English

• The urgent demand for power batteries and large-scale energy storage has imposed challenging requirements on the fast-charging capability, long cycle life, and cost of cathode materials used lithium-ion battery [1-3]. Among the current state-of-the-art commercial cathodes, spinel-type LiMn₂O₄ (LMO) holds a substantial market share due to its great structural stability and low cost [4, 5]. Despite these advantages, LMO suffers from severe Metal dissolution, structural degradation, and capacity fading during prolonged cycling, particularly at elevated temperatures [6, 7]. Moreover, at high charge/discharge rates, LMO exhibits apparent polarization and a rapid decline in electrochemical performance, limiting its widespread commercial application [8, 9].

  Extensive research has focused on understanding the origins of the capacity fading and increased polarization of LMO. A clear consensus has been reached that its rapid capacity decay is predominantly rooted in surface side reactions such as Mn dissolution [10]. In addition, high polarization during fast charging is mainly attributed to blocked or excessive Li⁺ diffusion pathways within the particles [11]. Moreover, repeated Li⁺ insertion/extraction over extended periods induce substantial stress from the anisotropic expansion and contraction of materials, resulting in intergranular cracks [12]. These cracks directly impede Li⁺ migration pathways and permit organic electrolytes to infiltrate the particle interior, causing severe side reactions and overall structural degradation [13].

  To overcome these challenges, developing single-crystal LMO has garnered significant attention in both academic and industrial research. Compared to polycrystalline LMO, single-crystal LMO exhibits superior electrochemical performance due to its coherent and stable crystal structure, appropriate particle size, and reduced specific surface area [14]. Furthermore, single-crystal LMO exhibits higher tap density, enhanced mechanical properties, and improved processability [15].

  Currently, various methods are employed to synthesize single-crystal LMO, including high-temperature solid-state calcination, Li salt flux methods, and hydrothermal synthesis. Specifically, the high-temperature calcination method (> 900 ℃) is advantageous for industrial production but results in excessively large single-crystal LMO particles (> 8 µm) [16]. Such large particle sizes hinder lithium-ion migration kinetics, leading to considerable polarization and capacity decline. Moreover, numerous oxygen vacancies and other crystal defects can be created on the particle surface under elevated temperature, adversely affecting electrochemical reactions. The Li salt flux method often involves adding large amounts of Li₂CO₃, LiCl, LiNO₃, and other lithium salts (with a Li/TM ratio greater than 0.6) during sintering to obtain single-crystal particles [17]. However, this method can produce particles with unstable crystal structures, such as non-spinel phases, and usually necessitates a washing step that dissolves Mn from the LMO surface. The hydrothermal method allows for precise control over the synthesis of small-sized, uniquely shaped Mn₃O₄, Mn(OH)₂, and other precursors, with the resulting single-crystal LMO retaining the morphology of these precursors [18]. Nonetheless, this method is complex, inconsistent, and challenging for large-scale production, rendering it unsuitable for commercial applications.

  Therefore, it is imperative to develop a method for synthesizing single-crystal LMO that integrates remarkable electrochemical and physical properties, cost-effectiveness, and scalability. This paper reports a strategy combing high-temperature pretreatment of precursors with BaO fluxing to synthesize moderately sized, densely surfaced single-crystal particles at conventional temperatures. These particles exhibit high crystallinity, structural stability, and a mall specific surface area, which effectively address stress accumulation and structural collapse caused by prolonged Li⁺ insertion /extraction, while mitigating Mn dissolution to preserve the integrity of the surface Li⁺ storage structure. Besides, the well-defined, wide, and short lithium-ion transport pathways confer rapid Li⁺migration capability and minimal polarization. Consequently, BaO-fluxed single-crystal LMO demonstrates exceptional cycling and rate performance. After 500 and 300 cycles at 1 C rate at room and high temperatures, the discharge specific capacity remains at 97.7 and 100.2 mAh/g, respectively. Moreover, at ultra-high discharge rates of 10 C and 20 C, the capacity retention rates are 91.4% and 88.8%, respectively.

