Preparation and characterization of the Li_1.12K_0.05Mn_0.57Ni_0.24Nb_0.02O₂ cathode material with highly improved rate cycling performance for lithium ion batteries

Abstract

In this work, Li_1.12K_0.05Mn_0.57Ni_0.24Nb_0.02O₂ (LMN–K/Nb) as a novel and high energy density cathode material is successfully synthesized and applied in lithium ion batteries. By combining interlayer exchange and elemental analysis, it can be confirmed that K⁺ and Nb⁵⁺ substitution is respectively in the lithium layer and transition metal (TM) layer since H⁺ replaces the cations that remain in the lithium layer rather than those in the TM layer. The effect of K⁺ and Nb⁵⁺ co-substitution on the kinetic behavior of insertion/extraction of Li⁺ is evaluated by electrochemical impedance spectroscopy (EIS), the galvanostatic intermittent titration technique (GITT) and galvanostatic charge/discharge (GCD). LMN–K/Nb delivers an initial capacity of 145 mA h g⁻¹ at 5C rate and 112 mA h g⁻¹ at 10C rate, and maintains 83.1% after 400 cycles at 5C rate and 82.5% at 10C rate. By post-mortem analysis of long-term cycled LMN–K/Nb, K⁺ and Nb⁵⁺ are recognized to play a role in suppressing the irreversible side reactions in LLOs during cycling. This work demonstrates that dual elemental substitution into the lithium layer and TM layer is a feasible strategy to enhance the performance of LLO cathode materials.

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1. Introduction

Mn-Based Li-rich layered oxides (LLOs) have been considered as one of the most promising cathode materials due to their higher specific capacities (>250 mA h g⁻¹) and energy densities, and low cost in comparison with the commercial LiCoO₂, LiFePO₄ and ternary Ni–Co–Mn cathode materials.^1,2 For practical application in lithium ion batteries (LIBs), LLOs meet the problems of low initial coulombic efficiency, fast voltage fading, and inferior rate cycling stability,^3,4 mainly due to the reasons including dissolution of transition metal ions into the electrolyte, low conductivity and irreversible formation of the spinel phase arising from the transition metal ions migrating from the transition metal (TM) layer to the lithium layer during long charging/discharging cycles.^5,6

In the past few years, researchers have been found that elemental substitution/doping into the LLO host structure is an effective method to improve structural stability and electrochemical performance.^7–11 Two types of substitution/doping have been discussed in the literature, i.e. substitution of lithium ions in the lithium layer by alkali metals and replacing transition metal elements (i.e. Ni, Co and Mn) in the TM layer by hetero metal elements. Qing et al. reported that sodium substitution at lithium sites could improve the electrochemical performance of the 0.5Li₂MnO₃·0.5LiNi_1/3Co_1/3Mn_1/3O₂ material.⁹ Another work by Zheng et al. confirmed that potassium substitution at lithium sites of Li_1.232Mn_0.615Ni_0.154O₂ effectively enhanced the structural stability and electrochemical performance.¹⁰ In our previous study on Li_1.167−xK_xMn_0.583Ni_0.25O₂ (x = 0, 0.025, 0.05 and 0.075) materials, we also proved that K⁺ substitution results in small perturbations in the host structure and improved the lithium diffusion ability and hindered the detrimental interphase growth of the spinel phase during cycling.¹¹ The substitution by K⁺/Na⁺ with a larger radius is efficient in significantly improving the cycling performance of LLOs, but the specific capacity improvement is still limited.

Besides the heteroatomic substitution into the lithium layer, some studies have shown the substitution into the TM layer by alien ions, such as Al³⁺, Fe³⁺, Ca²⁺, Ru⁴⁺, Nb⁵⁺, etc.^12–16 Passerini's group studied Li_1.2Mn_0.585Ni_0.185Fe_0.03O₂ in which iron substitution effectively mitigated the voltage decay and capacity degradation.¹³ Dong et al. synthesized Li_1.2(Ni_0.13Co_0.13Mn_0.54)_1−xNb_xO₂ (x = 0, 0.01, 0.02, 0.04) in which appropriate Nb substitution accelerated the lithium-ion diffusion ability and stabilized the crystal structure.¹⁴ The substitution in the TM layer has significantly improved the specific capacity of LLOs, but an improvement in the long-term cycling performance still needs more investigation.

