Material Preparation

Pristine Li_1.2Ni_0.13Co_0.13Mn_0.54O₂ material was purchased from Fuli. The AlPO₄-coated Li_1.2Ni_0.13Co_0.13Mn_0.54O₂ material was synthesized using a facile sol-gel method. Different concentrations of phosphorus pentoxide (P₂O₅, 99.99%, Aladdin) and aluminum nitrate hydrate (Al(NO₃)₃·9H₂O, analytical reagent, Sinopharm) were dissolved in ethanol. Ammonium bicarbonate (NH₄HCO₃, reagent grade, Aladdin) was used to adjust the pH to 7, and the solution was stirred at 60 ℃ for 3 h. Salting-out was carried out at -10 ℃ using liquid nitrogen, and the mother liquor was collected after filtration. Pristine Li_1.2Ni_0.13Co_0.13Mn_0.54O₂ material was added and stirred at 80 ℃ for 3 h to obtain the LMR@APO precursor. This precursor was sintered in a muffle furnace at 400 ℃ for 4 h to obtain the Li_1.2Ni_0.13Co_0.13Mn_0.54O₂ material coated with AlPO₄. Only the amounts of Al(NO₃)₃·9H₂O and P₂O₅ need to be changed to obtain Li_1.2Ni_0.13Co_0.13Mn_0.54O₂ materials with different AlPO₄ coating contents.

Material Characterizations

Scanning electron microscopy (SEM; S-4800N, Hitachi, Japan) and energy-dispersive X-ray spectroscopy (EDS) mapping was performed to observe the morphology of the samples. High-resolution transmission electron microscopy (HRTEM) was conducted using a transmission electron microscope (G2 F20 S-TWIN, Tecnai FEI, USA). X-ray diffraction (XRD) measurements were performed for phase analysis (D8 Advance X, Bruker, USA; Cu-Ka radiation, λ=1.542 Å) in the 2θ range of 10–80 ° at room temperature. X-ray photoelectron spectroscopy (XPS; ESCALAB250 Thermo Fisher, USA) was performed to determine the elemental composition of the samples. The ex-situ hXAS test for the Mn K-edge of the LMR sample was performed at the Shanghai Synchrotron Radiation Facility (SSRF) BL16U1 beamline, with a photon flux of 3.7×10¹² photons s^-1 at the sample position, using the transmission mode [20]. The hXAS test for the Mn K-edge of the long-cycling LMR samples was performed at the SSRF BL14W beamline with a photon flux of 2.6×10¹² photons s^-1 at the sample position using the transmission mode [23]. The ex-situ hXAS tests for the Ni K-edge and Co K-edge of the LMR sample were performed at the SSRF BL20U1 beamline with a photon flux of 10¹² photons s^-1 at the sample position using the fluorescence mode [24]. All the hXAS tests were calibrated using the K-edge data of the transition metal foils corresponding to the tested elements. The Mn L-edge sXAS of the LMR sample was obtained at the SSRF BL08U1A beamline using the TEY mode with photon fluxes of 1.5×10⁸ photons s^-1 and an approximate pressure of 10^-6 Torr [25, 26]. Calibration was performed using Mn L-edge data of the standard MnO₂ sample. The O K-edge sXAS was performed at the SSRF BL02B02 beamline using a combination of the TEY and TFY modes with a photon flux at the sample position of 3×10¹⁰ photons s^-1 and an approximate pressure of 5×10^-9 Torr [27]. Calibration was performed using the O K-edge data of the standard SrTiO₃ sample. The XAS data were processed using Athena software from the IFEFFIT software package for energy calibration, noise removal, background subtraction, and normalization [28]. Wavelet-transform analysis was performed using wtEXAFS software. The Mn K-edge TXM-XANES test of the LMR samples was carried out at the SSRF BL18B beamline, with a photon flux of 1.3×10¹⁰ photons s^-1 at the sample position, field of view of approximately 20 μm × 20 μm, and spatial resolution of 60 nm [29]. The active material from the electrode sheet was obtained from batteries disassembled in an argon-filled glove box. The material was scraped off with a scalpel and then placed in a blue-capped vial, and dimethyl carbonate (DMC) solvent was added. After sealing, the blue-capped vial was removed from the glove box for ultrasonication. Finally, before testing, the sample was delivered to the microtip of the glass tube using an industrial dispenser.

