Surface BO₃ Configuration in Li‐Rich Cathode Materials Enabling Highly‐Stable Anionic Redox Reactions

ABSTRACT

Li‐rich Mn‐based layered oxides (LRMOs) are considered promising cathode candidates for next‐generation high‐energy‐density lithium batteries, owing to their high capacity and low cost. However, they are plagued by lattice‐oxygen release and surface‐driven structural degradation, which lead to low initial coulombic efficiency and poor cycling stability. Here, a B‐heterogeneous coordination structure is incorporated into the Li‐rich materials, forming a ≈4 nm surface layer enriched in BO₃ units while retaining BO₄ units within the bulk. Both of tetrahedral BO₄ and trigonal BO₃ display stronger bonding interaction than those of transition metal (TM)─O bonds (i.e., Mn─O, Ni─O, and Co─O), while surface BO₃ further strengthens the B─O bonds compared with bulk BO₄, thus robustly anchoring lattice oxygen to suppress irreversible oxygen loss. Benefiting from this synergistic heterogeneous coordination, the modified LRMOs deliver a high reversible capacity of ∼300 mAh g⁻¹ at 0.1C, an enhanced initial Coulombic efficiency of 93.5% and excellent capacity retention of 85.8% after 300 cycles at 1C. This work demonstrates the surface BO₃ structure as an effective paradigm to reconcile oxygen‐redox activity with long‐term stability in high‐energy‐density lithium batteries.

A spatially heterogeneous B─O coordination is constructed in Li‐rich Mn‐based oxide cathodes. The surface BO₃ units exhibit ultrahigh B─O bond strength to robustly anchor lattice oxygen, while bulk BO₄ units stabilize the layered structure during lithiation/delithiation. This bonding hierarchy enables reversible oxygen redox with long‐term stability.

1. Introduction

Driven by the swift proliferation of the electric vehicle industry, the development of next‐generation cathode materials is imperative for realizing lithium batteries with substantially higher energy density and lower manufacturing costs [1, 2, 3]. Among the available candidates, LRMOs are particularly attractive as their combined TM and anionic oxygen redox delivers specific capacities above 300 mAh g⁻¹, beyond the intrinsic limit of TM‐exclusive systems [4, 5]. However, the surface lattice oxygen in LRMOs is thermodynamically unstable at high voltage, leading to irreversible oxygen loss (O^2– → O₂) and progressive structural degradation [6]. The released oxygen further reacts with the electrolyte, triggering severe parasitic reactions, TM dissolution, and the formation of an unstable cathode–electrolyte interphase (CEI) [7, 8]. The coupled bulk and interfacial instabilities result in a low initial coulombic efficiency (ICE, typically below 85%) and rapid capacity fading, severely limiting the practical application of LRMOs in advanced lithium batteries [9].

To date, two principal strategies have been extensively investigated to mitigate these degradation pathways: surface coatings and bulk lattice engineering. Surface coatings have been widely explored to physically block oxygen escape and suppress interfacial side reactions, including Li₃PO₄, LiNbO₃, Li₂WO₄, and Li₂B₄O₇ [7, 8, 10, 11, 12]. An alternative strategy is to introduce dopants into the host lattice to strengthen metal–oxygen bonding and thereby enhance oxygen stability [13, 14]. For example, strong‐bonding motifs such as Mo–O, Al–O and Ti–O have been introduced into the bulk lattice to improve oxygen stability [15, 16, 17, 18, 19, 20]. However, surface coatings are often ineffective in suppressing bulk TM migration, whereas bulk‐strengthened TM–O bonding alone rarely curbs lattice‐oxygen loss at the surface. The homogenized design philosophy is intrinsically mismatched with the practical “surface‐to‐bulk” degradation nature of LRMOs [21, 22]. The surface region, where oxygen release and structural instability are most severe, requires a more robust bonding environment than the bulk. Moving beyond the conventional homogeneous doping paradigm, it is thus crucial to develop a spatially heterogeneous functional modification strategy that tailors the bonding environment along the particle depth and specifically reinforces the vulnerable surface region.