  A cost-effective, scalable two-step method for synthesizing single-crystal LiMn₂O₄ (SC-LMOB) using BaO flux was developed guided by theoretical insights. As illustrated in Fig. 1, polycrystalline Mn₃O₄ (MO) secondary particles, composed of aggregated small primary particles, were sintered at 900 ℃ to form large aggregated Mn₃O₄ particles, denoted as S-MO. Subsequently, S-MO was uniformly mixed with Li₂CO₃ and calcined, enabling the primary S-MO particles to react and grow evenly, resulting in polycrystalline LiMn₂O₄ (PC-LMO). During this process, BaO flux was introduced, which facilitated the fusion and growth between adjacent grains at the grain boundaries of the primary S-MO particles, eliminating grain boundaries and ultimately forming single-crystal LiMn₂O₄ with surface-doped Ba²⁺, referred to as SC-LMOB.

  Figure 1

  

  Figure 1.  Schematic diagram of the fabrication of S-MO, SC-LMOB, and PC-LMO. The red arrows and wavy lines indicate the role of BaO flux in promoting crystal fusion and growth. The gray and light green spheres represent the Li source and Ba source, respectively.

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  To investigate the effects of high-temperature sintering on the morphology and phase composition of the MO precursor, as well as the impact of BaO flux on the synthesized LMO, we characterized MO, S-MO, PC-LMO, and SC-LMOB using field emission scanning electron microscopy (FESEM) and X-ray diffraction (XRD) (Fig. 2). Fig. 2a indicates that the MO precursor consists of secondary particles with diameters of 1–5 µm, formed by the aggregation of rough, loose, nano-sized primary particles. After high-temperature treatment, the S-MO precursor transforms into secondary particles with diameters of 3–10 µm, composed of smooth, dense primary particles with diameters of 1–3 µm (Fig. 2b). Therefore, high-temperature sintering markable promotes the fusion and growth of the original MO primary particles. Upon mixing S-MO with Li₂CO₃ and calcining, PC-LMO inherits the morphology of S-MO, forming aggregates of dense, irregular polyhedral particles with sizes ranging from 5 µm to 20 µm (Fig. 2c). In contrast, SC-LMOB exhibits dispersed, single-crystal brick-like morphology particles with sizes ranging from 0.5 µm to 5 µm (Fig. 2d), attributed to the fluxing effect of BaO at the grain boundaries of the primary particles. The large-radius Ba²⁺ facilitate the fusion growth of primary particles by overcoming grain boundary energy, resulting in dense single-crystal particles, a phenomenon also corroborated by other studies [19].

  Figure 2

  

  Figure 2.  Morphology and phase composition of the precursors and products. FESEM images of (a) the polycrystalline MO precursor, (b) sintered S-MO precursor, (c) PC-LMO product, and (d) SC-LMOB product. XRD patterns of (e) MO and S-MO precursors and (f) PC-LMO as well as SC-LMOB products, with enlarged views of the (111), (311), and (400) peaks.

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  Figs. 2e and f compare the crystal structures of MO before and after sintering, and the impact of flux on LMO products. The XRD patterns of MO and S-MO correspond to the orthorhombic Mn₃O₄ phase (PDF #24–0734). Compared to MO, S-MO displays nearly double the peak intensities, indicating improved crystallinity after high-temperature sintering (Fig. 2e) [20]. The position of the main (222) peak remains unchanged, suggesting minimal impact on unit cell parameters. However, trace amounts of Mn₂O₃ (PDF #24–0508) appear in S-MO, likely due to the partial decomposition of Mn₃O₄ at high temperatures [21]. The XRD patterns of PC-LMO and SC-LMOB match perfectly with spinel LMO (PDF #88–1026) without impurity peaks, indicating the conversion of Mn₃O₄ and trace Mn₂O₃ to LMO (Fig. 2f) [22]. The diffraction peaks of SC-LMOB are sharper and more intense than those of PC-LMO, signifying enhanced crystallinity [23]. Additionally, the enlarged views of the (111), (311), and (400) peaks show that the peaks of SC-LMOB are sharper and more symmetric, with a lower I₍₃₁₁₎/I₍₄₀₀₎ intensity ratio, suggesting that the BaO flux reduces cation mixing and structural defects in LMO [24]. To further investigate the impact of BaO on the lattice parameters of SC-LMOB, we conducted XRD refinement for both samples (Tables S1 and S2 in Supporting information). The results revealed that a small amount of Ba²⁺ was doped into the Li site of SC-LMOB, with a doping ratio of Li: Ba = 0.9994:0.0006. Additionally, the lattice parameter and volume of SC-LMOB were 8.2317 Å and 557.80 ų, which are larger than those of PC-LMO at 8.2308 Å and 557.61 ų, respectively. The results indicate that the slight doping of large-radius Ba²⁺ ions into the Li layer of SC-LMOB increases its cell volume, which is beneficial for enhancing the rapid insertion/extraction of Li⁺ within the crystal lattice [25].