Since substitution in the lithium layer or TM layer brings different benefits to the host material, co-substitution with the aim of utilizing synergistic effects from two or more substituted elements would be a promising strategy to improve the material performance.^17–19 For example, Na and Al co-substitution has been reported to be an effective way to reduce the voltage decay of LLOs and inhibit the side reaction between the cathode materials and electrolyte during cycling.¹⁷ The co-substitution via Mg and La respectively in the lithium layer or TM layer has been proved to improve the structural stability and electrochemical performance of Li(Li_0.2Mn_0.54Co_0.13Ni_0.13)O₂.¹⁸ To the best of our knowledge, there is no research on the co-substitution of LLOs with K⁺ and Nb⁵⁺, the structural stability mechanism and the improved properties of LLOs from the synergistic effect of K⁺ and Nb⁵⁺ co-substitution is unknown.

In this work, K/Nb co-substituted Li_1.12K_0.05Mn_0.51Ni_0.24Nb_0.02O₂ is successfully synthesized, and its structural stability and electrochemical performance are compared with those of Li_1.17Mn_0.52Ni_0.25O₂. The synergistic effect of K/Nb co-substitution on structural stability is revealed through inductively coupled plasma-atomic emission spectrometry (ICP-AES), X-ray powder diffraction (XRD), scanning and transmission electron microscopy (SEM and TEM), X-ray photoelectron spectroscopy (XPS), electrochemical impedance spectroscopy (EIS), galvanostatic charge/discharge (GCD), etc. At 2C rate, LMN–K/Nb delivers 164 mA h g⁻¹ at the 1^st cycle and retains 98% of its initial capacity after 200 cycles, while LMN delivers only 123 mA h g⁻¹ at the 1^st cycle and retains only 65% of its initial capacity at the 200^th cycle. This work reports an approach for dual-ion substitution and provides it as an effective way to develop LLO cathode materials with stable cycling performance for energy storage applications.

2. Experimental section

2.1. Materials preparation

In this work, Li₂CO₃, MnO₂, NiCO₃ and absolute ethanol were bought from Sinopharm Chemical Reagent Co., Ltd and used without any pretreatment. Li_1.17Mn_0.58Ni_0.25O₂ and Li_1.12K_0.05Mn_0.57Ni_0.24Nb_0.02O₂ were prepared through a wet ball mill and high-temperature treatment. The parameters of ball milling are as follows: absolute ethanol used as a dispersant for the wet ball mill, NiCO₃, MnO₂ and Li₂CO₃ reagents in stoichiometric amounts mixed using a planetary ball mill (QM-1SP4-CL, Nanjing, China). Herein, a 5 wt% extra amount of lithium source besides its stoichiometric amount is used, compensating for lithium loss during high temperature treatment. The weight ratio of the milling ball with reactants, and the milling time and speed are 8 : 1, 10 h and 360 rpm, respectively. The precursor after ball milling was heated at 450 °C for 3 h and then at 900 °C for 12 h under an air atmosphere. The as-produced Li_1.17Mn_0.58Ni_0.25O₂ was simply named LMN. By using K₂CO₃ and Nb₂O₅ as the potassium source and niobium source, Li_1.12K_0.05Mn_0.57Ni_0.24Nb_0.02O₂ was synthesized by the same route and was named LMN–K/Nb.

2.2. Characterization

Crystallographic structures of LMN and LMN–K/Nb were obtained by powder X-ray diffraction (XRD, Dmax-2200, Cu-Kα radiation) in the angle range of 10°–80° at 0.02° per scanning step and 4s counting time at each step. The compositions of Li, Ni, Mn, K and Nb were obtained by inductively coupled plasma atomic emission spectrometry (ICP-AES, Plasma-400, PerkinElmer). The surface morphologies and microstructures of the samples were characterized by scanning electron microscopy (SEM, ZEISS Sigma) and transmission electron microscopy (TEM, JEM-2100). The surface elemental valences of samples were analyzed by X-ray photoelectron spectroscopy (XPS, Thermo Fisher Scientific, Escalab250Xi).

