Insights into the chemical and structural evolution of Li-rich layered oxide cathode materials

Abstract

Lithium-rich layered oxide cathodes have an advantage of high energy density. However, continuous capacity fading and voltage hysteresis have largely restricted their practical application. In this work, we investigated in depth the voltage-fading mechanism of Li-rich materials by in situ Raman spectroscopy, X-ray absorption near edge structure (XANES) spectroscopy, and high-resolution transmission electron microscopy (HRTEM). It has been found that the chemical and structural evolution of a Li_1.2Ni_0.15Co_0.1Mn_0.55O₂ cathode is quasi-reversible in a whole charge–discharge cycle. The structural evolution is evidently irreversible upon long-term cycling, resulting in the dissolution of cations from the lattices and structural collapse, which in turn leads to undesirable voltage fading. This finding is important for a better understanding of the redox reaction mechanisms of high-capacity Li-rich cathodes.

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1. Introduction

Li-Rich layered oxides, represented by the chemical formula xLiMO₂·(1 − x)Li₂MnO₃ (0 < x < 1.0, M = Mn, Ni, Co, etc.), have recently attracted significant research interest due to their extraordinarily high energy density, being considered as cathode materials for lithium ion batteries.^1,2 The ultra-high capacity of Li-rich compounds originates from the simultaneous cationic and anionic redox processes.^3,4

However, Tarascon and his co-workers⁵ have shown that the oxygen redox process is strongly associated with the structural evolution in Li-rich materials. In addition, redox-mediated oxygen vacancies can be induced by the formation and escape of oxygen gas at a high charge cutoff voltage, typically above 4.5 V for most Li-rich or Na-deficient materials.^5,6 Therefore, the oxygenic redox-mediated irreversible structural evolution is theoretically inevitable if we want to activate and take advantage of the anionic redox activity.^7,8 Therefore, the anionic redox activity endows Li-rich layered oxides with a satisfactory high capacity but unsatisfactory long-term cycling stability.^9,10 Ultimately, these restrictions still hinder the practical application of Li-rich layered oxides.

Nonetheless, various characterization methods have been previously developed to study the electrochemical degradation mechanisms of Li-rich materials.^11–13 However, the previous studies lack a comprehensive consideration of the evolution of cations and the local structure of Li-rich cathodes after short- and long-term charge/discharge cycles, which have been a long-standing hurdle for the understanding of the mechanistic link between the chemical/structural evolution and the voltage and capacity degradation of Li-rich materials. Therefore, it is of great importance to clarify the impact of cation evolution and the local structure on the electrochemical stability of Li-rich layered cathode materials.

Herein, to better understand the mechanistic link between the chemical/structural evolution and the voltage degradation of Li-rich materials, we employed different characterization techniques to investigate the structural characteristics of a representative layer-structured Li-rich cathode material Li_1.2Ni_0.15Co_0.1Mn_0.55O₂ before and after cycling. The results clearly revealed that the cation dissolution and the irreversible local structural evolution are mainly responsible for the degradation mechanism of Li-rich materials.

2. Experimental section

2.1. Chemicals

Commercially available chemical reagents were used as received without any further purification. Nickel(II) sulfate hexahydrate (NiSO₄·6H₂O), cobalt(II) sulfate heptahydrate (CoSO₄·7H₂O), manganese(II) sulfate monohydrate (MnSO₄·H₂O), sodium carbonate (Na₂CO₃), ammonium hydroxide (NH₃·H₂O), lithium carbonate (Li₂CO₃), poly(vinylidene fluoride) (PVDF), Super P, and 1-methyl-2-pyrrolidinone (NMP) were provided by Sinopharm Chemical Reagent Co. Ltd (Shanghai, China). Commercially available battery-grade chemicals, diethyl carbonate (DEC), dimethyl carbonate (DMC), ethylene carbonate (EC), and lithium hexafluorophosphate (LiPF₆) were obtained from Sigma-Aldrich Co. Ltd (Shanghai, China).