  High resolution transmission electron microscopy (HRTEM) was used to observe the crystal structures of PC-LMO and SC-LMOB at microscopic level (Figs. 3a-h). Fig. 3a demonstrates that PC-LMO is composed of multiple irregularly shaped particles bonded together. At a higher magnification, atomic images can be observed in the nanometer-scale regions of the particle surface (Fig. 3b). Enlarging the area within the yellow dashed box and performing a fast Fourier transform (FFT) produces Figs. 3c and d. Precise measurements indicate that the interplanar spacings shown by the two sets of white dashed lines in Fig. 3c are d₁ = 0.290 nm and d₂ = 0.206 nm, corresponding to the (220) and (400) planes of LMO, respectively [26]. The diffraction spots in the FFT also confirm that the crystal system is Fd-3 m, and the crystal zone axis of this particle is [011] (Fig. 3d) [27]. In contrast, SC-LMOB exhibits a single-particle brick-like morphology (Fig. 3e). At higher magnification, clear lattice fringes with an interplanar spacing of d₃ = 0.206 nm, corresponding to the (400) plane of LMO, are observed on the particle surface (Fig. 3g) [9]. Further FFT analysis reveals the presence of the (422), indicating the crystal zone axis is [011] (Fig. 3h). Additionally, energy-dispersive X-ray spectroscopy (EDS) mapping presents uniform distribution of Mn, O, and Ba on the surface of the SC-LMOB particle (Fig. 3i).

  Figure 3

  

  Figure 3.  Crystal structure and element distribution. (a) TEM and (b) HRTEM image of fresh PC-LMO particle. (c) HRTEM at selected region and (d) corresponding FFT image of PC-LMO. (e) TEM and (f) HRTEM image of fresh PC-LMO particle. (g) HRTEM at selected region and (h) corresponding FFT image of PC-LMO. (i) FESEM and EDS mapping of SC-LMOB.