Electrode Preparation and Cell Assembly

Electrochemical-performance tests were conducted using coin-type 2025 cells. Both cathode materials were composed of a mixture containing 80 wt% active material and 10 wt% poly (vinylidene difluoride) binder in a precise amount of N-methyl-2-pyrrolidone (NMP), 5 wt% Super P, and 5 wt% vapor-grown carbon fiber. The resulting slurry was uniformly cast onto aluminum foil and dried overnight at 70 ℃ under vacuum conditions to remove the NMP. Then, disc electrodes with a diameter of 12 mm were punched from the coated foil, achieving an average active-material mass loading of 2.5–3.5 mg cm^-2. Half-cell assembly was performed in an argon-enriched glovebox to ensure a contamination-free environment, utilizing pure lithium as the counter electrode. A polypropylene membrane (Celgard-2400) was used as the separator. The electrolyte solution consisted of 1 M LiPF6 dissolved in a mixture of ethylene carbonate, DMC, and diethyl carbonate in a volumetric ratio of 1:1:1. Precisely 100 μL of this type of electrolyte was added to each coin cell to maintain a uniform electrochemical behavior.

Electrochemical Measurements

All measurements for the electrochemical testing of the coin-type cells were performed at ambient room temperature. Galvanostatic cycling tests were performed at a constant current density of 1 C (1 C = 250 mA hg^-1) over 2.0–4.6 V using a NEWARE battery-testing system.

Structure and Morphology of AlPO₄-coated LMR Material

The XRD patterns of pristine, 0.5 wt% AlPO₄-, 1 wt% AlPO₄-, and 2 wt% AlPO₄-coated LMR samples are shown in Fig. 1a and S3. The patterns can be indexed with the O3-type layered structure of the hexagonal a-NaFeO₂, featuring an $R\overline{3}m$ space group. The additional reflections between 20° and 25°, corresponding to a weak superlattice, are attributed to the (020) and (110) lattice planes of the Li₂MnO₃-like component with a C2/m space group. This implies the existence of a honeycomb-like short-range ordered structure of LiMn₆ in the transition-metal (TM) layers. The observed separation of the (006)/(012) and (018)/(110) peaks indicates a well-crystallized layered structure in the lattice [22]. Compared to the original sample, coated samples exhibit similar diffraction peaks, indicating no significant change in the bulk structure. The fine-scanned XRD patterns of samples with different coating amounts at positions corresponding to the AlPO₄ reflection peaks are displayed in Fig. S4. No obvious reflections are observed owing to the amorphous state and low coating amount of the AlPO₄ coating layer. This means that the sample coated with AlPO₄ maintains the intact hexagonal a-NaFeO₂ structure with good arrangement after surface modification. Additionally, XPS measurements are conducted to provide more evidence of the successful modification of AlPO₄. The photoemission peaks of Al 2P and P 2p detected on the LMR@0.5% APO sample are illustrated in Fig. 1b, c. The Al 2p binding energy, measured at 73.4 eV, is between the reported values for the AlPO₄ material at 74.5 eV and the range observed for materials such as LiAl_xM_1-xO₂ (M = Co, Ni)[30]. This may be attributed to the diffusion of Al after the calcination process, resulting in the formation of a solid solution at the interface of the coating layer and underlying bulk material [31]. The P 2p photoemission peak is divided into two components at 134.5 eV and 133.2 eV, which are consistent with the reports of AlPO₄ and Li₃PO₄, confirming the presence of the AlPO₄ phase and accompanying Li₃PO₄ phase on the sample surface [31]. SEM images displayed in Fig. 1d, e reveal that both pristine and 0.5 wt% AlPO₄-coated LMR materials exhibit comparable morphology, suggesting that the coating process does not alter the morphology of the particles. The spherical particles, approximately 10 micrometers in diameter, consist of uniform primary nanoparticles. Figure 1f and S5 demonstrate the even distribution of Al, P, and O in the EDS elemental-mapping images, demonstrating the successful encapsulation of AlPO₄.