2.2. Preparation of the Li_1.2Ni_0.15Co_0.1Mn_0.55O₂ material

Firstly, a Ni_0.15Co_0.1Mn_0.55(CO₃)_0.8 precursor material was synthesized via co-precipitation by the modified Jean-Marie Tarascon's method.¹⁴ In detail, the desired stoichiometric amounts of MnSO₄·H₂O, CoSO₄·7H₂O and NiSO₄·6H₂O were dissolved in an aqueous solution (cation concentration of 2.0 M in total, Ni²⁺ : Co²⁺ : Mn²⁺ = 1.5 : 0.5 : 5.5 in molar ratio). An appropriate amount of Na₂CO₃ aqueous solution (2.0 M, CO₃²⁻ : M²⁺ = 1.1 in molar ratio) was used as the precipitant and added dropwise into the above-mentioned aqueous sulfate solution under continuous stirring. Ammonia solution (20 wt%) was fed into the reaction as a chelating agent. Then the mixed solution was kept at 80 °C and a pH of 8 for 12 h. The greyish yellow Ni_0.15Co_0.1Mn_0.55(CO₃)_0.8 precipitate obtained was centrifuged, collected, thoroughly washed with deionized water, and vacuum-dried at 120 °C for 12 h. Then the as-prepared Ni_0.15Co_0.1Mn_0.55(CO₃)_0.8 precursor was grind-mixed with Li₂CO₃ (Li : M = 1.6 in molar ratio). Then the obtained mixed powders were pre-calcined at 500 °C for 6 h, followed by sintering at 900 °C for 15 h under air atmosphere to afford the grayish black Li_1.2Ni_0.15Co_0.1Mn_0.55O₂ material.

2.3. Electrochemical measurements

Firstly, the cathode electrode material slurry was fabricated with 80 wt% of Li_1.2Ni_0.15Co_0.1Mn_0.55O₂ active material, 10 wt% of PVDF binder, and 10 wt% of conductive Super P additive in NMP. Then the slurry was pasted onto aluminium foil. The cathode electrode laminates were pressed down for an active material loading of 4–5 mg cm⁻² with a scraper blade and a calender. Once calendered, the electrodes were thoroughly dried for 12 h at 120 °C under vacuum. The obtained electrode material was cut into disks with a diameter of 12 mm. Then CR2016 coin-type cells were assembled in an argon-filled glove box with the as-prepared Li_1.2Ni_0.15Co_0.1Mn_0.55O₂ electrode as the cathode, metallic lithium foil as the anode, a polypropylene micro-porous film (Cellgard 2400) as the separator, and 1 M LiPF₆ (EC/DMC/DEC, 1 : 1 : 1 volume ratio) as the electrolyte. The galvanostatic discharge–charge measurements were performed with the as-assembled CR2016 coin-type cells on a battery test system with a potential range of 2.25 to 4.75 V (vs. Li/Li⁺) at 25 °C.

2.4. Characterization

XRD (powder X-ray diffraction) characterization of the Li_1.2Ni_0.15Co_0.1Mn_0.55O₂ material was performed on a Bruker D8 XRD instrument (Bruker Corporation, German) for Rietveld refinement and a Rigaku Smartlab 9KW XRD instrument for general measurements on the electrodes. The morphologies of the Li_1.2Ni_0.15Co_0.1Mn_0.55O₂ sample were observed on a FEI Tecnai F30 high resolution transmission electron microscope (HRTEM, FEI Corporation, America) and a TESCAN Mira 3 scanning electron microscope (SEM, TESCAN Corporation, Czech Republic). In situ Raman spectroscopic characterization was carried out on a Renishaw inVia microscopic Raman spectrometer (Renishaw Corporation, UK) with a 532 nm wavelength laser, a laser beam of ≈1 μm diameter, a 50 mW power, a CCD camera, and an optical Leica microscope. A 50× objective lens and a 1800 lines per mm grating were used to focus the incident beam. The N₂ adsorption–desorption isotherms were measured using a Micromeritics ASAP 2020 analyzer at 77 K (liquid nitrogen temperature). The specific surface area of the cathode sample was determined using the Brunauer–Emmett–Teller (BET) model, and pore size distribution was calculated using the Barrett–Joyner–Halenda (BJH) method. Electrochemical impedance spectroscopy (EIS) measurements on the Li_1.2Ni_0.15Co_0.1Mn_0.55O₂ cathode were carried out using a Bio-Logic VMP3 electrochemical workstation (Bio-Logic Science Instruments Inc., France) with a 5 mV perturbation and a frequency range of 10⁵–10⁻² Hz. XANES and nano-CT experiments were performed at the Shanghai Synchrotron Radiation Facility (SSRF) BL08U1A and the Beijing Synchrotron Radiation Facility (BSRF) 4W1A, respectively.

