Unique insights into the design of low-strain single-crystalline Ni-rich cathodes with superior cycling stability

Significance

In order to support the growth of electrochemical energy storage applications, improved cathode materials of Li-ion batteries are required to unlock a longer lifespan in combination with enhanced energy density. Developing micro-sized single-crystalline Ni-rich cathodes (SCN) has emerged as the mainstream to improve the volumetric energy density and safety. However, uneven Li-ion distribution and stress concentration result in intragranular crack generation with a limited cycle life. Herein, an optimal particle size is predicted by simulating the stress distributions and a unique strategy is proposed to synthesize the target-sized SCN (m-NCM83). Accordingly, the m-NCM83 exhibits superior cycling stability with high structural integrity. This work provides crucial guidance for the design and synthesis of high-energy density SCN with superior cycling stability.

Abstract

Micro-sized single-crystalline Ni-rich cathodes are emerging as prominent candidates owing to their larger compact density and higher safety compared with poly-crystalline counterparts, yet the uneven stress distribution and lattice oxygen loss result in the intragranular crack generation and planar gliding. Herein, taking LiNi_0.83Co_0.12Mn_0.05O₂ as an example, an optimal particle size of 3.7 µm is predicted by simulating the stress distributions at various states of charge and their relationship with fracture free-energy, and then, the fitted curves of particle size with calcination temperature and time are further built, which guides the successful synthesis of target-sized particles (m-NCM83) with highly ordered layered structure by a unique high-temperature short-duration pulse lithiation strategy. The m-NCM83 significantly reduces strain energy, Li/O loss, and cationic mixing, thereby inhibiting crack formation, planar gliding, and surface degradation. Accordingly, the m-NCM83 exhibits superior cycling stability with highly structural integrity and dual-doped m-NCM83 further shows excellent 88.1% capacity retention.

Developing micro-sized single-crystalline Ni-rich cathodes, LiNi_xCo_yMn_1-x-yO₂ (x ≥ 0.8), has emerged as the mainstream to further improve the volumetric energy density and safety of lithium-ion batteries (1–3). Such cathodes can fully eliminate the grain boundary effects among primary particles of traditional polycrystalline microspheres (4, 5). This merit not only greatly boosts their compaction density by decreasing the porosity during the rolling process but also markedly inhibits the electrolyte deterioration and the rock-salt phase generation on the surface by minimizing the cathode-electrolyte interface side reactions and the lattice oxygen loss (6). It should be noted that the particle size directly impacts their capacity retention at different rates and cycle life (7–9). Oversized single-crystalline Ni-rich cathodes will suffer from extended Li-ion diffusion paths, easily causing the uneven Li-ion distribution and stress concentration during the charge/discharge process with the formation of intragranular cracks (10–12). Recently, Xiao et al. employed a diffusion-induced stress model by density functional theory calculations to predict the optimal single-crystalline particle size of LiNi_0.76Mn_0.14Co_0.1O₂ (13). The result indicated that the strain energy generated by lithium-ion diffusion is consistently less than the fracture energy when the practical size is controlled below 3.5 µm, thus avoiding the irreversible crystal-plane gliding and microcracks. However, the above theoretical prediction has not yet been verified by experiments to date. This is primarily because the lithiation temperature of single-crystalline Ni-rich cathodes is typically ~100 °C higher than that of polycrystalline counterparts (14, 15). The resulting serious Li/O loss makes it difficult to synthesize highly ordered layered single-crystalline Ni-rich cathodes (16, 17).

At present, some reports found that the addition of fluxing agents such as sulfates and nitrates can lower the melting point of lithium salts (18–21), and meantime, some additives like cerium oxide and strontium oxide can regulate the surface energy of crystal facets (7, 22). These two ways are helpful for realizing the synthesis of single-crystalline Ni-rich cathodes at relatively low lithiation temperature by expediting the nucleation and growth of grains. Nevertheless, the use of such additives will inevitably introduce electrochemically inert impurities, which will lead to poor thermal stability and severe electrochemical polarization especially at high rates (23, 24). The production costs will also increase. Typically, the washing process is usually involved to remove fluxing agents from single-crystalline Ni-rich cathodes (25–27). However, Jeff et al. discovered that this process triggers the severe Li⁺/H⁺ proton exchange, causing partial loss of active lithium, narrowing of the interlayer spacing along the c-axis, and formation of electrochemically inert rock-salt phases (28). These negative factors greatly reduce the reversible capacity and retard Li-ion diffusion kinetics. Therefore, the size optimization and the corresponding synthesis technique of high-quality single-crystalline Ni-rich cathodes have not been well-solved so far.