2.3. Electrochemical measurements

The electrochemical properties of the samples were measured by both types of half-cell (CR2016-type coin cell) and full cell (pouch-type cell). The positive electrode was composed of active materials, polyvinylidene fluoride (PVdF) and carbon black in a weight ratio of 8 : 1 : 1 and N-methyl-2-pyrrolidone (NMP) as a solvent. The mixture was stirred for 24 h, and the slurry was coated on Al foil. The electrode was dried at 100 °C for 12 h under vacuum. In a half-cell, the area and thickness of the positive electrode are 1.21 cm² and 25 μm, respectively. The mass loading of the electrode is 1.8 mg cm⁻². By using metallic lithium foil as a negative electrode, Celgard 2500 as a separator, 1 M LiPF₆ dissolved in dimethyl carbonate (DMC), ethyl methyl carbonate (EMC) and ethylene carbonate (EC) (1 : 1 : 1 in volume) as an electrolyte, the coin cells were assembled in an argon-filled glove box where both water and oxygen concentrations are less than 0.1 ppm. Electrochemical impedance spectroscopy (EIS) was performed on an electrochemical workstation (PARSTAT 2273, Princeton, USA), and the frequency range was from 10⁵ Hz to 10⁻² Hz. Using the galvanostatic intermittent titration technique (GITT), the curves of the cathode material during the 3^rd cycle were recorded at a time interval of 30 min by using a LAND CT2001A battery testing system.

Considering the practical application, we have assembled a full cell (pouch type cell) to evaluate the electrochemical properties of LMN–K/Nb. In this work, commercial graphite (Sinopharm Chemical Reagent Co., Ltd) was used as a negative electrode in a pouch-type cell. The positive and negative electrodes were cut in squares, and the side lengths of the squares for the negative and positive electrodes are 35 and 40 mm, respectively. The designed capacity of the negative electrode was 1.1 times that of the positive electrode. The slurry of the negative electrode was prepared by mixing 80 wt% of commercial graphite, 10 wt% of PVdF and 10 wt% of carbon black. Generally, secondary sealing is necessary for lithium-rich cathode materials in a pouch-type full cell, because the gas accumulated after several charge/discharge cycles swells the bag and must be removed. In this work, secondary sealing for the pouch-type full cell was applied after 5 charging/discharging cycles to remove the gas accumulated during the redox reaction. A LAND CT2001A battery testing system (Wuhan LANHE, P.R. China) was used to record the galvanostatic charge/discharge profiles at 25 °C.

3. Results and discussion

Fig. 1a shows the XRD patterns of LMN and LMN–K/Nb. Each sample presents clear and sharp reflections, indicating that both samples are high crystallized. Most reflections can be indexed to the hexagonal NaFeO₂ structure (R3̄m), besides the superlattice diffraction peaks between 20 and 25° that correspond to the Li₂MnO₃ phase (C2/m).^20,21 No impurity peaks are found in LMN–K/Nb, indicating that K⁺ and Nb⁵⁺ are successfully substituted into the crystal lattice without affecting the host structure of LLOs. As shown in the inset image in Fig. 1a, the (003) peak of LMN–K/Nb shifts by a few degrees lower than that of LMN, indicating that the interlayer spacing of LMN–K/Nb has a slightly bigger value than that of LMN. This may be due to the fact that the intercalated K⁺ expands the interlayer spacing of LMN–K/Nb because K⁺ (1.38 Å) has a much larger ionic radius than Li⁺ (0.76 Å). The expanded interlayer spacing can increase the diffusion ability of lithium ions.²²

Fig. 1 XRD patterns (a), XPS survey spectra (b), and high-resolution XPS spectra of Mn 2p (c) and Ni 2p (d) of LMN and LMN–K/Nb. The K 2p (e) and Nb 3d (f) XPS spectra of LMN–K/Nb at different etching times (seconds).

In order to further study the structure of LMN and LMN–K/Nb, the XRD patterns were refined based on the two-phase mode of LiNi_0.5Mn_0.5O₂ (R3̄m) and the Li₂MnO₃ monoclinic phase (C2/m) by using the General Structure Analysis System (GSAS). The values of the R_wp factor (less than 10%) confirm the reasonable refinement.²⁰ The intensity ratios I₍₀₀₃₎/I₍₁₀₄₎, as summarized in Table S1,† are greater than 1.2, indicating that both samples have an ordered layered structure with low cation mixing.²² The smaller I₍₀₀₃₎/I₍₁₀₄₎ value of LMN–K/Nb than that of LMN again indicates the substitution of K⁺ and Nb⁵⁺ into the lattice structure of LLOs, which reduces the ordering layered structure to some extent. As detailed in the inset in Fig. S1,† the well-splitting reflections of (006)/(102) and (108)/(110) show that both samples have a typical layered structure.²⁰ In addition, the lattice parameter c of LMN–K/Nb (14.2754 Å, R3̄m phase) is higher than that of LMN (14.2608 Å), which corresponds to the (003) peak shift of LMN–K/Nb in Fig. 1a. The large interplanar spacing of LMN–K/Nb is conducive to the rapid diffusion of Li⁺, thereby showing excellent rate performance.