3. Results and discussion

3.1. Physicochemical properties of as-prepared Li_1.2Ni_0.15Co_0.1Mn_0.55O₂ cathode material

The as-prepared pristine Li_1.2Ni_0.15Co_0.1Mn_0.55O₂ material was characterized using scanning electron microscopy (SEM), transmission electron microscopy (TEM), powder X-ray diffusion (XRD) and N₂ adsorption–desorption isotherms. The SEM image (Fig. 1A) and TEM image (Fig. 1B) show that the Li_1.2Ni_0.15Co_0.1Mn_0.55O₂ sample is composed of secondary particles of ∼2 μm. The Rietveld refinement of the XRD pattern of the pristine Li_1.2Ni_0.15Co_0.1Mn_0.55O₂ is presented in Fig. 1C. The XRD pattern displays relatively sharp peaks, implying high crystallinity. The obtained XRD reflections are in line with a rhombohedral structure (R3̄m, LiMO₂) and a monoclinic structure (C2/m, Li₂MnO₃), indicating that the fine layered structure of the Li_1.2Ni_0.15Co_0.1Mn_0.55O₂ material is crystallized with rhombohedral symmetry with an R3̄m space group and monoclinic symmetry with a C2/m space group. The N₂ adsorption–desorption isotherm and the BJH pore-size distribution are shown in Fig. 1D. The N₂ adsorption–desorption isotherm shows a hysteresis loop due to the capillary condensation of N₂ gas in the pores of the Li_1.2Ni_0.15Co_0.1Mn_0.55O₂ sample.¹⁵ The BJH average pore diameter is 5.76 nm, and the BET specific surface area is 20.61 m² g⁻¹, suggesting that the Li_1.2Ni_0.15Co_0.1Mn_0.55O₂ sample exhibits typical characteristics of mesoporous materials.

Fig. 1 Characterization and electrochemical performance of Li_1.2Ni_0.15Co_0.1Mn_0.55O₂: (A) SEM image, (B) TEM image, (C) XRD and Rietveld refinements, and (D) N₂ adsorption–desorption isotherm (inset: BJH pore-size distribution); (E) charge–discharge profiles at different cycle numbers; (F) specific capacity, coulombic efficiency, and average discharge voltage as a function of cycle number.

The electrochemical properties of the Li_1.2Ni_0.15Co_0.1Mn_0.55O₂ cathode material were investigated in a half cell configuration. Galvanostatic discharge–charge cycling was performed under a current density of 20 mA g⁻¹ between 2.25 V and 4.75 V. The results are shown in Fig. 1E. As shown in Fig. 1E, two plateaus with charge voltages around 4.0 V and 4.5 V can be observed during the first charging process, which are typical characteristics of Li-rich layered materials.^16,17 In addition, the plateau (at 4.5 V) disappears in the following charging processes, revealing a structural variation in the first charging process, which may be ascribed to the oxygen redox process.¹⁸ The Li_1.2Ni_0.15Co_0.1Mn_0.55O₂ cathode material exhibits good cycling performance as it exhibits high coulombic efficiencies (close to 100%) in the following cycles. However, as shown in Fig. 1F, the Li_1.2Ni_0.15Co_0.1Mn_0.55O₂ cathode material suffers from capacity fading and voltage decay during cycling, which are common problems encountered in Li-rich layered cathode materials.^19,20

3.2. Characterization and analysis of the Li_1.2Ni_0.15Co_0.1Mn_0.55O₂ cathode material during its first charge–discharge cycle