In this work, we first predict the optimal particle size of single-crystalline LiNi_0.83Co_0.12Mn_0.05O₂ is 3.7 µm by elucidating the stress distributions as well as the relationship between strain energy and fracture free-energy under various states of charge. After building the intrinsic correlations among lithiation temperature, time, and particle size, a unique strategy is proposed to synthesize the target-sized single-crystalline Ni-rich cathodes (m-NCM83) with low cationic mixing and poor oxygen vacancies by introducing high-temperature short-duration pulses (975 °C, 10 min) into the 800 °C lithiation process. The m-NCM83 significantly reduces the stress concentration and Li/O loss, exhibiting superior cycling stability with highly structural integrity even under harsh operation conditions.

Results and Discussion

Failure Mechanism and Size Theoretical Prediction of Single-Crystalline Ni-Rich Cathodes.

The degradation mechanism of single-crystalline Ni-rich cathodes mainly comes from two aspects. First, the prolonged high-temperature lithiation step inevitably causes the spontaneous reduction of unstable Ni³⁺ ions on particle surface with abundant oxygen vacancies arising (29), which will aggravate the migration of transition metal ions and the resultant lattice gliding during de-/lithiation, as shown in Fig. 1A. Second, there is a great difference in Li-ion concentration distribution from surface to bulk phase because of a long diffusion path in micro-sized single-crystalline particles (30, 31). The inhomogeneous phase transition and lattice strain within the particles occurs successively, giving rise to the stress accumulation and concentration (32). As shown in Fig. 1B, the mechanical structure of the particles is eventually destroyed with the generation of microcracks. Therefore, the suitable particle size is crucial for improving the cycling stability of single-crystalline Ni-rich cathodes. To fill the gap in addressing the two issues, taking NCM83 as an example (33, 34), we first built a diffusion-induced stress model based on Fick’s second law and current distribution theory (SI Appendix, Fig. S1). The distribution of Li-ion concentration within a single particle was simulated under various SOCs (states of charges) (Fig. 1C). It is observed that the uneven Li-ion distribution is most obvious at 10% SOC. The corresponding stress distributions were given based on the COMSOL simulation that Li-ion diffusion and intercalation models were established by the dynamic change of Young’s modulus (34). Fig. 1D shows that tangential stress peak values always appear near the particle surface under different SOCs. Among them, the largest stress happens at 10% SOC, which was selected to visually display the 3D internal stress distribution in Fig. 1E. The above simulation results highlight that the uneven Li-ion distribution is a culprit to stress concentration.

Fig. 1.

Failure mechanism and size theoretical prediction of single-crystalline Ni-rich cathodes. Schematic illustration of (A) the crystal-plane gliding caused by oxygen loss, and (B) intragranular cracks due to inhomogeneous Li-ion distribution. (C) Li-ion concentration distribution and (D) tangential stress distribution of NCM83 at various SOC. (E) 2D visualization of the stress distribution in NCM83 at 10% SOC. (F) Fitting results of Li-ion concentration distribution at 10% SOC. (G) Surface energy and strain energy curves as a function of particle sizes. (H) Gibbs free energy of particle fracture at different sizes.

According to Fick's second law, the mathematical relationship between Li-ion concentration (C_Li) and particle size (R) was obtained by fitting the simulated data at 10% SOC, delivering a typical error function relationship in Fig. 1F. This result can quantify the stress concentration within the particle and identify the key parameters affecting crack generation. Specifically, the strain energy of NCM83 can be identified as:

$$\begin{array}{c}{\left.\mathrm{\Pi }\right|}_{T_P}=\frac{4}{3}\times \pi \times R\times \left[\frac{a{\times E}_0\times {\left(C_R-C_0\right)}^2}{1-v}\right]\hfill \\ \times {\int }_0^r{\xi }^2\times \frac{1}{E_Y}\times R \text{dR}.\hfill \end{array}$$ [1]

The surface energy can be expressed as:

$$\gamma =\left(E_0+kC\right)\times a/10.$$ [2]