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  The electrochemical performance of samples with and without BaO flux at room temperature (25 ℃) were investigated to understand the impact of BaO on the Li⁺ storage and transport capabilities of SC-LMOB (Fig. 4 and Figs. S1-S4 in Supporting information). Fig. 4a examines the effect of BaO flux addition on the rate performance of the material. At a cutoff voltage of 4.3 V and 0.2 C (with 1 C = 148 mAh/g), the discharge capacities of SC-LMOB and PC-LMO are 112.9 and 105.4 mAh/g, respectively. It is evident that the discharge capacity of SC-LMOB consistently higher than that of PC-LMO. SC-LMOB delivers a reversible discharge capacity of 106.9 mAh/g from 0.2 C to 5.0 C (94.7% of the capacity retention), while PC-LMO drops to 84.6 mAh/g (79.1% of the capacity retention). Unprecedentedly, even at the ultra-high discharge rate, SC-LMOB still exhibits superior rate performance with discharge capacity of 103.2 and 100.3 mAh/g (with capacity retention of 91.4% and 88.8%) at 10 C and 20 C, respectively. In contrast, PC-LMO rapidly drops to only 72.3 and 40.7 mAh/g, with 68.6% and 38.6% capacity retention. Selected discharge profiles of PC-LMO and SC-LMOB at different rate are presented in Figs. 4b and c. The mid-point voltage of SC-LMOB retain above 3.86 V at 20 C, while PC-LMO decays to 3.27 V. Electrochemical impedance spectroscopy (EIS) tests (Fig. 4d) were carried out to understand the boost rate performance of SC-LMOB. The Nyquist plots of PC-LMO and SC-LMOB can be divided into two semi-circles in the high and medium frequency region, and a line in the low frequency region [28]. The equivalent circuit is shown in the inset, where R_e represents the solution resistance, R_f and CPE1 are the resistance and capacitance of solid electrolyte interface, R_ct and CPE₂ represent the charge transfer resistance and double layer capacitance, and W₁ represents the Li⁺ diffusion resistance of the cathode [29]. The fitting results (Fig. S1) reveal that both the R_f and R_ct of SC-LMOB are lower than those of PC-LMO, indicating faster surface Li⁺ insertion in SC-LMOB [30]. The diffusion coefficient of Li⁺ in bulk of cathode can be calculated by using Eq. 1 [31]:

  $ D=0.5 * R^2 T^2 A^{-2} n^{-4} F^{-4} C^{-2} \sigma^{-2} $

  (1)

  where D is Li⁺ diffusion coefficient, R is the gas contents (J K⁻¹ mol⁻¹), T is measurement temperature (K), A is the effect area (cm²) of cathode, n is the number of electrons per reaction species (n = 1 for LiMn₂O₄), F is Faraday constant, C is the concentration of Li⁺ (mol/cm³), and σ is Warburg factor, which obeys Eq. 2 [31]:

  $ Z^{\prime}=R_e+R_{c t}+\sigma \omega^{-1 / 2} $

  (2)

  Figure 4

  

  Figure 4.  Electrochemical characterization of PC-LMO and SC-LMOB within the voltage range of 3.0–4.3 V at 25 ℃. (a) Rate performance of PC-LMO and SC-LMOB at room temperature. Discharge curves of (b) PC-LMO and (c) SC-LMOB at 0.2, 2, 5, 10, and 20 C, respectively. (d) Electrochemical impedance spectroscopy (EIS) and the corresponding equivalent circuit diagram for the two samples. dQ/dV curves of (e) PC-LMO and (f) SC-LMOB at 0.2 C for the 1^st cycle and at 1 C for the 100^th, 200^th, 300^th, 400^th, and 500^th cycles, respectively. (g) Cycling performance of both samples at 1 C.

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  As shown in Fig. S2, Z' and ω ^−1/2 are plotted to obtain the value of σ. Based on the σ and Eq. 1, the diffusion coefficient of PC-LMO is calculated to be 3.66×10⁻¹¹ cm²/s, while that of the SC-LMOB is 4.34×10⁻¹¹ cm²/s. Therefore, the low surface resistance and the high internal Li-ion diffusion coefficient of SC-LMOB, attributed to the short and intact ion diffusion channels inside its small-sized single-crystal particles, as well as the large cell volume on its surface. Fig. S3 shows the electronic conductivity tests of the two samples at 25 ℃, where SC-LMOB exhibits an electronic conductivity of 2.55×10⁻⁵ S/cm, significantly higher than the 2.06×10⁻⁵ S/cm of PC-LMO. This may be due to the band bending at the BaO/LMO interface, causing electron accumulation on the surface, which can accelerate electron exchange reactions during Li⁺ insertion/extraction. Therefore, the superior rate performance of SC-LMOB compared to PC-LMO is due to the combined enhancement of both ionic and electronic conductivity.