3.2.1. In situ Raman. In situ Raman spectra of the Li_1.2Ni_0.15Co_0.1Mn_0.55O₂ electrode at different voltage states during its first charging and discharging cycle can be seen in Fig. 2. The Raman peak at 518 cm⁻¹ belongs to the A1g bond (δ_O–M–O), shifts to 480 cm⁻¹ (A1g′) during de-lithiation and re-lithiation processes, and eventually returns to 518 cm⁻¹. This indicates the reversible variation of the lattice parameters during the Li⁺ extraction and re-insertion processes, which is, however, more drastic during the anionic redox process.²¹ Meanwhile, the broad Raman peak at 618 cm⁻¹ can be attributed to E_g vibrations (υ_MO6). The intensity of this peak is not obviously decreased until Li_1.2Ni_0.15Co_0.1Mn_0.55O₂ reaches the second charge plateau, reflecting a layered-to-spinel phase transition during the anionic redox process.^22,23 Then this Raman peak regains its intensity during the discharge process of the Li_1.2Ni_0.15Co_0.1Mn_0.55O₂ cathode, a fact that can be related to the reversible structural variation during the whole charge–discharge process. Therefore, the cationic redox process can activate the variation of the lattice parameters without any variation of the crystal phases, whereas the anionic redox process can induce variations both on the lattice parameters and crystal phase. The Raman peaks at 700, 716, 742, 891, 905, 936, and 974 cm⁻¹ can be ascribed to LiPF₆, EC, solvated EC with Li⁺ (EC-Li⁺), DEC, solvated DEC with Li⁺ (DEC-Li⁺), DMC, and solvated DMC with Li⁺ (DMC-Li⁺), respectively.²⁴ We can conclude that the electrolyte undergoes an oxidation decomposition during the charging process and a reversible reduction resurrection during the discharging process. Notably, an intense Raman peak at ∼850 cm⁻¹ is observed for a charging voltage above 4.4 V, increasing gradually until the end of charging, and then decreasing until the discharging voltage falls below 3.95 V. These peaks can be attributed to the typical peroxo O–O (O₂²⁻) bond stretching, caused by the oxygen redox process in Li_1.2Ni_0.15Co_0.1Mn_0.55O₂.^25,26 Interestingly, the Raman peak corresponding to peroxo O–O (O₂²⁻) slightly shifts to higher wavenumbers during charging, and then gradually returns to its original position during discharging. This indicates that the bond length is shortened when the O²⁻ in Li_1.2Ni_0.15Co_0.1Mn_0.55O₂ is partially oxidized to O₂²⁻, whereas the O²⁻/O⁻ redox process is reversible during charging/discharging.^27,28 The in situ Raman spectra provide definite evidence that reversible oxygen evolution and structural variation are involved in the whole charging/discharging process of Li_1.2Ni_0.15Co_0.1Mn_0.55O₂.

Fig. 2 In situ Raman spectra of Li_1.2Ni_0.15Co_0.1Mn_0.55O₂ throughout the first galvanostatic cycle. The corresponding voltage profiles are shown for clarity.

3.2.2. Ex situ XANES. Ex situ X-ray absorption near-edge structure (XANES) spectra of the Li_1.2Ni_0.15Co_0.1Mn_0.55O₂ electrodes were obtained at various designated voltages, providing valuable information about the valence state and the redox chemistry of Li_1.2Ni_0.15Co_0.1Mn_0.55O₂ upon its first charge–discharge process. As shown in Fig. 3, the shape of the spectrum of the O K edge peak in the pristine state is distinctly different from those of other states, whereas the spectra of other elements show a similar shape throughout all states of the charge/discharge process. This is probably due to the presence of oxygen vacancies in the pristine material, which is in accordance with previous reports.^29,30 Upon charging from the pristine open-circuit voltage (OCV) → C 4.35 V, the spectra of Ni L₃, Mn L₃, and Co L₃ edge peaks shift rightward to higher energy levels, without any significant shift of the O K edge peak, suggesting the oxidation of the transition metal (TM) ions during this charging process. In contrast, when charging through C 4.35 V → C 4.75 V, a significant change of the shape (broaden) and position (rightward shift) of the O K peak can be observed, whereas the position of TM L₃-edge peaks changes negligibly, indicating that the oxygen redox instead of cationic redox process is mainly involved in this charging plateau.³¹ It is also worth noting that the increase in the pre-edge intensities may be assigned to the influence of the anionic redox reaction and the stronger hybridization of the higher-valent transition metal and oxygen.³² When discharging through C 4.75 V → D 3.75 V, all the TM L₃ and O K edge peaks shift to lower energy levels, which is related to the simultaneous reduction of oxygen and TM ions, reflecting the appearance of a hybridized “TM-O redox” process.³³ Upon further discharging through D 3.75 V → D 2.25 V, a continuous leftward shift of the XAS peaks can be observed. The TM L₃ edge peaks reach their initial positions with a decrease in their intensity, indicating the weakening of the hybridization strength between the transition metal and oxygen during this process.²⁰ The evolution of the XAS spectra indicates that both the transition metal cations and oxygen ions participate in the charge compensation reactions during the reversible redox cycles. The redox reactions of the transition metal cations proceed at a lower voltage, whereas in the high-voltage region between 4.35 V and 4.75 V, electron transfer occurs mainly at the oxygen sites. Then, the simultaneous reduction of oxygen and transition metal ions can be achieved by the hybridized “TM-O redox” process. As for one whole charge/discharge process, the chemical evolution of the Li_1.2Ni_0.15Co_0.1Mn_0.55O₂ cathode is quasi-reversible.