Here, we report a unique B–O heterogeneous coordination structure in LRMOs featuring BO₃ units at the surface (≈4 nm) and BO₄ units in the bulk, which is realized via a simple one‐step solid‐state reaction using NaBH₄ as the B source (Figure 1a). Density functional theory (DFT) calculations reveal that, compared with M–O bonds (M = Ti, V, Cr, P, Al, S, Cu, Zn) introduced by other dopants in LRMOs, both BO₃ and BO₄ units exhibit significantly higher B–O bond energies, with BO₃ showing the highest bond energy among all the considered M−O units (Figure 1b). The exceptionally strong surface BO₃ bonds provide robust anchoring, effectively suppressing oxygen release, while bulk BO₄ tetrahedra act as “structural pillars” to restrain TM migration without over‐stabilizing lattice oxygen, thereby preserving reversible oxygen redox activity. As a result, the modified LRMO (denoted as LRMOB) delivers an enhanced ICE of 93.5% at 0.1C, along with a high reversible capacity of ∼300 mAh g⁻¹ at 0.1C and superior cycling stability (capacity retention improved from 65.7% to 85.8% after 300 cycles at 1C). This heterogeneous‐coordination strategy offers a promising and scalable avenue for the design of next‐generation high‐energy‐density and durable cathode materials.

FIGURE 1.

Material synthesis and optimization strategy. (a) Schematic diagram of B‐species incorporation into the host lattice. (b) DFT calculation of bond energies for different dopant elements.

2. Results and Discussion

Li_1.2Ni_0.13Co_0.13Mn_0.54O₂ (LRMO) was synthesized via a solid‐state reaction using Li₂CO₃ and carbonate precursor obtained through a co‐precipitation process. The modified sample (LRMOB) was prepared by applying a NaBH₄ pretreatment to the LRMO. Detailed synthesis procedures are provided in the Supporting Information.

The scanning electron microscopy (SEM) images of the LRMO and LRMOB (Figure S1) reveal that both samples possess a well‐defined spherical morphology with an average diameter of approximately 10 µm, consisting of densely packed nanosized primary particles. Notably, the boronization process preserves the integrity of the microspherical architecture without inducing any noticeable morphological degradation (Figure S2). Furthermore, scanning transmission electron microscopy with energy‐dispersive x‐ray spectroscopy (STEM‐EDS) elemental mapping and focused ion beam‐scanning electron microscopy (FIB‐SEM) analysis of the LRMO (Figures S3 and S4) confirm the uniform distribution of Co, Ni, Mn, and O elements within the primary particles. To verify the bulk incorporation of B in LRMOB, cross‐sectional lamellae of the secondary particles were prepared using FIB‐SEM (the sample preparation process is illustrated in Figures S5 and S6). Subsequent STEM‐EDS elemental mapping clearly demonstrates that Co, Ni, Mn, O, and B are homogeneously distributed throughout the bulk region of LRMOB, confirming successful B doping within the bulk (Figure S7). Similarly, depth analysis x‐ray photoelectron spectroscopy (XPS) was conducted (Figure S8), showing that B remains detectable even after etching to a depth of approximately 50 nm, confirming the incorporation of B within the bulk structural phase [10, 11]. Furthermore, quantitative analysis by inductively coupled plasma optical emission spectroscopy (ICP‐OES) indicates that the boron content in LRMOB is generally consistent with the nominal feeding ratio (Table S1), further confirming the successful incorporation of B into the material. X‐ray diffraction (XRD) patterns of the LRMO and LRMOB are presented in Figure 2a,b. Both samples exhibit well‐defined diffraction peaks characteristic of the α‐NaFeO₂‐type layered structure with a space group of R‐3m. In addition, the minor reflections observed in the range of 20°–22° can be attributed to the Li₂MnO₃‐type monoclinic phase (C2/m) [23, 24, 25]. As shown in Figure S9, the enlarged (003) peak of LRMOB shifts toward a lower diffraction angle compared with that of the LRMO, indicating that B ions are successfully incorporated into the host lattice [26]. Additionally, both the lattice parameter c and the unit‐cell volume exhibit a slight increase for LRMOB (Table S2). The widening of the Li layers observed in the high‐angle annular dark‐field scanning transmission electron microscopy (HAADF‐STEM) images further corroborates this conclusion. Specifically, the Li‐layer spacing in LRMOB reaches 0.488 nm, exceeding the 0.477 nm spacing observed in the LRMO, thereby demonstrating this incorporation expands the (003) interplanar spacing and facilitates the broadening of Li‐ion diffusion channels in LRMOB (Figure S10). Meanwhile, the Raman spectra further confirm that the boronization treatment does not induce any phase transformation. For the two samples, two prominent peaks located at approximately 595 and 480 cm⁻¹ correspond to the A_1g and E_g vibrational modes of the R‐3m layered structure, while the additional weak features observed in the range of 350–450 cm⁻¹ are associated with the A_g mode of the monoclinic Li₂MnO₃ phase (Figure S11) [7, 12].