Fig. 1Full-screen View

(Color online) (a) XRD patterns of pristine LMR and LMR@0.5% APO samples. (b) P 2p and (c) Al 2p XPS spectra of LMR@0.5% APO sample. SEM images of (d) pristine and (e) LMR@0.5% APO samples. (f) EDS element mapping of the LMR@0.5% APO sample. HRTEM images of (g) pristine LMR and (h) LMR@0.5% APO samples

TEM was also employed to further obverse the AlPO₄ coating layer. In Fig. 1g, h, TEM images of both the pristine and 0.5 wt% AlPO₄-coated LMR samples show well-defined equidistant fringes. This corresponds to a d-spacing of 0.47 nm on the (003) plane, indicating good crystallinity in bulk crystal[32]. The TEM image of the 0.5 wt% AlPO₄-coated LMR sample shows an approximately 8-nm-thick coating layer within an amorphous phase. Additionally, this coating layer is uniformly distributed across the surface of the bulk material.

Electrochemical Performance Measurements

The electrochemical performances of both the pristine and 0.5 wt% AlPO₄-coated LMR samples are shown in Fig. 2. Figure 2a compares the cycling performances of pristine and LMR@0.5% APO samples with different coating amounts at a 1 C rate. To fully activate the lithium-rich manganese material, a low rate of 0.1 C was employed for the first two cycles, with a working-voltage range set between 2.0 and 4.7 V. For subsequent cycles, a 1 C rate was employed with a working-voltage range set between 2.0 and 4.6 V. The pristine LMR sample has an initial discharge specific capacity of 293.43 mA h g^-1; however, it begins to decay sharply after approximately 100 cycles, and the capacity is only 39.12 mA h g^-1 after 200 cycles. The LMR@0.5% APO sample exhibits enhanced cycling stability and specific capacity during cycling, with an initial discharge specific capacity of 305.77 mA h g^-1. After 200 cycles at 1 C, the capacity remains at 173.36 mA h g^-1, with a capacity-retention rate of 73.9% relative to the third cycle at the beginning of the 1 C cycling. This is because the side reactions caused by the complete exposure of the pristine LMR-material surface to the electrolyte, the irreversible structural phase changes on the surface of the LMR material, and local structural distortion at high voltages during the cycling process are suppressed by the AlPO₄ coating layer. Figure 2b presents the initial charge-discharge curve of the pristine and 0.5 wt% AlPO₄-coated LMR samples at 0.1 C (1 C = 250 mA h g^-1) within the voltage range of 2.0–4.7 V. The initial charge curve is characterized by a sloping section below 4.4 V and plateau above 4.5 V. This lower sloping section is associated with the oxidation of Co³⁺ and Ni²⁺ ions. The plateau above 4.5 V corresponds to the phase transition involving the anionic redox reaction, which is typically accompanied by the liberation of oxygen [33, 34]. The pristine LMR sample exhibits an initial charging capacity of 318.45 mA h g^-1 at 0.1 C; however, it incurs a substantial irreversible capacity of 25.22 mA h g^-1. The specific initial discharge capacity is 293.43 mA h g^-1, with a Coulombic efficiency of 92.1% for the initial cycle. Conversely, the 0.5 wt% AlPO₄-coated LMR sample presents an initial discharge capacity of 305.77 mA h g^-1, along with an improved Coulombic efficiency of 92.6%. This enhancement suggests that the AlPO₄ coating effectively stabilizes the surface of the sample. It also mitigates the irreversible loss of lattice oxygen and curtails side reactions at the electrode–electrolyte interface, which contribute to capacity decline [35], thereby improving the initial discharge capacity and overall Coulombic efficiency. To further investigate the electrochemical reactions throughout the charge and discharge cycles, the corresponding differential capacity versus voltage (dQ/dV) curves and charge-discharge curves of both samples at different numbers of cycles are shown in Fig. 2c and S6. The LMR samples do not exhibit significant differences in the overall distribution of the peak profiles after the coating modification before approximately 100 cycles. The first peak at 4.05 V is attributed to the oxidation of Ni²⁺/Ni⁴⁺ and Co³⁺/Co⁴⁺, and the second peak at 4.5 V is attributed to the anionic oxidation in the initial charge process. Owing to the inversion of cationic and anionic redox reactions, the first peak at 4.3 V is attributed to the reduction of the transition metal, whereas the second peak at approximately 3.3 V is attributed to the reduction of anionic ligands and the reduction of Mn⁴⁺ to Mn³⁺ owing to the loss of O₂ [9] during the discharge process. The additional reduction peak at 2.75 V is due to lithium-ion intercalation of the spinel phase. This indicates the formation of a spinel phase on the surface during the coating process [36], which can be substantiated by subsequent sXAS analysis. The dQ/dV plots of the subsequent cycles indicate that after 100 cycles, the oxidation peak of the pristine LMR sample significantly shifts to the right and the reduction peak significantly shifts to the left, suggesting that substantial polarization affects the progression of redox reactions and leads to rapid capacity decrease. In contrast, the 0.5 wt% AlPO₄-coated LMR sample maintains good electrochemical activity with no significant shifts in the positions of the oxidation and reduction peaks, demonstrating the well-maintained structure because of the AlPO₄ coating.