Fig. 3 Ex situ XANES of Li_1.2Ni_0.15Co_0.1Mn_0.55O₂ during the first cycle: Co L₃ (A), Mn L₃ (B), Ni L₃ (C), and O K (D) edges.

3.3. Characterization and analysis of the Li_1.2Ni_0.15Co_0.1Mn_0.55O₂ cathode material during long-term cycle-aging

3.3.1. HR-TEM and SAED. The HR-TEM images of the Li_1.2Ni_0.15Co_0.1Mn_0.55O₂ electrodes before and after cycle-aging revealed that the cycling deterioration of Li_1.2Ni_0.15Co_0.1Mn_0.55O₂ is mainly associated with its phase transformation during electrochemical cycling. Detailed lattice fringes of the pristine Li_1.2Ni_0.15Co_0.1Mn_0.55O₂ material can be observed in Fig. 4A, revealing the typical (111) crystal planes of the Li₂MnO₃ phase and the typical (003) crystal plane of the LiMO₂ phase, respectively.^34,35 The high crystallinity of the pristine Li_1.2Ni_0.15Co_0.1Mn_0.55O₂ material is further verified using selected-area electron diffraction (SAED) (Fig. 4B). However, the HR-TEM image of the cycled material shows regions with highly degraded and randomly distributed lattice fringes, verifying the phase transformation from the original highly crystallized layered structure into a spinel-like structure. The phase transformation might have resulted from the Li and oxygen vacancies formed during the charging processes. These vacancies facilitate the migration of the transition metal ions from their original sites into the lithium layer, resulting in the structural rearrangement and irreversible local structural changes.³⁶ The bulk degradation and electrochemical performance deterioration of the Li-rich cathode are probably associated with this phase transformation.³⁷

Fig. 4 HRTEM images of Li_1.2Ni_0.15Co_0.1Mn_0.55O₂ before (A) and after 90 cycles (C); and (B) and (D) the corresponding SAED patterns.

3.3.2. TEM-EDS. Normalized using the C peak intensity, Fig. 5 shows the spectra of the TEM-energy dispersive spectroscopy (EDS) of the Li_1.2Ni_0.15Co_0.1Mn_0.55O₂ cathode and the corresponding elemental analysis (in molar ratio) before and after 90 times charge–discharge. Before cycling, the pristine Li_1.2Ni_0.15Co_0.1Mn_0.55O₂ material possesses almost the same molar ratio of O : Mn : Ni : Co = 71.8 : 19.3 : 5.37 : 3.53 as the nominal one. As shown in Fig. 5B, after cycling, the retention rate of O, Mn, Ni, and Co is 90.4%, 89.5%, 67.3%, and 95.1%, respectively. Therefore, we can draw the conclusion that the electrochemical cycling is responsible for the observed change in the chemical composition in Li_1.2Ni_0.15Co_0.1Mn_0.55O₂. This finding is consistent with the phase transformation from the layered structure (Li[Li_xM_1−x]O₂, where x ≥ 0) to the defective spinel structure (Li_xM₂O₄) as extended cycling occurs,³⁸ which is also further corroborated by the HR-TEM observation. In fact, the oxygen loss and dissolution of transition metals, especially for Ni, cause a deterioration of the structural stability, which contributes to cell capacity fading.³⁹

Fig. 5 TEM-EDS profiles (A) and chemical composition (B) of the Li_1.2Ni_0.15Co_0.1Mn_0.55O₂ cathode before and after cycling; (C): typical TEM image of a cycled Li_1.2Ni_0.15Co_0.1Mn_0.55O₂ cathode for two-spot EDS analysis; and (D) EDS profile of a cycled Li_1.2Ni_0.15Co_0.1Mn_0.55O₂ cathode.

EDS analysis was also conducted at two different locations of the aged Li_1.2Ni_0.15Co_0.1Mn_0.55O₂ cathode with a spot size of 10 nm (Fig. 5C). The two selected areas display distinctly different crystal structures, as the A zone possesses better crystallinity than the B zone. Normalized using the C peak intensity, the obtained EDS spectra are shown in Fig. 5D. Apparently, the content of each element in the A zone is higher than that in the B zone, signifying the irreversible structural and chemical evolution during cycling. The loss of elements is differentiated within the entire area, resulting in distortion of crystal periodicity. It can be conclusively stated that the lattice transformation and chemical evolution are the main degradation mechanisms for the Li-rich cathodes. Interestingly, the variation of the molar ratio of Co content is relatively negligible at different zones or after cycling, suggesting the good stability of the Co element in the lattice. Therefore, appropriate doping with the Co element may be an effective method to improve the structural stability of Li-rich materials.