The Li : K : Mn : Ni : Nb atomic ratios for LMN and LMN–K/Nb were measured by ICP. The values are respectively Li_1.17Mn_0.59Ni_0.25 and Li_1.11K_0.05Mn_0.58Ni_0.24Nb_0.02 (the value after each metal element is the calculated atomic ratio), and they are very close to the theoretical values of Li_1.17Mn_0.58Ni_0.25 and Li_1.12K_0.05Mn_0.57Ni_0.24Nb_0.02 for LMN and LMN–K/Nb, respectively. Until now, it was a challenge to directly prove the substitution of heteroelements into the lithium layer or TM layer of LLOs. The cations that remain in lithium layers could be exchanged with H⁺ by soaking the samples in acidic solution, instead of those in TM layers.²³ In this work, the substituted site of K⁺ and Nb⁵⁺ in the lithium layer and TM layer is studied by using this method. To do so, 500 mg of the sample was dispersed in 500 mL of HCl solution (1 mol L⁻¹) under magnetic stirring for 72 h, and the HCl solution was renewed every 24 h. Then, the sample was centrifuged, washed with deionized water and dried at 80 °C. Before and after H⁺ exchange, the atomic ratios of metal elements were measured by ICP. The calculated atomic ratios for LMN and LMN–K/Nb after H⁺ exchange are Li_0.36Mn_0.57Ni_0.25 and Li_0.30K_0.01Mn_0.58Ni_0.24Nb_0.02, respectively. Lithium atomic ratios in both LMN and LMN–K/Nb dramatically reduced to about 0.3, and the K atomic ratio in LMN–K/Nb significantly decreased from 0.05 to 0.01 after H⁺ exchange. However, the Nb atomic ratio in LMN–K/Nb shows no change before and after H⁺ exchange. Since H⁺ replaces the cations that remain in the lithium layer rather than those in the TM layer, the aforementioned results may strongly confirm that K⁺ and Nb⁵⁺ substitution is respectively in the lithium layer and TM layer.

The chemical valences of elements at the sample surface and near-surface are evaluated by XPS coupled with Ar⁺ etching. The survey spectra of LMN and LMN–K/Nb as shown in Fig. 1b present the relative metal elements without any impurities. The Mn 2p spectra (Fig. 1c) of LMN and LMN–K/Nb show the same characteristic peaks respectively corresponding to Mn 2p_1/2 at 653.0 eV and Mn 2p_3/2 at 641.5 eV. The Ni 2p spectra (Fig. 1d) of LMN and LMN–K/Nb show the same characteristic peaks at 855.0 and 872.5 eV, which correspond to Ni 2p_3/2 and Ni 2p_1/2, respectively. The peak positions of Mn 2p and Ni 2p in LMN and LMN–K/Nb corresponding to the presence of Mn⁴⁺ and Ni²⁺ are very close to the previous report,²⁴ and the results indicate that the substitution has no effect on the valence state of Mn and Ni.

For LMN–K/Nb, the peaks at 292.4 eV and 295.5 eV in K 2p spectra (Fig. 1e) assigned to K 2p_3/2 and K 2p_1/2,²⁵ and the peaks at 206.3 eV and 209.1 eV in Nb 3d spectra (Fig. 1f) are respectively assigned to Nb 3d_5/2 and Nb 3d_3/2.²⁶ In addition, Fig. 1e and f show the K 2p and Nb 3d peaks of LMN–K/Nb at different Ar⁺ etching times of 0, 150 and 300 s (the relative etching depth in the range of 0–150 nm). Because of the same quantity of K and Nb elements detected at different depths of the sample crystal, there is strong evidence that K and Nb elements are uniformly incorporated into the structure but do not simply reside on the sample surface.