  To further investigate the role of BaO flux in enhancing the room-temperature electrochemical and structural stability of single-crystal LMO, we examined the dQ/dV profiles, long-cycle performance and charge/discharge curves of PC-LMO and SC-LMOB at room temperature (25 ℃) (Figs. 4e-g and Fig. S4). Figs. 4e and f display the dQ/dV curves for the first activation cycle at 0.2 C and for the 100^th, 200^th, 300^th, 400^th, and 500^th cycles at 1 C for both samples. Both samples exhibit prominent peaks around 4.0 V and 4.1 V, linked to the Mn³⁺/Mn⁴⁺ redox couple and Li⁺ intercalation/de-intercalation processes [32]. For PC-LMO, the peak separations at 4.0 V and 4.1 V in the first cycle are 0.04 V and 0.03 V, respectively, whereas for SC-LMOB, they are 0.03 V and 0.02 V. This indicates that SC-LMOB experiences lower Li⁺ electrochemical polarization and better reversibility, attributed to its superior Li⁺ intercalation/de-intercalation kinetics [33]. With increasing cycle numbers, the redox peak intensities of PC-LMO decrease rapidly, indicating rapid structural degradation [34]. This is closely related to the phenomenon of its charge-discharge curve that its discharge plateau decreases quickly and becomes insignificant (Fig. S4). In contrast, the peak contractions for SC-LMOB are much slower, particularly the reduction peak at 4.0 V for Mn⁴⁺→ Mn³⁺, which maintains nearly constant peak area after 500 cycles at 1 C, resulting in the excellent structural stability of SC-LMOB.

  The long-cycle performance (Fig. 4g) further highlights the differences in electrochemical stability between the two samples. SC-LMOB exhibits a reversible discharge capacity of 97.7 mAh/g at 1 C after 500 cycles, with a capacity retention of 85.3%, notably higher than that of PC-LMO, which has a discharge capacity of 65.3 mAh/g and a capacity retention of 62.3%. The superior room-temperature electrochemical cycling performance of SC-LMOB is related to its appropriate particle size and small specific surface area, which reduce structural collapse and polarization-induced capacity fade due to surface side reactions.

  LMO typically suffers from rapid electrochemical degradation at high temperatures due to severe Mn dissolution [5]. Fig. 5 and Figs. S5 and S6 (Supporting information) exhibit the electrochemical performance of PC-LMO and SC-LMOB at 55 ℃. Therefore, it is crucial to improve the high-temperature electrochemical stability of the material. Fig. S5 illustrates the rate performance of the two samples at 55 ℃. The initial discharge capacities of PC-LMO and SC-LMOB are 108.4 and 115.5 mAh/g at 0.2 C and 55 ℃, respectively. When the discharge rate improved to 5 C, discharge capacity of PC-LMO decreased to 85.1 mAh/g, with a capacity retention of 78.5%. In contrast, SC-LMOB maintained a higher discharge capacity of 109.5 mAh/g at 5 C, corresponding to a capacity retention of 94.8%. The results indicate that, without the BaO flux, PC-LMO suffers a more rapid rate capability decline in high temperatures than room temperatures. Moreover, Figs. 5a and b display the charge-discharge curves for the first activation cycle at 0.2 C and for the 100^th, 200^th, and 300^th cycles at 1 C for both samples within the voltage range of 3.0–4.3 V. SC-LMOB has a higher initial discharge capacity (115.0 mAh/g) compared to PC-LMO (108.2 mAh/g). Additionally, SC-LMOB exhibits higher midpoint voltages during charging (4.08 V) and discharging (4.06 V) compared to PC-LMO (4.07 V and 4.03 V, respectively), indicating better grain size and lower polarization. After 300 cycles at 1 C, PC-LMO shows significant polarization, with its charge-discharge plateaus becoming completely blurred and midpoint voltages shifting to 4.15 V and 3.94 V, respectively. In contrast, SC-LMOB maintains distinct charge-discharge plateaus, with midpoint voltages of 4.01 V and 4.03 V, demonstrating its improved resistance to electrochemical degradation under severe high-temperature cycling conditions [35].