The specific meaning of each symbol is detailed in SI Appendix. Following the minimum energy principle, the particle will crack when the strain energy exceeds twice the surface energy (fracture energy). The relationship between Li-ion concentration and particle radius was brought into the above two formulas. As shown in Fig. 1G, both strain energy and fracture energy increase with particle size, respectively, and the two curves intersect at R = 3.7 µm. The cracking free energy of NCM83 under different particle sizes was further calculated (Fig. 1H). The corresponding value approaches zero, indicating that the maximum particle size of NCM83 is 3.7 µm without cracking under different SOC, which will guarantee a high de-/lithiation reversibility for achieving a long cycle life.

Theoretical Guidance and Microstructure Characterizations.

To realize the synthesis of the target-sized single-crystalline NCM83, the Ostwald ripening formula:

$$R_t^3=R_0^3+kt,$$ [3]

was employed, where R_t denotes the final particle size, R₀ signifies the initial particle size, and k represents the particle growth rate coefficient (35, 36). Here, R₀ value can be defined as the particle size after the precursor is calcinated at 800 °C for 10 h because of its well-layered structure under relatively lower lithiation temperature (SI Appendix, Fig. S2). Based on this calcination condition, the temperature is further increased to 850 °C, 900 °C, 950 °C, and 1,000 °C, respectively, for another 5 h to obtain the corresponding R_t. According to the equation:

$$\text{ln}k=\text{ln}{k}_0+\left(-\frac{Q_B}{R}\right)\times \frac{1}{T},$$ [4]

the linear relationship between lnk and 1/T is obtained (SI Appendix, Fig. S3). After substituting it into Eq. 3, the correlational curves were achieved between the particle size (R_t) of NCM83 and the calcination temperature (T) and time (t), as shown in Fig. 2A. We find that the particles rapidly grow at initial stage and then gradually level off. Meantime, the higher the calcination temperature, the faster the initial growth rate of particle size (SI Appendix, Fig. S4). Although a calcination temperature of 1,000 °C for 1 min can reach the particle size of 3.7 µm, the severe adhesion and agglomeration between particles can be observed (SI Appendix, Fig. S5). Therefore, we chose a calcination temperature of 975 °C for 10 min to reach the desired particle size of 3.7 µm and then 800 °C for 10 h to round off the layered structure. The latter stage has little effect on particle size (SI Appendix, Fig. S4). Fig. 2B gives the SEM (scanning electron microscope) image of the as-obtained m-NCM83 with a size of about 3.68 µm and a span value of 0.71 (SI Appendix, Fig. S6). The EBSD (electron backscattered diffraction) image in Fig. 2C reveals a single-crystalline phase arrangement for each m-NCM83 particle, indicating the successful preparation of single-crystalline m-NCM83 cathodes. We also synthesized the b-NCM83 and s-NCM83 with mean particle sizes of 7.04 µm and 2.53 µm as controls (SI Appendix, Fig. S7).

Fig. 2.

Theoretical guidance and microstructure characterizations. (A) The NCM83 particle size variation over time at different lithiation temperatures. (B) SEM images and (C) EBSD orientation (Euler angle) mappings of m-NCM83. (D) XRD refined spectra, (E) R values, intensity ratios of (003) and (104) peaks, (F) crystal structure parameters, Li/Ni mixing values obtained from the refined XRD results, (G) EPR results, and (H) O 1s XPS spectra of s-NCM83, m-NCM83, and b-NCM83. (I) The O/Ni ratio from EDS line scanning signal intensities for m-NCM83 and b-NCM83.

The XRD (X-ray diffraction) patterns of three samples exhibit a typical layered structure without impurities, as shown in Fig. 2D. Among them, the m-NCM83 shows a highest (003)/(104) peak intensity and a lowest R value (Fig. 2E). It also gives a largest c/3a value (Fig. 2F), indicating a more orderly layered structure (37). Furthermore, the Li/Ni disorder of m-NCM83 is only 2.16%, much lower than b-NCM83 (5.38%) and s-NCM83 (2.29%). Fig. 2G shows the EPR results of the three samples, and the quantitative assess was also performed by integrating EPR curves (SI Appendix, Fig. S8). The results indicate that the oxygen vacancy concentration of m-NCM83 decreases by 67% compared to b-NCM83. Similar trends are also observed in the O 1s XPS spectra that m-NCM83 has a slightly stronger metal-oxygen (TM-O) bond peak with approximate oxygen vacancy content compared with s-NCM83, but much better than b-NCM83 (Fig. 2H). Meanwhile, the TEM-EDS (transmission electron microscopy-energy dispersive spectrometry) analysis further manifests that the O/Ni signal of m-NCM83 remains unchanged at the depth of 600 nm, while a clear and continuous decline for b-NCM83 on 100 nm surface depth, as shown in Fig. 2I. These highly consistent results indicate the introduction of short-time high temperature for m-NCM83 has no major impact on its ordered layered structure and oxygen vacancy but can gain the target-sized particles.