The crystal morphology of LMN and LMN–K/Nb are characterized by SEM, TEM and HRTEM (Fig. 2). The particles of both samples are homogeneous with similar crystal shapes and the sizes in the range of 200–400 nm (Fig. 2a, b, a′ and b′). The results show that the K⁺ and Nb⁵⁺ substitution plays a small role in crystal morphology. As shown in the HRTEM images in Fig. 2a′′ and b′′, both samples show clear lattice fringes. The lattice distances for LMN and LMN–K/Nb are 0.471 and 0.479 nm corresponding to the (003) crystal plane.² A minor bigger lattice distance of (003) in LMN–K/Nb than that in LMN may be due to K⁺ with a bigger ionic radius than Li⁺ in lithium layers.

Fig. 2 SEM, TEM and HRTEM images of LMN (a, a′ and a′′) and LMN–K/Nb (b, b′ and b′′).

The effects of K⁺ and Nb⁵⁺ co-substitution on the electrochemical properties of LMN–K/Nb are studied in comparison with LMN. Fig. 3a and b show the Nyquist plots of half cells with LMN and LMN–K/Nb as positive electrodes, and the cells have been discharged to 3.5 V at the 3^rd cycle and 50^th cycle, respectively. The Nyquist plots are well refined by an equivalent circuit in Fig. S2,† and the exact values of R_e (ohmic resistance), R_sf (resistance of SEI film) and R_ct (charge transfer resistance) are summarized in Table S2.† It is noted that both samples show no obvious change in R_e values during cycling. However, the R_sf and R_ct values of both LMN and LMN–K/Nb at the 50^th cycle are increased two times those at the 3^rd cycle, arising from their phase change from the layered to the spinel structure and irreversible side reactions during cycling.^11,27 For example, the R_ct values of LMN and LMN–K/Nb are increased from 505.71 and 90.63 Ω at the 3^rd cycle to 1353.32 and 266.16 Ω at the 50^th cycle. At the 3^rd and 50^th cycles, LMN–K/Nb shows much smaller R_sf and R_ct values than LMN, which might be due to the fact that the synergistic effect of K⁺ and Nb⁵⁺ co-substitution effectively increases the structural stability and lithium diffusion ability.

Fig. 3 Nyquist plots of samples discharged to 3.5 V at the 3^rd cycle (a) and 50^th cycle (b). GITT curves (time interval: 30 min) at the 3^rd cycle (c) and the calculated diffusion coefficients of Li⁺ at 0.1C (d).

The effect of K⁺ and Nb⁵⁺ co-substitution on the kinetic behavior of insertion/extraction of Li⁺ is further evaluated by GITT. Each current pulse (charge or discharge) of 10 min at a current density of 0.1C followed by keeping at a 30 min open circuit was repeatedly employed until the charge/discharge voltages reach 4.9 V/2.5 V, respectively. Based on the GITT curves of each sample at the 3^rd cycle (Fig. 3c), the lithium diffusion coefficient of Li⁺ (D_Li⁺) is calculated according to the following equation:^11,28

wherein m and M respectively represent mass (g) and molecular mass (g mol⁻¹) of the active materials, V_M is the molar volume deduced from the crystallographic data (cm³ mol⁻¹), A is the surface area (cm²) of the electrode, ΔE_s is the change in the steady-state voltage and ΔE_τ is the voltage change in a single step GITT experiment. The D_Li⁺ values of LMN and LMN–K/Nb at various states of discharge are plotted in Fig. 3d. It can be seen that the D_Li⁺ values of both samples are slowly decreased from ∼1.2 × 10⁻¹¹ to ∼1.5 × 10⁻¹² cm² s⁻¹ during the discharging process, and LMN–K/Nb shows a higher D_Li⁺ value than LMN. According to numerous research results,^22,27,28 it can be concluded that the extended interlayer spacing can improve the Li⁺ diffusion capability. Since the radius of K⁺ (1.38 Å) is larger than that of Li⁺ (0.76 Å), the LMN–K/Nb with an extended interlayer spacing exhibits a higher diffusion coefficient of Li⁺, which helps it to exhibit excellent rate performance.