FIGURE 2.

Structural characterization of the LRMO and LRMOB. (a) and (b) Rietveld refinements of XRD patterns for LRMOB and the LRMO. (c) EPR spectra of the LRMO and LRMOB. (d) and (e) HAADF image and corresponding O K‐edge EELS line‐scan intensity map of the LRMO and LRMOB particles. (f) ¹¹B solid‐state SS‐NMR spectrum of LRMOB. (g) Structural models with V_O at two distinct lattice sites. (h) System energies of models containing V_O at two distinct lattice sites.

To verify the formation of surface BO₃ coordination structures, electron paramagnetic resonance (EPR) spectra (Figure 2c) reveal that LRMOB exhibits a markedly stronger oxygen‐vacancy (V_O) signal compared with the LRMO [27, 28, 29, 30, 31]. Furthermore, the intensity of the V_O signal increases progressively with the amount of NaBH₄ introduced during synthesis (Figure S12), suggesting that the boronization treatment effectively facilitates the generation of oxygen vacancies. Figure S13 presents the O 1s spectra of the samples. The peaks at 529.4 and 530.5 eV are attributed to lattice oxygen and V_O, respectively, while the peak at 531.5 eV corresponds to adsorbed oxygen [16]. Notably, the characteristic signals of V_O are clearly observed in the treated samples. To further investigate the nanoscale spatial distribution of V_O, HAADF‐STEM imaging and electron energy‐loss spectroscopy (EELS) line‐scan analyses were conducted using STEM. As illustrated in Figure 2d,e, both the LRMO and LRMOB maintain the typical layered structure across their surface and bulk regions. In contrast, for LRMOB, the pre‐edge peak of the O K‐edge at approximately 533 eV exhibits a gradual decrease in intensity from the bulk toward the surface, whereas no discernible variation is observed for the LRMO. This observation indicates that V_O is primarily concentrated near the surface of LRMOB, forming a defective layer approximately 4 nm thick [8, 32, 33, 34]. This observation aligns well with the findings derived from XPS etching (Figure S14). The solid state nuclear magnetic resonance (SSNMR) spectrum shown in Figure 2f reveals a prominent peak at 0 ppm, corresponding to 4‐coordinated B ions, alongside a smaller peak at 10 ppm, which can be attributed to 3‐coordinated B ions [35, 36]. This feature likely arises during the boronization process: in the presence of V_O, B can coordinate with three O atoms and one V_O, forming a BO₃‐V_O‐type 3‐coordinated environment. Figure 2 g,h further confirms that the calculations show that when V_O and B coexist, the configuration where V_O occupies a three‐coordinated site is energetically more favorable than the 4‐coordinated adjacent site.