Electrochemical Performance and Cycling Stability Forecast in Coin-Type Half Cells.

The coin-type half-cell was selected to evaluate the electrochemical performances of m-NCM83, b-NCM83, and s-NCM83. Fig. 3A offers the first charge and discharge curves of the three samples at 25 °C. We observe that the charge voltage platform decreases with the reduction of particle size, implying the gradually lowered electrochemical polarization (38). Therefore, the m-NCM83 exhibits a highest Coulombic efficiency of 87.7% with a smallest irreversible capacity of 32.5 mAh g⁻¹. In Fig. 3B, the m-NCM83 shows a slightly lower reversible specific capacity at 0.2C (190.3 mAh g⁻¹) compared to s-NCM83, but it possesses the highest specific capacity of 118.6 mAh g⁻¹ at 10C, much higher than that of b-NCM83 (78.0 mAh g⁻¹) and s-NCM83 (107.5 mAh g⁻¹). The linear relationships between peak current and square root of the scanning rate in various CV curves indicates that the m-NCM83 shows the larger slope with the faster Li⁺ transfer kinetics than b-NCM83 (SI Appendix, Fig. S9). The Li-ion diffusion coefficient (D_Li⁺) was further quantified according to the galvanostatic intermittent titration technique (GITT) curves (SI Appendix, Fig. S10). As shown in Fig. 3C, both m-NCM83 and s-NCM83 display almost the same D_Li⁺ before rate performance test. After that, slight decrease has been seen for m-NCM83, but a sharp decrease for s-NCM83. This may be because the large specific surface area of s-NCM83 causes more side reactions with the electrolyte and the formation of rock-salt phase on the surface. The stability of the three samples was performed by capacity retention under different operation conditions. As shown in Fig. 3D, m-NCM83 exhibits a capacity retention of 91.6% after 100 cycles at 0.5C within 2.7 to 4.3 V, while only 74.5% and 83.6% for s-NCM83 and b-NCM83. Even operating at 55 °C or within 2.7 to 4.6 V, the m-NCM83 still exhibits high-capacity retentions of 89.6% and 90.5%, respectively (Fig. 3 E and F), remarkably better than the other two samples. For the m-NCM83, a 90.7% of initial capacity is still maintained after 100 cycles at 3C (SI Appendix, Fig. S11), which is much higher than the other two samples. The long cycle life of m-NCM83 was predicted by the capacity decay model by Jeff et al. (39), which can be described as:

$$\begin{array}{c}{Q}_d\left(t\right)=Q_0\left(1-A\sqrt{t}\right)-\frac{1}{2}\left({\left.\frac{\mathit{dQ}}{\mathit{dV}}\right|}_L+{\left.\frac{\mathit{dQ}}{\mathit{dV}}\right|}_U\right)\hfill \\ \ast \left(\mathrm{\Delta }V\left(t\right)-\mathrm{\Delta }{V}_0\right).\hfill \end{array}$$ [5]

Fig. 3.

Electrochemical performance and cycling stability forecast in coin-type half cells. (A) The initial charge–discharge curves, and (B) reversible specific capacity at 0.2C to 10C of m-NCM83, b-NCM83, and s-NCM83. (C) Li-ion diffusion coefficients calculated by GITT data before and after the rate test of m-NCM83 and s-NCM83. Capacity retention at 0.5C after 100 cycles at (D) 25 °C, (E) 55 °C within 2.7 to 4.3 V, and (F) at 25 °C within 2.7 to 4.6 V for the three samples. (G) Linear fitting of ΔV versus time, (H) the first dQ/dV versus V curve, and (I) the fitting curve of capacity decay versus time when cycling at 0.5C within 2.7 to 4.3 V for m-NCM83.