  Figure 5

  

  Figure 5.  Electrochemical characterization of PC-LMO and SC-LMOB within the voltage window of 3.0–4.3 V at 55 ℃. Charge/discharge curves of (a)PC-LMO and (b) SC-LMOB at 0.2 C for the 1^st cycle and at 1 C for the 100^th, 200^th and 300^th cycles, respectively. dQ/dV curves of (c) PC-LMO and (d) SC-LMOB at 0.2 C for the 1^st cycle and at 1 C for the 100^th, 200^th and 300^th cycles, respectively. (e) High temperature cycling performance of both samples at 1 C.

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  Figs. 5c and d show the dQ/dV curves for both electrodes at different cycle numbers. In the first cycle, the peak separations of the two redox pairs are similar for both samples. However, after 300 cycles at 1 C, the peak separations around 4.0 V and 4.1 V for PC-LMO increase to 0.06 V. Under the same conditions, peak separations of SC-LMOB change only to 0.04 V. Notably, the redox peak areas for PC-LMO shrink prominently more than those for SC-LMOB, indicating that PC-LMO without BaO flux undergoes rapid structural degradation at high temperatures. Fig. 5e compares the high-temperature long-cycle performance of the two electrodes. PC-LMO suffers severe electrochemical degradation, with its discharge capacity rapidly dropping to 53.8 mAh/g at 1 C after 300 cycles, corresponding to a capacity retention of 53.0%. In contrast, SC-LMOB maintains a discharge capacity of 100.2 mAh/g with a capacity retention of 86.0% under the same conditions. Furthermore, after 300 cycles, the discharge midpoint voltage of the PC-LMO drops rapidly from about 4.00 V to 3.94 V, while that of the SC-LMOB is maintained at about 4.05 V with only a slight decay (Fig. S6).

  These findings demonstrate that the superior high-temperature electrochemical performance of SC-LMOB is closely related to its small specific surface area and well-defined, wide Li⁺ diffusion pathways that effectively mitigates severe Mn dissolution and the Jahn-Teller effect during long high-temperature cycling, thereby maintaining structural integrity [36].

  The HRTEM images of the PC-LMO and SC-LMOB samples after 300 cycles at high temperatures are presented in Fig. S7 (Supporting information). As observed in Fig. S7a, PC-LMO exhibits localized cracks under these harsh conditions. Enlarging the crack area, marked by red circles, reveals relatively intact lattice fringes in the crystal interior far from the cracks. However, as the examination moves closer to the cracks, the lattice fringes become increasingly blurred (Fig. S7b). To further investigate, three regions along the white arrow were analyzed using FFT images, labeled b-Ⅰ, b-Ⅱ, and b-Ⅲ. The FFT results show a transition from a well-defined Fd-3 m spinel phase in the crystal interior to a defect-spinel phase, and eventually to a polycrystalline structure near the crack. This indicates that under high-temperature cycling, PC-LMO undergoes rapid structural degradation and crack formation. This process is exacerbated by Mn dissolution, leading to electrolyte infiltration, further structural degradation, and blocked Li⁺ diffusion pathways. In contrast, SC-LMOB maintains its brick-like morphology (Fig. S7c). Enlarged views and selected FFT images indicate that SC-LMOB retains a largely intact Fd-3 m spinel lattice structure from the bulk to the surface, with only minor surface lattice defects (Fig. S7d). This comprehensive analysis demonstrates that SC-LMOB exhibits significantly better structural stability at high temperatures compared to PC-LMO, which accounts for the substantial differences in their high-temperature electrochemical performance.

  In summary, we developed a synthesis technique combining high-temperature precursor treatment with a BaO flux to produce single-crystal LiMn₂O₄ cathodes. High-temperature sintering of the Mn₃O₄ precursor promotes particle growth, which benefits the morphology of the subsequent single-crystal product. The BaO flux remarkably facilitates the fusion and growth of the primary LiMn₂O₄ particles, resulting in single-crystal brick-like particles with moderate grain size, complete lattice structure, and high crystallinity. Compared to polycrystalline particles, these single-crystal particles have a smaller specific surface area, faster Li⁺ diffusion rate, and stable structure. These properties help mitigate issues such as Mn dissolution and excessive polarization, which lead to structural degradation under high-temperature and ultra-high charge/discharge rate conditions. The results indicate that the BaO flux synthesis of single-crystal materials is an effective strategy to address problems like surface Mn dissolution, structural collapse, and large polarization caused by rapid Li⁺ insertion/extraction in LiMn₂O₄. This approach is conducive to achieving satisfactory electrochemical performance and accelerating the widespread commercial application of single-crystal LiMn₂O₄.