To further ascertain the cause for improved cycling and rate performance by K⁺ and Nb⁵⁺ co-substitution, galvanostatic charging and discharging of LMN and LMN–K/Nb were performed. The cells are cycled in the voltage range of 2.5–4.9 V at various current rates. The initial charge–discharge curves at 0.1C (1C = 200 mA g⁻¹) of both samples present two charging steps (Fig. 4a), i.e. the charging below 4.5 V corresponds to the reaction of Ni²⁺/Ni⁴⁺ in the LiNi_0.5Mn_0.5O₂ layered phase, and the other voltage plateau of ∼4.6 V is attributed to the extraction of Li⁺ and O evolution from the lithium rich Li₂MnO₃ phase.²⁹ The latter often shows poor reversibility and it leads to a large irreversible capacity in the first cycle. The initial specific discharge capacities of LMN and LMN–K/Nb are 211 and 244 mA h g⁻¹, and the coulombic efficiencies (CEs) are 59.6 and 77.2%, respectively. Since Nb could offer stronger metal oxygen bonds, which is beneficial for suppressing the O evolution of LLOs.¹⁴

Fig. 4 The 1^st charge/discharge curves (a), the cycling performance at 0.1C and 2C (b), and 5C and 10C (c), and the cycling performance of pouch-type full cells at 0.1C (d).

Fig. 4b and c show a comparison of the cycling performance and coulombic efficiency (CE) of LMN and LMN–K/Nb at various current rates. LMN–K/Nb presents much higher specific capacity than LMN at each rate. At 0.1C, LMN–K/Nb delivers 246 mA h g⁻¹ at the initial cycle and 231 mA h g⁻¹ at the 100^th cycle, while LMN only delivers 209 mA h g⁻¹ at the 1^st cycle and 203 mA h g⁻¹ at the 100^th cycle. At 2C rate, LMN–K/Nb delivers 164 mA h g⁻¹ at the 1^st cycle and retains 98% of its initial capacity after 200 cycles, while LMN delivers only 123 mA h g⁻¹ at the 1^st cycle and retains only 65% of its initial capacity at the 200^th cycle (Fig. 4b).

At higher current rates of 5C and 10C, LMN–K/Nb delivers two times the specific capacity than LMN through 400 cycles (as shown in Fig. 4c). At 5C, LMN–K/Nb delivers 145 mA h g⁻¹ at the 1^st cycle and retains 83.1% of its initial capacity at the 400^th cycle, while LMN delivers only 75 mA h g⁻¹ at the 1^st cycle and retains only 60% of its initial capacity at the 400^th cycle (Fig. 4c). At 10C rate, the specific capacity of LMN decreases to 48 mA h g⁻¹ and 28 mA h g⁻¹ at the 1^st and 400^th cycles, respectively, while LMN–K/Nb still delivers a high capacity of 112 mA h g⁻¹ at the 1^st cycle and retains 92 mA h g⁻¹ at the 400^th cycle. The much better rate cycling performance of LMN–K/Nb than LMN should be attributed to the success of K⁺ and Nb⁵⁺ co-substitution.

To evaluate the practical application of the synthesized cathode materials, full cells (pouch type cell) with the prepared materials as a positive electrode and commercial graphite as a negative electrode are assembled. A formation cycle at a current rate of 0.01C is expected as the solid electrolyte interphase (SEI) layer is grown.³⁰ The cycling performance of full cells is measured between 2.5 and 4.9 V at 0.1C rate, and the cycling performance is presented in Fig. 4d. LMN–K/Nb shows much higher capacity than LMN, especially, the capacity retention of LMN–K/Nb is 88.3% after 100 cycles which is much higher than that of LMN of only 64.8%. As shown in Fig. S3,† the rate performance of the LMN and LMN–K/Nb samples is evaluated at 0.1C, 0.2C, 0.4C, 0.5C, 0.8C and 1C. By increasing the current density from 0.1 to 1C, the discharge specific capacity of the samples is relatively reduced. For example, the discharge capacities of LMN are 196.6, 171.1, 149.0, 134.2, 111.9, 94.3 and 166.2 mA h g⁻¹, but LKMNO delivers higher values of 216.6, 194.5, 162.5, 147.8, 130.3, 117.5 and 205.8 mA h g⁻¹ at 0.1, 0.2, 0.4, 0.5, 0.8, 1 and 0.1C, respectively. The full cell performance further confirms that K⁺ and Nb⁵⁺ co-substitution efficiently improves the capacity contribution and cycling stability of LLO cathode materials.