The $Q_0\left(1-A\sqrt{t}\right)$ accounts for capacity loss due to the consumption of active lithium for the formation of solid electrolyte interface membrane. The rest component represents capacity loss caused by the ever-increasing polarization. The detailed parameter descriptions are provided in SI Appendix. Base on this model, the data of m-NCM83 in 100 cycles (Fig. 3D) are brought into the Eq. 5. After fitting, the relationship (ΔV(t)) of the average charge–discharge voltage difference (ΔV) and the cycle time (t) was built in Fig. 3G. Meantime, the values of ${\left.\frac{\mathrm{d}\mathrm{Q}}{\mathrm{d}\mathrm{V}}\right|}_{\mathrm{L}}$ and ${\left.\frac{\mathrm{d}\mathrm{Q}}{\mathrm{d}\mathrm{V}}\right|}_{\mathrm{U}}$ were obtained by substituting the upper and lower cut-off voltages into the first dQ/dV versus V curve in Fig. 3H. The above data were substituted into the Eq. 5, the real-time capacity Qd(t) over the cycle time (t) for m-NCM83 can be expressed as Q_d(t) = 179.4*(1-0.02258*t^−1/2)- 0.03649*t. As shown in Fig. 3I, we can then predict the cycle life of m-NCM83 that will maintain 80.25% capacity after operating for 1,000 h.

Microstructure Analysis and Interface Stability after Cycling.

In order to unveil the enhancement mechanism of cycling stability, we delved into the microstructure and interface chemistry of m-NCM83 and b-NCM83 after 100 cycles at 0.5C. The SEM images of the cycled samples are provided in SI Appendix, Fig. S12 A and B. Contrary to b-NCM83, there is no observable crack in m-NCM83 with slight planar gliding for m-NCM83. The cross-sectional SEM images of m-NCM83 are also provided in Fig. 4 A and B, showing crack-free single-crystalline particles. An ordered lattice fringe with an interval of 0.47 nm is observed in Fig. 4C. The corresponding geometrical phase analysis (GPA) displays a minor internal lattice strain (Fig. 4D), implying a highly ordered layered structure of m-NCM83. However, there are obvious cracks in the interior of some particles of b-NCM83 with fresh exposure surface (Fig. 4 E and F). Meantime, we can see distinct rock-salt phase and lattice distortion (Fig. 4G) as well as the large strain inside the particles (Fig. 4H). The XRD patterns after cycling were displayed in SI Appendix, Fig. S12 C–F. The (003) peak shift in m-NCM83 is only 0.06° with slightly increased splitting gaps of (006)/(012) and (018)/(110) peaks, which is more obvious for the other two samples. The result further verifies the m-NCM83 has superior structural stability. These characterization results are highly consistent. The XPS analysis was performed to compare the surface chemistry of m-NCM83 and b-NCM83 after cycling. The peak intensities of Li_xPO_yF_z and LiF/NiF₂ greatly decrease for m-NCM83 (Fig. 4I), indicating less interface side reactions that result in P-O-F, C-O, and C=O peaks are greatly reduced compared with b-NCM83 (SI Appendix, Fig. S13). Meantime, the m-NCM83 exhibits lower Ni, Co, and Mn dissolution on Li anode than b-NCM83 (Fig. 4J). The in situ DEMS (differential electrochemical mass spectrometry) analysis was also carried out to monitor the gas generation in the first charge process. As shown in Fig. 4K, only traces of CO₂ and no O₂ are observed for m-NCM83 but obvious signals are detected when charging to 4.18 V for b-NCM83, thus significantly reducing the lattice oxygen loss with high safety. The alleviated gas production of m-NCM83 is mainly attributed to the fewer oxygen vacancies and the inhibited intragranular cracking. The former can hinder lattice oxygen escape by decreasing the lattice gliding, while the latter can effectively mitigate interfacial side reactions. Combined with the microstructure analysis, the big-sized particles have more serious side effects mainly owing to the rock-salt phase formation, more lattice oxygen loss and intragranular microcrack generation. Therefore, the particle size optimization is crucial for improving electrochemical performances and prolonging cycle life.

Fig. 4.