Fig. 2Full-screen View

(Color online) (a) Cycle performance of pristine LMR and LMR@0.5% APO samples at 1 C. (b) Initial charge and discharge curves (at 0.1 C). (c)dQ/dV profiles (at 0.1 C). dQ/dV profiles from the 3rd to 200th cycle at 1 C of (d) pristine LMR and (e) LMR@0.5% APO samples

Electronic Structure Characterized by Combined Synchrotron-based X-ray Absorption Spectroscopy

To further explore the modification mechanism of AlPO₄ coatings on LMR materials, ex-situ synchrotron-based sXAS and hard XANES technology are employed to investigate the evolution of surface and bulk Mn in the LMR materials during the electrochemical cycling process. The hard-XANES test was performed at beamline BL20U of the SSRF using the transmission mode [24]. The transmission mode involves the ratio of the ionization-chamber currents before and after the sample, reflecting the interaction between the X-rays and material. Through hard-XANES testing, we can obtain information about the lowest unoccupied states of the 3d transition-metal elements in the bulk of the material and determine their average oxidation states [37-39]. The sXAS test was performed at beamline BL08U1A of the SSRF using the TEY mode. In the TEY mode, we can measure the photogenerated current on the material surface using a picoammeter and determine the transition of 3d transition-metal elements from the 2p states to 3d states. Because information on the unoccupied 3d orbital states is directly related to the electrochemical performance of cathode materials, sXAS can provide strong support for the study of cathode materials [40, 41]. As shown in Fig. 3b, during the charging process, the white-line peak of the Mn K-edge shifts marginally to the right-hand side, whereas the half-height energy moves to the left-hand side. Because of the complexity of the Mn K-edge spectrum, the oxidation state of Mn cannot be determined accurately. The position of the K absorption edge corresponds to the transition from the Mn 1s orbital to the first dipole-allowed transition orbital; therefore, a significant enhancement in the absorption intensity occurs at the edge position. Therefore, the absorption edge can be accurately determined by identifying the position of the strongest peak in the first derivative, thereby determining the oxidation state. During the charging process, the maximum value of the first derivative, which is the position of the absorption edge, undergoes a slight shift to the right-hand side of less than 2 eV, which may be owing to the oxidation of low-valence Mn produced by element segregation during the sample-synthesis process. When discharged to 2.75 V, the Mn K-edge XANES spectrum shifts significantly back towards the pristine position, corresponding to the reduction of Mn on the surface during discharge. Notably, in the parts of the charge–discharge curves where the capacity is not provided, specifically during charging from OCV to 3.7 V and discharging from 2.75 V to 2 V, no significant shift in the Mn K-edge XANES is observed, indicating that the average valence state does not significantly change during these processes. The ex-situ sXAS spectrum in Fig. 3c shows that when charging from OCV to 3.7 V, a marked decrease occurs in the valence state of surface Mn from approximately Mn⁴⁺ to Mn²⁺. This valence state remains at approximately +4 during the charging plateau above 4.4 V. Upon discharging to 2.75 V, features of Mn²⁺ and Mn⁴⁺ reemerge, which corresponds to the reduction of Mn⁴⁺ observed in the dQ/dV plots. Further discharge to 2.0 V results in a rapid reduction of the surface valence state to Mn²⁺. This evolutionary trend is clearly demonstrated in the differential spectrum shown in Fig. 3d. Integrating