3.3.3. EIS studies. The EIS characterization of a cycle-aged cell was performed at a state of charge of 100% at 25 °C. Fig. 6A displays the specific equivalent circuit model,⁴⁰ and the obtained Nyquist plots and detailed resulting electrochemical impedance values are given in Fig. 6B and Table 1, respectively. The inset of Fig. 6B displays the magnified regions at high frequencies in the Nyquist plots. The intersection with the x-axis, which can be attributed to the ohmic resistance (R_s) of the battery, increases monotonically with the increasing cycling number, reflecting the continuously increasing internal resistance of the battery during cycle-aging. Two semicircles can be observed in a frequency range of 25 Hz to 200 kHz. The left one can be assigned to the cumulative solid–electrolyte interfaces (SEI) of both the electrodes, whereas the right semicircle can be ascribed to the solid diffusion in the electrode materials. A widening of the semicircles can be observed in Fig. 6B, indicating the general trend of impedance increase upon increasing the number of cycles, which is related to the aging of the cell.⁴¹

Fig. 6 EIS analysis of pristine and cycling-aged Li_1.2Ni_0.15Co_0.1Mn_0.55O₂: a specific equivalent circuit model (A), Nyquist curves (B) and evolution of R_e, R_SEI, and R_S (C) as a function of cycle numbers.

Table 1 Simulated electrochemical impedance parameters of Li_1.2Ni_0.15Co_0.1Mn_0.55O₂

Number of cycles	R s (Ω)	R SEI (Ω)	R e (Ω)	R ct (Ω)	Fitting coefficient
1	48.07	74.96	229.3	2760	97.79
2	49.97	77.25	234.4	2661	97.14
5	51.29	78.24	234.7	2058	96.63
10	51.66	78.8	235.2	2018	96.89
20	52.46	79.35	235.2	1981	96.70
30	52.9	79.69	235.3	1586	96.98
50	53.54	80.19	235.4	1495	96.45
70	53.94	80.72	236.8	1356	97.24
90	54.59	81.29	237	1251	96.06

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Fig. 6C shows the absolute values of impedances at characteristic frequencies upon increasing the number of cycles. The green curve shows the resistance of solid diffusion in the cathode electrode material (R_e), whereas the inset displays the ohmic resistance (R_s) of the cell and the resistance of the SEI (R_SEI). A distinct increase during the initial several cycles and a slow increase in the following cycles can be observed in each impedance curve due to the progressive aging of the cell. It can be conclusively stated that the increase of the R_s, R_SEI, and R_e values reflects the increase of the internal ohmic resistance, thickness of the SEI, and resistance of solid diffusion in the electrode material during cyclic aging, respectively. It should be clarified that the rise of R_e is not solely caused by the structural variations of the cathode material; the identified isolated area in the surfaces of the electrodes (SEI/electrode interfaces) is also crucial.⁴² The formation and growth of the SEI contribute significantly to the increase of the impedance parameters of the aging cell, and a robust SEI on the electrodes can be established after 10 charge–discharge cycles. However, as suggested in Fig. 6B, the degree of R_e increase is larger than that of R_SEI, revealing the existence of cathode degradation during cycling.⁴³

3.3.4. XANES. The X-ray absorption near-edge structure (XANES) spectra at the O K edge and L₃-edges of Mn, Co and Ni were collected at the pristine and 90 charge–discharge cycling-aged Li_1.2Ni_0.15Co_0.1Mn_0.55O₂ electrode (Fig. 7). An obvious leftward shift toward lower energy levels accompanied by a decreased intensity of the pre-edge peaks can be observed after long-term cycling. This suggests that the average content and valence state of Mn, Co, and Ni cations in Li_1.2Ni_0.15Co_0.1Mn_0.55O₂ are continuously reduced upon increasing the number of cycles.⁴⁴ As previously reported, the dissolution of cations in the electrolyte is more severe for the cations of lower valence states.^45,46 This transition metal cation dissolution is a common detrimental factor towards voltage decrease and capacity decay.^47,48 Moreover, the peak at 534 eV, shown in the O K-edge XAS spectra in Fig. 7D, can be associated with Li₂CO₃, which is a major constituent of the SEI layer, indicating the growth of the CEI layer during cycling.^49,50 The decreased pre-edge peaks and similar-shaped O K-edge spectrum before and after cycling can be explained by taking into account the surface reconstruction due to the formation of oxygen vacancies and rock-salt/spinel phases.²⁰ The evolution of the XANES spectra in the cycling-aged Li_1.2Ni_0.15Co_0.1Mn_0.55O₂ electrode reveals its structural degradation during cycling.