Microstructure analysis and interface stability after cycling. (A) Low and (B) high-magnification cross-section SEM images, (C) HRTEM images and (D) the corresponding strain state obtained by GPA for m-NCM83 after 100 cycles at 0.5C, and (E–H) the results for b-NCM83 at the same conditions. (I) F 1s XPS spectra, (J) the mass of transition metal (TM) deposited on Li anode after 100 cycles at 0.5C, and (K) in situ DEMS curves during the first charge process at 0.1C for m-NCM83 and b-NCM83, respectively.

Modifications and Electrochemical Performances in Coin-/Pouch-Type Full-Cells.

The K and Ti co-doped m-NCM83 was carried out to deeply investigate the application performances, in which K and Ti elements are located in the lithium layer and transition metal layer (40, 41), respectively. The dual-doping is expected to widen the Li-ion transport channel and suppress the Li/Ni disorder (Fig. 5A), thereby enhancing the rate and cycle performances. The average particle size of KT-m-NCM83 is 3.64 µm, evidencing that the K and Ti co-doping has no effect on particle size (SI Appendix, Fig. S14). Elemental maps of SEM (SI Appendix, Fig. S15) and TEM images (Fig. 5B) demonstrate the uniform distribution of K and Ti in m-NCM83 without local enrichment. The elemental compositions of KT-m-NCM83 have also been proven by ICP analysis (SI Appendix, Table S2), which is consistent well with designed components. Refined XRD result of KT-m-NCM83 verifies the well-layered structure without impurity phases, as shown in Fig. 5C and SI Appendix, Table S3. The KT-b-NCM83 and KT-s-NCM83 are also prepared for comparisons. In coin-type half-cells, the KT-m-NCM83 delivers a high specific capacity of 122 mAh g⁻¹ at 10C without large fluctuations in comparison with m-NCM83, also better than the other two samples (Fig. 5D). After 100 cycles at 0.5C, the KT-m-NCM83 exhibits an excellent capacity retention of 94.6% (Fig. 5E), much higher than KT-b-NCM83 (89.9%) and KT-s-NCM83 (83.8%). Nyquist plots of electrochemical impedances at different cycles are given in SI Appendix, Fig. S16. The corresponding surface film (R_sf) and charge transfer (R_ct) resistances are illustrated in Fig. 5F. During 100 cycles, both R_sf and R_ct of KT-m-NCM83 display a low initial value with a slight increase, but the other two samples have the opposite. These observations imply the superior interfacial and structural stability of KT-m-NCM83, which are in good agreement with the previous descriptions. The pouch-type full-cells are further assembled using KT-m-NCM83 as cathode and commercial graphite as anode. As shown in Fig. 5G, the capacity retention reaches 76.2% when operating from 0.1C to 3C. In particular, the KT-m-NCM83 exhibits superior cycle life with 88.1% retention of the first reversible capacity at 0.5C after 1,000 cycles, as shown in Fig. 5H. The Coulombic efficiency is always maintained at above 99.5%. The impressive merits indicate that the m-NCM83 has huge potentials for practical applications.

Fig. 5.

Modified cathodes and corresponding electrochemical performance. (A) Schematic of the crystal structure of K and Ti co-doped m-NCM83 (KT-m-NCM83). (B) STEM-EDS maps and (C) Refined XRD results of KT-m-NCM83. (D) The rate capability, (E) cycle performance, and (F) impedance changes of KT-s-NCM83, KT-m-NCM83 and KT-b-NCM83. (G) The rate capability and (H) the cycle stability of KT-m-NCM83 at 0.5C in pouch-type full-cells.

Conclusions

In summary, we exploit a unique strategy for the controllable synthesis of micro-sized single-crystalline Ni-rich oxide cathodes with highly ordered layered structure by introducing high-temperature short-duration pulse into the lithiation process. This is guided by the intrinsic correlations between particle size and lithiation time under various calcination temperature, which is a fitting of experimental data based on the Ostwald ripening formula. Then, taking NCM83 as an example, the stress distributions at various states of charge and their relationship with fracture free-energy are obtained by a diffusion-induced-stress simulation according to Fick’s second law and current distribution theory. An optimal particle size of 3.7 µm can be predicted and successfully has been synthesized, which vastly decreases strain energy, Li/O loss, and cationic mixing compared with big-/small-sized counterparts. The microcracks, planar gliding, and surface degradation are thereby greatly inhibited even at high state of charge, exhibiting a remarkably prolonged cycle life even operating at 55 °C or within 2.7 to 4.6 V. After modifying by dual-doping of K and Ti, the application performances have been further strengthened with 88.1% capacity retention at 0.5C after 1,000 cycles in pouch-type full cells. This work provides a crucial guidance for the design and synthesis of high-energy density single-crystalline Ni-rich oxide cathodes with superior cycling stability.