These results indicate that in LRMOB, B tends to adopt a three‐coordinated configuration at the surface, whereas in the bulk it predominantly exists in a 4‐coordinated state. On this basis, we performed DFT calculations for four possible configurations in which the V_O occupies one of the four oxygen sites surrounding a BO₄ unit, denoted as site 1, site 2, site 3, and site 4 (Figure S15). The computed total energies reveal the stability order to be site 1 > site 2 > site 3 > adjacent site >site 4. Among these, site 4 is the least stable because the three coordinating O atoms lie within the same TM layer; the incorporation of B at this position exerts steric pressure on the neighbouring Mn, causing a significant displacement from its original TM site. In contrast, the relative stabilities of the other three configurations correlate with their respective B–O bond lengths; the structure in which V_O occupies the site associated with the longest B–O bond is the most energetically favorable. Consequently, the resulting BO₃ motif preserves the three shortest B–O bonds of the original BO₄ tetrahedron. Furthermore, the transition from a tetrahedral to a trigonal‐planar coordination further shortens these B–O bonds, thereby significantly enhancing the ability of B to anchor the surrounding lattice oxygen. The L‐edge EELS spectra of Mn, Co, and Ni in the LRMO show no discernible edge shifts, indicating that the valence states of the TM ions remain uniform throughout the material (Figure S16b–e). In contrast, the L‐edge spectra collected from LRMOB exhibit a gradual shift of the Mn L‐edges toward lower energy from the bulk to the surface (Figure S16g–j), implying the progressive reduction of Mn ions coordinated with oxygen at the surface. XPS analyses (Figure S17) of Ni, Co, and Mn for both the LRMO and LRMOB further corroborate these findings; while the Ni and Co valence states are essentially identical in the two materials, the Mn 2p spectrum of LRMOB displays a red shift relative to that of the LRMO, confirming the reduced Mn state in LRMOB. The reduction of the Mn valence state from 3.90 to 3.82 is corroborated by the Mn 3s spectra shown in Figure S18 [37]. To further directly probe the Mn valence state and local coordination environment, Mn K‐edge X‐ray absorption fine structure (XAFS) measurements were performed for LRMO and LRMOB. As shown in Figure S19, the X‐ray absorption near edge structure (XANES) spectra show that the Mn K‐edge of LRMOB exhibits a slight shift toward lower energy relative to that of LRMO, indicating a slightly reduced average Mn valence state in LRMOB. The observed TM valence state features further reinforce the conclusion that LRMOB possesses a heterogeneous B‐coordination configuration. Neutron powder diffraction (NPD) data and corresponding Rietveld refinement analyses were performed to confirm the structural models, with the unit cell parameters summarized in Figure S20 and Table S3. The corresponding crystal structures have been deposited with the CCDC/FIZ Karlsruhe Access Structures service, with the assignment of deposition numbers to each structure provided in the subsequent description following Figure S20.