the analysis of differences in Mn signals between the bulk and surface indicates that the surface Mn²⁺ is primarily generated in the low-voltage sloped sections of the electrochemical curves, where no capacity is provided, as the average valence state of Mn in the bulk remains unchanged. This suggests that the significant surface enrichment of Mn²⁺ results from intrinsically chemical reactions, rather than electrochemical reactions. Furthermore, this observation aligns with findings in the literature that a weakly acidic LiPF₆ electrolyte promotes the disproportionation of Mn³⁺ to form Mn⁴⁺ and the more soluble Mn²⁺, which is more likely to be desorbed and adsorbed on the surface [21]. This further supports the hypothesis of chemically driven changes in Mn valence at the surface during electrochemical cycling. Subsequently, we tracked the evolution of surface Mn during extended cycling. As shown in Fig. 3e, the surface Mn oxidation state was consistently at +2 before the rapid capacity decline (within 100 cycles), and it then increased to +4 after 200 cycles. This is attributed to ongoing disproportionation reactions at the surface during electrochemical cycling, where Mn⁴⁺ ions are retained on the surface and Mn²⁺ ions are solubilized. This disrupts the electrolyte and leads to interference between the anode and cathode. This disproportionation reaction leads to increased polarization and renders the material electrochemically inactive, triggering a rapid capacity decrease. After 200 cycles, the surface Mn²⁺ is completely solubilized in the electrolyte, and the surface signal of the sample is dominated by Mn⁴⁺ [7]. The improvement mechanism of the AlPO₄-coating layer is then investigated. By utilizing ex-situ Mn L-edge XANES, the evolution of Mn on the surfaces of pristine LMR and LMR@0.5% APO samples are compared in Fig. 3f and g, which show distinct differences. First, the feature peak corresponding to Mn³⁺ in the coated sample before cycling is significantly stronger than that in the pristine sample, indicating the formation of a spinel-like phase on the surface owing to the coating process. During the charging process, the coated sample exhibits a slower rate of Mn oxidation. This suggests that the AlPO₄-coating layer can modulate the reaction degree, thereby suppressing the overcharge and surface structural changes related to high-voltage oxygen release. Furthermore, when charged to 4.7 V, a small Mn³⁺ signal appears on the surface of the pristine LMR sample, corresponding to the so-called reductive elimination mechanism, where surface transition metals are reduced owing to oxygen release at high charging states. The oxygen-deficient surface structure results from the irreversible release of oxygen, which causes a more severe reduction of Mn at low voltages. In contrast, the LMR@0.5% APO sample maintains Mn⁴⁺ on its surface when charged to 4.7 V, indicating that the AlPO₄-coating layer effectively stabilizes the surface structure, suppresses surface-oxygen evolution, and enhances the initial Coulombic efficiency [42]. Finally, when discharged to 2.75 V, no Mn²⁺ signal is detected, even though the coated-sample surface similarly presents the Mn³⁺ signal resulting from the reduction of Mn⁴⁺. Moreover, an Mn²⁺ signal can be clearly observed in the Mn L-edge sXAS of the cycled LMR@0.5% APO sample (Fig. S7), which is different from that of the cycled LMR sample. This implies that the AlPO₄-coating layer effectively obstructs the direct interaction between the electrode and electrolyte [10, 43]. This inhibits side-reaction-induced Mn disproportionation and significantly mitigates the capacity decay caused by Mn²⁺ dissolution, which enhances the cycle stability.