Materials and Methods

Materials Preparations.

The Ni_0.83Co_0.12Mn_0.05(OH)₂ (TM(OH)₂) precursor was bought from Haianzhichuan Co. Ltd. TM(OH)₂ was mixed with LiOH·H₂O (5% excess) and then sintered at 800 °C for 10 h with short-duration pulses at 975 °C for several minutes under a pure O₂ atmosphere to obtain the m-NCM83. For contrast, the cathodes with relatively smaller and bigger particles sizes were prepared by continuously calcining at 850 °C (s-NCM83) and 900 °C (b-NCM83) for 12 h, respectively. The K and Ti dual-doped m-NCM83 was also prepared by solid-phase calcination process. Typically, the TM(OH)₂, LiOH·H₂O (5% excess), and potassium titanium oxalate were first mixed for 10 min. Then, the KT-m-NCM83, KT-b-NCM83, and KT-s-NCM83 were separately obtained at different calcination conditions, which were consistent with m-NCM83, b-NCM83, and s-NCM83, respectively.

Material Characterizations.

The elemental compositions of the powders were identified by an inductively coupled plasma atomic emission spectrometer (ICP-AES, Agilent 725). The crystalline structure of all samples was confirmed by Powder XRD (Bruker D8 Advance, Cu Kα radiation) with a scan rate of 2° min⁻¹. The detailed structural information was analyzed via the General Structure Analysis System Rietveld refinement software. The morphology, microstructure, and element distribution of the samples were analyzed by a SEM (Helios G4 UC), and the TEM (FEI Talos F200X) with attachment of EDS. The cross-section and TEM samples were prepared by a triple ion beam milling system (Leica EM TIC 3X) and focused ion beam etching technique (FIB, TESCAN GALA 3), respectively. The crystal orientation of as-prepared samples was also measured via EBSD in transmission mode (t-EBSD). The particles size distributions of cathodes were measured by laser scattering particle size analyzer (Microtrac S3500SI) and nanoparticle size and zeta potential analyzer (Malvern Zetasizer Nano ZS90). D₅₀ refers to the particle size at which the cumulative distribution of a sample reaches 50% and is commonly used to represent the average particle size of a powder. The surface element compositions and valence were characterized via X-ray photoelectron spectra (XPS, AXIS Ultra DLD spectrometer). In situ DEMS experiments were performed on Swagelok cells with DEMS-QAS 100 (Shanghai Linglu Instruments Co. Ltd.) and Gamry instruments (Interface 1010E).

Electrochemical Measurements.

The prepared powders were mixed with carbon black (Super-P) and poly (vinylidene fluoride) in a mass proportion of 8:1:1 in N-methyl-2-pyrrolidone solvent. The resulting slurry was then coated on aluminum foil and dried in a vacuum environment with an active material loading of 3.0 to 4.0 mg cm⁻². The coin-type half cells were assembled in an argon-filled glove box with as-obtained electrode, lithium metal, and polypropylene separator (Celgard 2400). The electrolyte consisted of 1.2 M LiPF6 in a mixture of ethylene carbonate/ethyl methyl carbonate (VEC:VEMC = 3:7) with 2 wt.% vinylene carbonates. The galvanostatic charge and discharge measurements were performed by the LANDCT2001A battery test system at 25 °C (1C = 180 mAh g⁻¹). The commercial graphite (Shanshan Technology) was used as anode materials for making pouch-type full cells. The loading mass of cathode material was about 13.2 mg cm⁻¹ with a N/P ratio in the range of 1.05 to 1.10. The full cells were evaluated between 2.7 and 4.3 V at 25 °C under various current densities. Electrochemical impedance spectra experiment was performed using Autolab PGSTAT302N electrochemical workstation with the range from 100 kHz to 0.01 Hz.

Supplementary Material

Appendix 01 (PDF)

Data, Materials, and Software Availability

All study data are included in the article and/or SI Appendix.