To assess the influence of the reductive boronization on the electrochemical performance, coin‐type half cells were assembled and tested within the voltage range of 2–4.8 V. As shown in Figure 3a, the initial charge–discharge profiles at 0.1C (1C = 200 mA g⁻¹) exhibit the typical redox behavior characteristic of LRMOB and the LRMO. The sloping region below 4.5 V primarily corresponds to the oxidation of TM cations, associated with the oxidation of Ni²⁺ and Co³⁺ to Ni⁴⁺ and Co⁴⁺, respectively. Meanwhile, the pronounced plateau near 4.5 V is ascribed to the oxidation of lattice oxygen anions (O^2– → O₂ ^n– → O₂, 0< n < 4). Due to the inhibitory effect of B on oxygen activity, the oxygen‐redox‐derived capacity of LRMOB (124.2 mAh g⁻¹) is lower than that of the LRMO (138.6 mAh g⁻¹). Nevertheless, the strong lattice‐oxygen bonding associated with the surface BO₃ units in LRMOB significantly mitigates irreversible oxygen loss, thereby markedly enhancing the ICE. Consequently, LRMOB delivers a high reversible capacity of 298.44 mAh g⁻¹ along with an exceptional ICE of 93.5%, in contrast to the LRMO, which exhibits a reversible capacity of 280.94 mAh g⁻¹ and a lower ICE of 82.1%. Moreover, benefiting from the strong bonding effect of the surface BO₃ structure on lattice oxygen, LRMOB maintains an impressive capacity retention of 85.8% after 300 cycles at 1C, which is significantly higher than that of the LRMO (65.7%) (Figure 3b). Similarly, the energy density decay of LRMOB is markedly lower than that of LRMO (Figure 3c). The capacity retention and energy density data are presented as mean values ± s.e.m. and the sample size is 3. The dQ/dV curves further indicate that LRMOB exhibits higher reversibility than the LRMO (Figure 3e,f). In addition, because interstitial B doping effectively suppresses TM migration, the voltage fade of LRMOB is limited to only 1.07 mV per cycle, markedly lower than that of the LRMO (1.48 mV per cycle) (Figure S21). Among the samples prepared with different NaBH₄ contents in an Ar atmosphere, LRMOB (i.e., LRMOB‐0.50%) shows the best cycling stability and rate performance compared with LRMO‐Ar, LRMOB‐0.25%, and LRMOB‐1.00% (Figure S22). Compared with the LRMO and other reported LRMO‐based materials, the modified LRMOB sample exhibits pronounced advantages in terms of ICE, cycle life, and capacity retention (Figure 3d) [7, 12, 26, 38, 39, 40, 41, 42, 43, 44]. The LRMOB electrode delivers discharge capacities of 298.3, 274.9, 250.2, 231.2, 205.7, and 165.1 mAh g⁻¹ at current rates of 0.1, 0.2, 0.5, 1, 2, and 5C, respectively, which are markedly higher than those of the other samples (Figures 3g and S22b). To gain deeper insight into the kinetic behavior of LRMOB and the LRMO, the galvanostatic intermittent titration technique (GITT) was employed to evaluate the Li⁺ diffusion coefficient (D_Li ⁺) (Figure S23 and Table S4). The evolution of the cell voltage and the corresponding D_Li ⁺ values during the initial charge–discharge process is shown in Figure S23c,d. The Li⁺ diffusion coefficient decreases sharply at around 4.5 V during charging and 3.3 V during discharging, which can be attributed to the activation and structural rearrangement of the Li₂MnO₃ phase; this activation step requires higher energy and thus leads to a lower D_Li ⁺. Overall, the D_Li ⁺ values of LRMOB are consistently higher than those of the LRMO throughout the initial cycle, indicating faster Li⁺ transport and a lower charge‐transfer energy barrier for LRMOB due to the multifunctional interfacial layers, consistent with its superior rate capability.

FIGURE 3.

Electrochemical performance of the LRMO and LRMOB. (a) Initial charge‐discharge curves at 0.1C. (b) Capacity retention and corresponding coulombic efficiency at 1C. (c) Energy density retention at 1C. (d) Comparison of LRMOB with other reported LRMOs in terms of ICE and cycling stability. (e) and (f) dQ/dV profiles of the LRMO and LRMOB at 1C. (g) Rate performance of the LRMO and LRMOB. Data are presented as mean values  ±  s.e.m. and the sample size is 3.

