anion or the presence of oxygen deficiency. Therefore, these findings may provide a vital hint for the strategy to reconstruct surface layer of Li-rich cathodes, and breakthroughs in today’s state-of-art Li-ion batteries.

Acknowledgement This work was financially supported by NSFC (Grants 21571176, 21611530688,


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21771171, 21671077 and 21025104) and the Scientific Start Foundation of Longyan University (LB2018014).

Supporting Information: The Supporting Information is available free of charge on the ACS Publications website at DOI: Additional materials and methods, ICP results, Calculated XRD patterns, XPS spectra, electrochemical data and EIS spectra

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