It is well known that IR drop is an important criterion to evaluate the electrochemical performance of electrode materials during redox reactions.^13,31 As shown in Fig. 5a and b, the IR drop values of LMN and LMN–K/Nb are respectively 0.26 and 0.14 V from the 1^st to 200^th cycles at 2C rate. Compared with the discharge curves for hundreds of cycles, the discharge voltages of LMN and LMN–K/Nb have all decreased along with an increase in charging/discharging cycles. The starting discharge voltages of LMN and LMN–K/Nb decrease from the 1^st cycle to the 400^th cycle at 5C rate and the decay values are respectively 0.43 and 0.27 V (Fig. 5c). A similar trend of the sample at 10C rate can be found in Fig. 5d. LMN–K/Nb presents a smaller IR drop than LMN, indicating that K/Nb co-substitution effectively mitigates the voltage decay and capacity degradation. In addition, Fig. 5e shows a comparison of the specific capacity and energy density of LMN–K/Nb with the reported LLO cathode materials in detail. Herein, more cobalt-free LLOs are deliberately selected and different modification methods are selected as examples, such as substitution by V, F and Nd/Al and coating by CoF₂, Li₂ZrO₃ and Li₄Ti₅O₁₂.^32–42 It is evident that LMN–K/Nb exhibits higher specific capacity and energy density than most of those previously reported. Therefore, the results show that K⁺ and Nb⁵⁺ co-substituted LLO cathode materials have huge advantages in practical applications, especially in devices that require high energy densities.

Fig. 5 Discharge curves at the 1^st, 50^th, 100^th, 200^th and 300^th cycles of LMN (a) and LMN–K/Nb (b) at 2C rate, and the histogram shows a variation trend of the starting discharge voltage after different cycles at 5C (c) and 10C (d) rates, respectively. The median voltage against specific capacity and energy density of selected LLO cathode materials (e).

In order to understand the effect of K⁺ and Nb⁵⁺ co-substitution on enhancing the structural stability of LLOs during cycling, the differential capacity curves (vs. voltage dQ/dV) of LMN and LMN–K/Nb at various cycles (1^st, 5^th, 10^th, 20^th, 30^th and 50^th cycles) are plotted in Fig. 6. The peaks denoted as P₁ and P₂ were observed at 3.8 and 4.6 V during the first charging process. P₁ is attributed to the redox reaction of Ni²⁺/Ni⁴⁺ in the LiNi_0.5Mn_0.5O₂ phase, and P₂ corresponds to an irreversible process of Li extraction and O₂ release from the Li₂MnO₃ structure,⁴³ which disappeared after the first cycle. From the 2^nd cycle, a new peak at 3.35 V (denoted as P₃) is observed in each sample, which is attributed to the redox reaction of Mn³⁺/Mn⁴⁺.^29,44 During discharge, three peaks are observed at 4.31, 3.79 and 3.25 V (respectively denoted as P₄, P₅ and P₆). P₄ is related to Li occupation at the tetrahedral sites, and P₅ is ascribed to Li occupation at the octahedral sites accompanied by the Ni^4+/3+/2+ and Mn⁴⁺/Mn³⁺ redox reaction.²² P₆ not only corresponds to the redox reaction of Mn³⁺/Mn²⁺ but also can be used to evaluate the phase transition from the layered structure to the spinel structure.⁴⁵ The intensity and displacement of the P₆ peak are directly related to the degree of phase transition.

Fig. 6 The dQ/dV profiles of LMN (a) and LMN–K/Nb (b) at 0.1C; schematic diagrams of the possible phase evolution routes for LMN (c) and LMN–K/Nb (d), and tetrahedral sites in the Li layer (e).

The appearance of Mn³⁺ in LLOs during discharge along cycling leads to the disproportionation reaction and the formation of Mn⁴⁺ and Mn²⁺.²⁸ The as-produced Mn²⁺ easily migrates into the lithium layer to form the spinel phase during the delithiation process. From the 1^st cycle to the 50^th cycle, LMN–K/Nb presents a weaker peak intensity and smaller deviation of the P₆ peak than LMN (Fig. 6a and b), indicating the better structural stability of LMN–K/Nb. The hindering effect of K⁺ substitution on the formation of the spinel phase is explained in Fig. 6c–e. For LMN, during the charging (delithiation) process, the removal of lithium ions results in the formation of trivacancies, and adjacent Mn ions tend to migrate into the tetrahedral vacancies surrounded by these trivacancies. As a result, the spinel phase is irreversibly formed (Fig. 6c and d).²⁵ Because K⁺ has a larger ion radius than Li⁺, K⁺ substitution into the Li layer plays a pillar role in preventing the migration of Mn²⁺ from the TM layer to the lithium layer (Fig. 6d). This effectively inhibits the formation of the spinel phase and improves the cycling performance of LMN–K/Nb.