To elucidate the evolution of lattice oxygen and further evaluate the stability of oxygen anions, we conducted the following characterizations. Ex situ XPS measurements of the O 1s region were performed to probe the oxygen‐redox behavior of both electrodes during the first cycle. As shown in Figure 4a, the peaks located at approximately 529.7, 532.0, and 533.7 eV are assigned to lattice O^2–, oxygenated surface species, and electrolyte decomposition products, respectively. When charged to 4.5 V, a shoulder peak attributable to oxidized oxygen species (O₂ ^{n –}) appears at around 531 eV in both samples [45, 46]. The fraction of the O₂ ^{n −} component in LRMOB reaches 8.01%, in contrast to 7.31% in the LRMO. Upon further charging to 4.8 V, the O₂ ^n– contribution in LRMOB increases to 12.54%, indicating a stronger stabilization of high‐valent oxygen and reduced oxygen loss, which can be ascribed to the strong anchoring effect of the BO₃ units on the surface lattice oxygen. In comparison, the O₂ ^{n –} signal in the LRMO accounts for only 10.79% of the surface oxygen at 4.8 V, implying more severe oxygen release. After discharging to 2.0 V, the O₂ ^{n –} peak disappears, consistent with the reduction of O₂ ^n– back to lattice O^2– (Figure S24). Figure 4b,c shows the density of states (DOS) of Mn and O, together with the total density of states, for the pristine LRMO and BO₃‐modified LRMOB. In the pristine LRMO, one of the O‐2p orbitals in the Li‐O‐Li configuration corresponds to a nonbonding state and is responsible for activating anionic redox. Upon introducing BO₃ units, the top of the O‐2p band shifts downward from –0.12 to –1.11 eV, indicating a reduced tendency for lattice oxygen to be oxidized to O₂. This downward shift stabilizes lattice oxygen and enhances the reversibility of anionic redox at high voltages. In addition, Li₂MnO₃ containing BO₃ units exhibits a narrower bandgap than pristine Li₂MnO₃, consistent with the improved electronic conductivity induced by B doping. We further compared the O‐2p band maximum and band gap of a Li₂MnO₃ model containing only oxygen vacancies (LRMO‐V_O) and a Li₂MnO₃ model containing only BO₄ units (LRMO‐B). LRMOB combines the short band gap of LRMO‐V_O with the low O‐2p band maximum of LRMO‐B, showing the lowest O‐2p band and band gap among all four models (Figure S25). These results demonstrate that surface BO₃ groups endow the material with highly reversible anionic oxygen activity together with enhanced electronic conductivity. To further evaluate the reversibility of anionic redox, differential electrochemical mass spectrometry (DEMS) was performed to monitor lattice oxygen release and quantify O₂ evolution during charging (Figure 4f,g). A detectable O₂ signal appears at high voltages as a result of lattice–oxygen oxidation. Notably, compared with the LRMO, LRMOB shows minimal O₂ release, which can be ascribed to the strong oxygen‐anchoring effect of the surface BO₃ units. Since each BO₃ unit incorporates one V_O, EPR measurements were carried out on cycled electrodes to assess the stability of the BO₃ unit after prolonged cycling. The cycled LRMO electrode exhibits a pronounced oxygen‐vacancy signal. In contrast, the LRMOB electrode collected after cycling shows only a modest increase in signal intensity (Figure 4d,e), indicating that the BO₃ units in LRMOB possess higher structural stability.

FIGURE 4.

Anion redox behaviors in the LRMO and LRMOB. (a) Evolution of O 1s XPS spectra at different voltage states. (b) and (c) Calculated DOS of LRMOB and the LRMO. (d) and (e) EPR spectra of LRMOB and the LRMO after 300 cycles, which were normalized based on carbon intensity. (f) and (g) Voltage profiles and the corresponding O₂ evolution of the LRMO and LRMOB obtained from in situ DEMS measurements at 0.1C.