  Declaration of competing interest

  The authors declare that they have no known competing financial interests or personal relationships that could have appeared to influence the work reported in this paper.

  CRediT authorship contribution statement

  Yuming Shu: Conceptualization, Data curation, Investigation, Methodology, Software, Writing – original draft, Writing – review & editing. Hanghang Lei: Data curation. Jiangnan Huang: Methodology. Qing Pan: Methodology. Baichao Zhang: Data curation. Yixin Xu: Methodology. Ye Zhou: Software. Guorong Hu: Resources. Yanbing Cao: Project administration. Guoqiang Zou: Resources. Wentao Deng: Project administration, Resources. Zhongdong Peng: Project administration, Resources, Supervision. Hongshuai Hou: Resources. Di Chen: Project administration, Resources. Xiaobo Ji: Funding acquisition, Project administration, Resources, Supervision.

  Acknowledgments

  This work is supported by National Key Research and Development Program of China (No. 2021YFB3502000) and the National Natural Science Foundation of China (Nos. 22309207, 52325405, U21A20284, 52261135632, 51874358 and 51772333).

  Supplementary materials

  Supplementary material associated with this article can be found, in the online version, at doi:10.1016/j.cclet.2024.110345.

  
  
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  Figure 1  Schematic diagram of the fabrication of S-MO, SC-LMOB, and PC-LMO. The red arrows and wavy lines indicate the role of BaO flux in promoting crystal fusion and growth. The gray and light green spheres represent the Li source and Ba source, respectively.

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  Figure 2  Morphology and phase composition of the precursors and products. FESEM images of (a) the polycrystalline MO precursor, (b) sintered S-MO precursor, (c) PC-LMO product, and (d) SC-LMOB product. XRD patterns of (e) MO and S-MO precursors and (f) PC-LMO as well as SC-LMOB products, with enlarged views of the (111), (311), and (400) peaks.

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  Figure 3  Crystal structure and element distribution. (a) TEM and (b) HRTEM image of fresh PC-LMO particle. (c) HRTEM at selected region and (d) corresponding FFT image of PC-LMO. (e) TEM and (f) HRTEM image of fresh PC-LMO particle. (g) HRTEM at selected region and (h) corresponding FFT image of PC-LMO. (i) FESEM and EDS mapping of SC-LMOB.

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  Figure 4  Electrochemical characterization of PC-LMO and SC-LMOB within the voltage range of 3.0–4.3 V at 25 ℃. (a) Rate performance of PC-LMO and SC-LMOB at room temperature. Discharge curves of (b) PC-LMO and (c) SC-LMOB at 0.2, 2, 5, 10, and 20 C, respectively. (d) Electrochemical impedance spectroscopy (EIS) and the corresponding equivalent circuit diagram for the two samples. dQ/dV curves of (e) PC-LMO and (f) SC-LMOB at 0.2 C for the 1^st cycle and at 1 C for the 100^th, 200^th, 300^th, 400^th, and 500^th cycles, respectively. (g) Cycling performance of both samples at 1 C.

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  Figure 5  Electrochemical characterization of PC-LMO and SC-LMOB within the voltage window of 3.0–4.3 V at 55 ℃. Charge/discharge curves of (a)PC-LMO and (b) SC-LMOB at 0.2 C for the 1^st cycle and at 1 C for the 100^th, 200^th and 300^th cycles, respectively. dQ/dV curves of (c) PC-LMO and (d) SC-LMOB at 0.2 C for the 1^st cycle and at 1 C for the 100^th, 200^th and 300^th cycles, respectively. (e) High temperature cycling performance of both samples at 1 C.

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