Fig. 3Full-screen View

(Color online) (a) Electrochemical curve of LMR material under ex-situ X-ray-absorption spectrum. (b) Ex-situ XANES spectra at Mn K-edge for pristine LMR sample during the first cycle. (c) Ex-situ XANES spectra at Mn L-edge for pristine LMR sample during the first cycle. (d) Differential spectrum of ex-situ XANES spectra at the Mn L-edge for pristine LMR sample during the first cycle. (e) XANES spectra at Mn L-edge for pristine LMR sample after 200 cycles. Ex-situ XANES spectra of the evolution of Mn for (f) pristine LMR and (g) LMR@0.5% APO samples

Extended X-ray-absorption fine structure (EXAFS) spectroscopy was used to investigate changes in the bulk average oxidation states and local coordination environments of pristine LMR and LMR@0.5% APO during long-term cycling. The EXAFS and XANES testing methods are similar, but the EXAFS test requires scanning over a wider energy range. Through EXAFS testing, we can simultaneously obtain information on the nearest-neighbor coordination and oxidation state of 3d transition-metal elements. This is important for cathode materials that undergo severe local structural distortion and changes in oxidation states during the long cycling process [44-46]. As shown in Fig. 4a, b, the Mn K-edge white-line peak of the pristine LMR sample shifts to the right-hand side, whereas that of the LMR@0.5% APO sample shifts to the left-hand side, indicating that the coating modification alters the evolution trend of the bulk oxidation state and local structure of the material during long cycling processes. The first derivative plot of XANES is also presented to obtain oxidation-state information more precisely through the position of the absorption edge. The Mn K-edge of the pristine LMR sample clearly shifts to the right-hand side, indicating an increase in the overall oxidation state of the sample during the long cycling process. However, the Mn K-edge of the LMR@0.5% APO sample exhibits a slight shift to the left-hand side, suggesting that its structure is well maintained during the long cycling process, with only minor electrochemical degradation occurring. However, compared with the obvious changes in the oxidation state of Mn on the surface, the change in the Mn K-edge in the bulk is less than 1 eV. Therefore, the Fourier-transform (FT) EXAFS (Fig. 4c, d) and wavelet transform analysis (Fig. 4e, f, g, h) of the pristine and LMR@0.5% APO samples are employed. The results show that the Mn-coordination structures in the bulk phase of both the coated and pristine samples demonstrate no significant differences in the bond lengths and coordination numbers corresponding to the Mn-O and Mn-TM shells after long-term cycling. This difference in the evolution of Mn in the bulk oxidation state and local structure indicates that the spatial uniformity of the Mn reaction during the extended cycling process is a key factor affecting cycling performance.

Fig. 4Full-screen View

(Color online) XANES spectra at Mn K-edge of pristine and cycled (a) LMR and (b) LMR@0.5% APO samples. FT-EXAFS of pristine and cycled (c) LMR and (d) LMR@0.5% APO samples. Wavelet transform of (e) pristine LMR and (f) LMR after 200 cycles; wavelet transform of (g) pristine LMR@0.5% APO and (h) LMR@0.5% APO after 200 cycles

Revealing the Reaction Homogeneity through Synchrotron-based Transmission X-ray Microscopy Technique Combined with XANES Spectroscopy

By employing synchrotron-based TXM-XANES testing at beamline BL18B of the SSRF, the spatial distribution of the Mn-related oxidation state was visualized at the nanoscale, allowing for direct observation of the evolution uniformity of the Mn-related oxidation state. This can be used to study the degradation mechanisms of LMR materials during the long cycling process as well as the modifying effects of the AlPO₄-coating layer on these materials. TXM-XANES is a synchrotron-based imaging method that integrates TXM with XANES. The spatial distribution of the oxidation states of micrometer-scale cathode materials can be characterized using X-rays with continuously varying energy for the transmission imaging of materials. This approach overcomes the limitations of traditional approaches, which may neglect detailed characterization by examining only specific areas or averaging values across the entire sample, potentially missing localized changes [29, 47]. The 2D TXM-XANES mapping offers a detailed aging-state analysis for each element by tracking the shifts in their respective XANES spectra [48]. Figure 5a, d show that the pristine LMR and LMR@0.5% APO samples exhibit a relatively uniform distribution of Mn oxidation states before cycling, with no local enrichment of Mn elements in a particular oxidation state. As shown in Fig. 5b, e, the oxidation-state distributions of the pristine LMR and LMR@0.5% APO samples evolved based on different trends after long-time cycling. Moreover, Fig. S8 illustrates that the average XANES curves extracted from all pixel points of the pristine LMR and LMR@0.5% APO samples are consistent with the XANES curves obtained by conventional transmission-mode XANES before and after long-term cycling. This indicates that the differences in the bulk oxidation states of LMR particles and enhancement of cycling stability by the AlPO₄-coating layer are closely related to the uniformity of the Mn reaction. After 200 cycles, the 2D oxidation-state mapping of Mn in the pristine LMR sample (Fig. 5b shows a decrease in the oxidation state at the center position corresponding to the bulk of the sample. A clear enrichment of high-valence Mn is observed at the edge positions corresponding to the sample surface, which leads to a slight increase in the average oxidation state. This result is consistent with the difference between the bulk and surface regions of the samples obtained from hard and soft X-ray absorption spectroscopy. This indicates that during the extended cycling process, the continuous disproportionation reaction leads to the dissolution of lower-valence Mn, resulting in significant aging-induced deactivation of Mn on the material surface. This deactivation slows the reaction kinetics and causes detachment from the conductive network, ultimately leading to rapid capacity decay. The 2D oxidation state mapping of Mn in the cycled LMR@0.5% APO sample (Fig. 5e) shows consistent color regions distributed evenly in the whole secondary particles, which indicates that no obvious chemical-phase separation of Mn occurs after 200 cycles. This implies that the aging states of Mn are uniform, signifying the isotropic reactivity of Mn in LMR@0.5% APO during cycling, which can be attributed to the protective effects of the AlPO₄-coating layer. As shown in Fig. 5c, f, by statistically analyzing the pixels corresponding to different oxidation states in the 2D oxidation-state distribution mappings of Mn, we can directly observe a clear coexistence of multiple valence states in the pristine LMR sample after cycling. This demonstrates that the distribution of Mn valence states is nonuniform. In contrast, in the LMR@0.5% APO sample, the valence-state distribution of different pixels is more concentrated with only one center, demonstrating a significant improvement in the uniformity of the reaction.