We further employed a series of characterization techniques to investigate the effect of the BO₃ structure on the chemical evolution of the cathode–electrolyte interphase. As shown in Figure 5a,b, high‐resolution transmission electron microscopy (HRTEM) analysis reveals that LRMOB forms a more compact and uniform CEI layer than the LRMO. Such a CEI layer promotes efficient Li⁺ transport and suppresses various interfacial side reactions, thereby protecting the internal structure of the LRMO and indicating a more stable interface during cycling. Compared with the LRMO, LRMOB exhibits a distinct TM‐O peak at ∼529.5 eV in the O 1s spectrum (Figure 5d), further indicating that the CEI passivation layer formed after prolonged cycling is thinner than that on the LRMO [47]. The Mn 2p signal of LRMOB is markedly weaker than that of the LRMO, indicating that its BO₃ unit effectively suppresses Mn‐ion dissolution from the layered structure and their subsequent surface deposition (Figure 5c) [48, 49]. In Figure S26a, the XPS C 1s spectrum shows a pronounced decrease in the intensity of the characteristic Li₂CO₃ peak at 290.5 eV for LRMOB, indicating fewer parasitic side reactions at the electrode‐electrolyte interface. Consistently, quantitative analysis of the LiF content (Figure S26b) reveals a marked reduction from 30.52% in the LRMO to only 18.26% in LRMOB, demonstrating that the multifunctional interfacial layers effectively suppress electrolyte decomposition and significantly mitigate interfacial side reactions.

FIGURE 5.

Interfacial evolution of the LRMO and LRMOB. (a) and (b) HRTEM images of LRMOB and the LRMO after 10 cycles at 1C. (c) and (d) Mn 2p and O 1s XPS spectra of the LRMO and LRMOB electrodes after 250 cycles at 1C. (e) and (f) DRT profiles transformation from in situ EIS data of LRMOB and the LRMO during the initial charge‐discharge process. (g) and (h) 3D distribution visualization images and the corresponding 2D overlay mapping of LRMOB and the LRMO.

Time‐of‐flight secondary ion mass spectrometry (TOF‐SIMS), combined with 3D surface reconstruction, 2D overlay mapping, and depth profiling, was employed to elucidate the effect of interfacial boronization on the chemical composition of the CEI layer after 250 cycles (Figures 5 g,h and S27–S29). Characteristic species such as Li₂CO₃ ^–, organic fragments (C₂HO^–, CHO₂ ^–, C₂F^–), P‐containing species (POF₂ ^–), and LiF₂ ^– are generated by the decomposition of the electrolyte solvent and LiPF₆ at high voltages. In addition, acidic HF can readily attack the electrode surface, inducing TM dissolution and structural degradation, ultimately leading to the formation of TM fluorides (MnF₃ ^–, NiF₃ ^–, CoF₃ ^–) that persist within the CEI layer. For the LRMO, the signals associated with these species are broadly distributed and penetrate deeply into the particles, indicative of severe interfacial side reactions. In contrast, LRMOB exhibits markedly attenuated signal intensity, which can be attributed to strong BO₃ bonding that suppresses the attack of high‐valent oxygen species on the lithium salt and thus inhibits electrolyte decomposition. To elucidate the evolution of interfacial Li⁺ kinetics during both charge and discharge, in situ galvanostatic electrochemical impedance spectroscopy (EIS) combined with distribution of relaxation times (DRT) analysis was carried out over the first cycle (Figure S30). This technique provides a time‐scale‐resolved diagnostic tool that markedly enhances the reliability of kinetic interpretation. Unlike conventional equivalent‐circuit fitting, the DRT approach directly analyzes the impedance spectra, identifies dynamic processes based on their characteristic time constants, and avoids subjective errors associated with manual circuit selection. Moreover, DRT analysis based on impedance measurements enables efficient deconvolution of individual electrochemical processes, with the corresponding impedance values obtained from the integrated area of each characteristic peak. Typically, peaks located at τ ≈ 10⁻³–10⁻¹ s and 10⁻¹–10 s are assigned to the charge‐transfer resistance (R _ct1) and interfacial diffusion resistance (R _ct2), respectively. Throughout the entire process, LRMOB shows a lower internal resistance than the LRMO. In the DRT spectra, the peaks labeled R _ct1 and R _ct2 correspond to charge‐transfer resistance at the CEI, while the R_CEI/SEI peak is associated with the CEI/SEI (solid electrolyte interphase). For the LRMO, R _ct2 rises sharply at high voltages, indicating irreversible O₂ release and concomitant structural phase transitions. In contrast, the interface resistance of LRMOB remains consistently lower than that of the LRMO, suggesting that the BO₃ unit effectively suppresses lattice‐oxygen evolution and the resulting interfacial side reactions.