Fig. 7 XANES of pristine and cycling-aged Li_1.2Ni_0.15Co_0.1Mn_0.55O₂: Co L₃ (A), Mn L3 (B), Ni L₃ (C), and O K (D) edges.

3.3.5. XRD. Fig. 8 displays the XRD patterns of the pristine and cycling-aged Li_1.2Ni_0.15Co_0.1Mn_0.55O₂ cathode materials. As shown in Fig. 8A, the diffraction peaks of pristine Li_1.2Ni_0.15Co_0.1Mn_0.55O₂ cathode electrodes at 36.62°, 36.69°, 36.76°, 36.85°, and 36.94° can be attributed to the (1 0 1) plane of C/2m, (2 0 −1) plane of R3̄m, (1 3 −1) plane of R3̄m, (0 0 6) plane of C/2m, and (1 1 −2) plane of R3̄m, respectively. During the initial several cycles, the obtained XRD patterns of cycling-aged Li_1.2Ni_0.15Co_0.1Mn_0.55O₂ electrodes are similar to those observed over fresh cells. However, the diffraction peaks of the cycled samples that slightly shifted to a lower diffraction angle greatly gained in intensity, as previously reported by Debasish Mohanty et al.⁵¹ This evolution of the XRD patterns can be ascribed to the increased lattice parameters along with increased atom disorder in the local lattice of the cycling-aged Li-rich material.⁵² Interestingly, a new main peak emerges at 36.50°, which is characteristic of the (3 1 1) lattice plane of the Li₄Mn₅O₁₂ spinel phase (JCPDS 46-0810). The latter suggests that the high cut-off voltage cycling induces an inevitable complex phase transition of the cycling-aged Li_1.2Ni_0.15Co_0.1Mn_0.55O₂ electrodes.⁵³ The atom loss in the lattice and layered-to-spinel phase conversion resulted in an increase in the lattice parameters. The overall structure of the cycling-aged Li_1.2Ni_0.15Co_0.1Mn_0.55O₂ cathode material still remains layered, but spinel domains of small size can be gradually generated and then locally integrated into the layered framework.

Fig. 8 The evolution of XRD patterns in a range of 36–37.5 as a function of cycling numbers for Li_1.2Ni_0.15Co_0.1Mn_0.55O₂: (A) galvanostatically cycled in voltage ranges of 2.25–4.75 V and (B) 2.25–4.35 V.

For better comparison, the Li_1.2Ni_0.15Co_0.1Mn_0.55O₂ cathode electrodes were galvanostatically tested in a potential window of 2.25–4.35 V (vs. Li⁺/Li). The XRD patterns of the cycling-aged electrodes are shown in Fig. 8B. The XRD patterns remain essentially the same in all cases, indicating that atom loss from the lattice and layered-to-spinel phase conversion can hardly be induced by a low cut-off voltage. As explained above, the chemical and structural evolution in the Li_1.2Ni_0.15Co_0.1Mn_0.55O₂ material has an important effect on the lithium deintercalation/intercalation process. Compared with the first cycle, lithium extraction/intercalation can occur at a lower voltage in the subsequent cycles, probably due to the migration of the lithium/transition metal cations in Li/M slabs from their original crystallographic sites to other available sites during the charge/discharge process. The latter can lead to re-ordering or even loss of transition metal cations and the formation of a spinel phase in some regions.⁵⁴ The present XRD results provide insight into the structural transformation of the cycling-aged Li_1.2Ni_0.15Co_0.1Mn_0.55O₂ illuminating the voltage-fading mechanism.