Fig. 5Full-screen View

(Color online) 2D valence-state mappings (pixel scale: 30 nm) at Mn K-edge of the pristine (a) LMR and (d) LMR@0.5% APO samples; 2D valence-state mappings at Mn K-edge of the cycled (b) LMR and (e) LMR@0.5% APO samples. Pixel energy-distribution histogram of (c) LMR and (f) LMR@0.5% APO samples

Structural Changes Caused by Anionic Redox Reactions at High Voltage Investigated by Synchrotron-based X-ray Absorption Spectroscopy

The oxygen-oxidation reaction in LMR materials at high voltages can provide extraordinary capacity; however, it can also lead to issues such as surface oxygen release, transition-metal reduction, and densification, thereby affecting the electrochemical performance and stability of the materials during subsequent cycles. Therefore, we used both hard and soft XAS to investigate the evolution of the electronic and local coordination structures of the transition-metal and oxygen ions at high voltages before and after surface modification [49-51]. The O K-edge XAS test was carried out at beamline BL02B02 of the SSRF using the combined TEY and TFY modes [27]. In this mode, we can measure the photogenerated current on the material surface and fluorescence from the bulk of the material simultaneously; therefore, the information from the surface and bulk of the materials can be obtained simultaneously [41]. LMR materials are prone to severe surface oxygen release during cycling; hence, simultaneous tracking of O in both the surface and bulk is crucial for the mechanistic study of LMR materials. As shown in Fig. 6a, c, the broad spectral features above 535 eV are related to the transitions from O 1s to the TM 4sp hybridized O 2p states. The two pre-edge peaks at approximately 530 and 532 eV are attributed to the transitions from O 1s to the unoccupied TM d orbitals [52]. The pristine LMR sample surface exhibits a distinct signal of the cathode electrolyte interphase layer before the reaction. Within this signal, the pre-edge peak corresponding to the hybridization of the transition-metal 3d orbitals has a low intensity, and a carbonate-species peak is observed at 534 eV [53]. Owing to the gradual oxidation of the transition-metal ions, the degree of hybridization between the TM 3d and O 2p bands increases as charging proceeds, resulting in a significant increase in the intensity of the pre-edge peaks corresponding to the 1s to t_2g and e_g transitions. When the pristine LMR sample is charged to a high-voltage plateau (>4.4 V) and undergoes oxygen-oxidation reactions, lattice oxygen-oxidation signals (531 eV) can be detected on the sample surface. However, compared with the bulk phase signals obtained in the TFY mode (Fig. 6c, a significant loss of oxidized lattice oxygen is observed. This surface oxygen release corresponds to surface Mn reduction at high voltages observed at the Mn L-edge, which further facilitates the reduction of Mn and potentially leads to further disproportionation during the discharge process. When discharged to 2.75 V, the sXAS spectrum of the pristine LMR sample differs from its original state and still exhibits a strong pre-edge peak. This indicates that many transition-metal ions remain on the surface near the end of the discharge process, corresponding to the dissolution of transition metals observed at the Mn L-edge. In contrast, the LMR@0.5% APO sample shows significantly suppressed surface redox reactions owing to the presence of the AlPO₄-coating layer, whereas redox reactions in the bulk phase remain unaffected (Fig. 6b, d. Additionally, no transition-metal ions remain on the surface during the discharge process, demonstrating that the AlPO₄-coating layer effectively suppresses the reactivity on the surface of the sample to stabilize the surface structure, inhibiting Mn disproportionation and dissolution at low voltages. The ex-situ Mn, Co, and Ni K-edge XANES spectra in Fig. 6e, f, i, j, m, n, show noticeable differences in the evolution trends of the electronic structure and local coordination structure of the TM ions in the pristine LMR and LMR@0.5% APO samples under high voltage. In both samples, the shift