The structural changes of the cycled LMN and LMN–K/Nb electrodes are evidenced by ex situ XRD. Fig. 7 shows the XRD patterns and enlarged characteristic peaks of both electrodes at the 1^st, 10^th and 50^th cycles. The layer-spinel phase transition can be clearly evaluated by the (003) and (104) peak shifts.⁴⁶ After the 1^st cycle, the (003) peak of LMN observed at 18.3° (2θ) still represents a layered structure; after 10 cycles it shifted to 18.1° (2θ), representing the mixed structure of the layer and spinel; and after 50 cycles it shifted to 17.9° (2θ) arising from the major spinel structure (Fig. 7a′). A similar trend of structural transition with cycling is also found by the (104) peak shifting (Fig. 7a′′). The characteristic layer structure of LMN–K/Nb was excellently maintained after dozens of cycles at the unchanged (003) and (104) peak positions (Fig. 7b′ and b′′). This clearly reveals that K⁺ and Nb⁵⁺ co-substitution can effectively suppress the irreversible spinel growth in the LLO cathodes during cycling. After dozens of cycles, the grown spinel phase in LMN and the unchanged layer structure in LMN–K/Nb are in accordance with the aforementioned P₆ in their dQ/dV results (Fig. 6a and b). The aforementioned results indicate that LMN has a large amount of Mn²⁺ dissolution, accompanied by phase transition from the layer to the spinel structure, and finally results in structural disruption and rapid capacity fading. For K⁺ and Nb⁵⁺ co-substitution, the structural stability and electrochemical properties of LMN–K/Nb are significantly improved, in which K⁺ substitution in the lithium layer effectively inhibits the phase transition and the Nb–O bonds arising from Nb⁵⁺ substitution effectively inhibits oxygen evolution and reduces the dissolution of Mn cations during cycling.

Fig. 7 XRD patterns and the enlarged characteristic peaks of LMN (a, a′ and a′′) and LMN–K/Nb (b, b′ and b′′) after different cycles, and both samples are discharged to 2.5 V.

Taking the above results into account, the significant improvement in the electrochemical performance of LMN–K/Nb can be attributed to the following synergistic effects: (1) K⁺ substitution in the lithium layer expands the interlayer spacing of the lithium layer, thereby showing excellent rate performance. (2) K⁺ can also inhibit the migration of transition metal ions, thus inhibiting the irreversible phase transition and improving the cycling stability. (3) The binding energy of Nb–O is higher than Ni–O and Mn–O, so Nb⁵⁺ substitution into the TM layer can inhibit the O evolution and stabilize the host material structure. Furthermore, K⁺ and Nb⁵⁺ are also recognized to play a role in suppressing the irreversible side reactions during cycling.

4. Conclusion

In this work, K⁺ and Nb⁵⁺ are successfully substituted into the lithium layer and TM layer of Li-rich cathode materials, and LMN–K/Nb shows improved structural stability and electrochemical properties in comparison with the pristine sample. It delivers initial capacities of 145 mA h g⁻¹ and 112 mA h g⁻¹ at current rates of 5C and 10C, respectively, and shows capacity retentions of 83.1% after 400 cycles at 5C. K⁺ substitution in the lithium layer expands the interlayer spacing of the lithium layer and inhibits irreversible phase transition, thus leading to improved rate and cycling stability. The binding energy of Nb–O is higher than those of Ni–O and Mn–O, so Nb⁵⁺ substitution into the TM layer can inhibit the O evolution and stabilize the material structure. Furthermore, K⁺ and Nb⁵⁺ are also recognized to play a role in suppressing the irreversible side reactions during cycling.

Conflicts of interest

There are no conflicts to declare.

Acknowledgements

The authors greatly appreciate the financial support from the National Natural Science Foundation of China (grant number 51802030) and the Natural Science Foundation of Jiangsu Province, China (Grant No. BK20161267).

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Footnote

† Electronic supplementary information (ESI) available. See DOI: 10.1039/d1nr06824e

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