3.3.6. 3D CT. The structural information about the internal structure and spatial distribution of the grain bulk of Li_1.2Ni_0.15Co_0.1Mn_0.55O₂ is confirmed by three-dimensional (3D) tomographic imaging. To illustrate this, virtual 3D cross-sectional slices of the grain bulk of Li_1.2Ni_0.15Co_0.1Mn_0.55O₂ particles are shown as yellow-coloured solid blocks in Fig. 9. In Fig. 9A, some cracks can be observed in the large secondary particles of the pristine Li_1.2Ni_0.15Co_0.1Mn_0.55O₂ sample. These particle fractures of the pristine Li_1.2Ni_0.15Co_0.1Mn_0.55O₂ sample resulted from the applied high pressure on the cathode during the fabrication process.⁵⁵

Fig. 9 Virtual 3D cross-sectional slices of the grain bulk of the Li_1.2Ni_0.15Co_0.1Mn_0.55O₂ particles before (A) and after 90 cycles (B).

It is noticeable, as shown in Fig. 9B, that fragmented pieces have been peeled off from the initial bulk particles, and thus the morphology of the secondary particles of Li_1.2Ni_0.15Co_0.1Mn_0.55O₂ changes significantly to small and irregular shaped sub-micron pieces with many voids after 90 charging–discharging cycles. It is worth noting that the progressive erosion on the electrode is not uniform with the depth and shape, which has been reported consistently by Chongmin Wang et al.⁵⁶ The intrinsic mechanism that causes such structural fragmentation of Li_1.2Ni_0.15Co_0.1Mn_0.55O₂ is related to the variation of the strain and stress due to the volume expansion and contraction during cycling.⁵⁷ In particular, the primary particles of Li_1.2Ni_0.15Co_0.1Mn_0.55O₂ undergo a lattice expansion/contraction strain caused by delithiation/lithiation during electrochemical cycling, resulting in numerous randomly distributed small voids due to the anisotropy and inhomogeneity of the primary particles. Upon deep electrochemical cycling, the discontinuous voids grow across forming cracks, and then the cracks propagate and link up to cause the fragmentation of the secondary particles of Li_1.2Ni_0.15Co_0.1Mn_0.55O₂. Eventually, the structural collapse of the grain bulk of Li_1.2Ni_0.15Co_0.1Mn_0.55O₂ is caused by the removal of the fragmented pieces of the secondary particles.

4. Conclusion

In this work, we investigated the capacity degradation and voltage fading mechanisms of a Li_1.2Ni_0.15Co_0.1Mn_0.55O₂ cathode as a representative Li-rich layered oxide cathode material for lithium ion batteries. Based on the in situ Raman and ex situ XANES characterization, the chemical evolution and structural transformation of the Li_1.2Ni_0.15Co_0.1Mn_0.55O₂ cathode are found to be quasi-reversible during a whole charge/discharge cycle. However, according to the results of the electrochemical measurements, HRTEM, SAED, TEM-EDS, EIS, XRD, XANES, and 3D CT of the pristine and cycling-aged Li_1.2Ni_0.15Co_0.1Mn_0.55O₂ cathode material, the activation of the anionic redox reaction could be an effective route to develop a high-capacity cathode. However, the anionic redox activity could also have various serious undesirable consequences, including irreversible layered-to-spinel structural transformation, atom disorder in the local lattice, oxygen vacancies, and cation dissolution, which eventually result in the lattice instability or even structural collapse of the grain bulk, and finally to voltage and capacity degradation. Our research findings also indicate that the promotion of the structural compatibility and stability, especially the stabilization of the oxygen lattice and reversible oxygen redox process, is the most critical factor towards constructing Li-rich materials with high capacity and good cycling stability.

Conflicts of interest

The authors declare no conflict of interest

Acknowledgements

This work was financially supported by the Basic Research Project of the Science and Technology Innovation Commission of Shenzhen (no. JCYJ20170412153139454), the National Natural Science Foundation of China (no. 21875097), Natural Science Foundation of Guangdong Province (no. 2018A030313725 and 2018B030322001), and the Professorial and Doctoral Scientific Research Foundation of Huizhou University (no. 2020JB025). In this work, XANES and Nano CT experiments were completed at the Shanghai Synchrotron Radiation Facility (SSRF) BL08U1A and the Beijing Synchrotron Radiation Facility (BSRF) 4 W1A, respectively. We thank Dr Zhenhua Chen, Dr Lijuan Zhang and Dr Xiangzhi Zhang from the SSRF for their help in XANES experiments and data analysis. We appreciate the support of the BSRF and the great help from Dr Kai Zhang and Dr Wanxia Huang.

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