to the right-hand side of the Ni K-edge is more pronounced than the shifts of the Co and Mn K-edges, which is consistent with Ni providing a larger capacity during the first plateau of the charging process. In the pristine LMR sample, the shifts in the Ni, Co, and Mn K-edges during the first charging plateau are larger than those in the LMR@0.5% APO sample, indicating that the AlPO₄-coating layer can also alleviate the overly rapid oxidation rate of TM ions in the bulk. This is consistent with the surface information obtained from sXAS mentioned above. During anionic oxidation, the valence states of the transition metals no longer undergo significant changes. The Fourier-transformed EXAFS of the Ni, Co, and Mn K-edges (Fig. 6g, h, k, l, o, p) illustrate that the AlPO₄-coating layer stabilizes the coordination structure effectively. The first coordination shell in the FT-EXAFS spectra of transition-metal oxides generally corresponds to TM-O bonds, and the second coordination shell generally corresponds to TM-TM bonds. Thus, as shown in Fig. 6g, h, compared to the pristine sample, the Mn-O and Mn-TM coordination corresponding to the first and second coordination shells, respectively, only undergo slight, uniform shifts in the LMR@0.5% APO sample. Moreover, as the reaction proceeds, the decrease in the intensity of the coordination peaks is even smaller. This demonstrates that the increase in the local structural disorder is alleviated, which is conducive to stabilizing the bulk structure and inhibiting the migration of transition metals and release of oxygen. In Fig. 6k, l, the Fourier-transformed EXAFS curves of the Ni K-edge show that Ni experiences a noticeable leftward shift in its first coordination shell Ni-O coordination peak due to providing a large number of electrons during the charging process. Compared to that of the pristine LMR sample, the Ni coordination structure of the coated sample charged to 4.4 V shows obvious distortion, owing to the retention of some Ni³⁺ in the bulk phase. After the pristine LMR sample is charged to 4.55 V, the Ni-O coordination peak in the sample continues to decrease, indicating progressively severe local structural distortion during the anionic redox process. In contrast, the Ni-O coordination peak in the coated sample exhibits a higher intensity than that in the pristine sample, and no significant change is observed in the intensity of the Ni-O coordination peak in the fully charged state. This indicates that the local structural distortion of Ni in the LMR@0.5% APO sample is effectively suppressed during the anionic redox stage. In Fig. 6o and p, the Co-O coordination also undergoes similar changes during anionic redox process. The changes in its coordination structure are minor compared to those of the pristine LMR sample, especially after being fully charged. The differences in the bond-length changes between the Ni and Co ions can be attributed to the different electronic structures of the Ni²⁺ and Co³⁺ ions. The strong σ interaction of the Ni-O bond renders it more sensitive to redox reactions than the weak π interaction of the Co-O bond [33]. In conclusion, the AlPO₄ coating layer can effectively stabilize the bulk structure of LMR materials under high voltages and alleviate local structural distortion and oxygen loss, which leads to structural collapse from the surface to the bulk during subsequent cycling, thereby effectively enhancing the electrochemical performance of the material.

Fig. 6Full-screen View

(Color online) Ex-situ XANES spectra during high-voltage plateau: ex-situ XANES spectra in TEY mode at O K-edge of pristine (a) LMR and (b) LMR@0.5% APO samples. Ex-situ XANES spectra in TFY mode at O K-edge of pristine (c) LMR and (d) LMR@0.5% APO samples. Ex-situ XANES spectra of pristine LMR sample at (e) Mn K-edge, (i) Ni K-edge, and (m) Co K-edge. Ex-situ XANES spectra of LMR@0.5% APO sample at (f) Mn K-edge, (j) Ni K-edge, and (n) Co K-edge. Ex-situ FT-EXAFS of pristine LMR sample at (g) Mn K-edge, (k) Ni K-edge, and (o) Co K-edge. Ex-situ FT-EXAFS of LMR@0.5% APO sample at (h) Mn K-edge, (l) Ni K-edge, and (p) Co K-edge