Beyond inhibiting surface O₂ release and stabilizing the surface structure, an important question is whether the heterogeneous B coordination can also suppress strain accumulation within the bulk. The Raman spectra of the LRMO and LRMOB after 250 cycles are shown in Figure S31. For LRMOB, the characteristic E_g and A_1g modes of the R‐3m layered structure are well preserved, accompanied by several weak bands in the 350–450 cm–¹ region that can be assigned to the A_g mode of the Li₂MnO₃ phase. The left‐shifted band at around 630 cm⁻¹ is ascribed to the A_1g mode of a spinel‐like phase. In contrast, for the LRMO, the symmetric M─O stretching band of the pristine material splits into two left‐shifted peaks at approximately 570 and 630 cm⁻¹, which are attributed to the A_1g modes of the R‐3m layered structure and the spinel‐like structure, respectively. Meanwhile, the series of weak Li₂MnO₃‐related bands vanish, and a new band emerges at around 375 cm⁻¹, which is assigned to the F_2g mode associated with the symmetric stretching of LiO₄ tetrahedra, providing strong evidence for the formation of a spinel‐like phase [22, 44, 50]. Compared with the LRMO, LRMOB exhibits a significantly smaller fraction of spinel‐like domains and more effectively preserves the layered framework, indicating that the BO₃ unit suppresses TM migration and irreversible surface oxygen release, thereby maintaining structural integrity during cycling. Consistently, SEM images (Figure S32) show that the secondary spherical particles of the LRMO become severely damaged and cracked after 250 cycles, whereas LRMOB largely retains an intact spherical morphology.

Furthermore, the LRMOB samples were bisected using FIB‐SEM, and the cross‐sectional observations revealed no tendency toward internal pulverization (Figure S32e,f). To evaluate the differences in long‐term cycling performance between LRMO and LRMOB, the evolution of interfacial resistances and lithium‐ion diffusion behavior was analyzed. After three activation cycles at 0.1C, the cells were subsequently cycled at 1C for 50 cycles, during which in situ EIS combined with DRT fitting was performed to investigate the evolution of interfacial resistance and lithium‐ion diffusion behavior over long‐term cycling. The results show that LRMOB exhibits a lower initial impedance and a more stable impedance evolution than LRMO, which can be attributed to the BO₃ structure effectively suppressing the overactivation of lattice oxygen, thereby mitigating parasitic electrolyte decomposition and stabilizing the CEI, leading to improved interfacial stability and enhanced charge‐transport kinetics (Figure S33). To further elucidate the effect of the heterogeneous B coordination on the bulk lattice, we performed in situ XRD measurements on LRMOB and the LRMO during the initial charge–discharge cycle. As shown in Figure S34, the asterisks denote the Be/BeO diffraction peaks arising from the window of the cell mold. Figures 6a,b and S35 displays the charge–discharge profiles of the in situ cells, together with the corresponding contour plots of the (003) and (101) reflections. The evolutions of the lattice parameters a and c obtained by Rietveld refinement are also summarized. For the LRMO, upon charging to 4.5 V, the (003) peak gradually shifts to lower angles, corresponding to a 1.196% increase in c‐spacing. This expansion is ascribed to the enhanced electrostatic repulsion between adjacent oxygen layers during Li extraction [13, 14]. Concurrently, the a‐spacing decreases by 0.702% due to the oxidation‐induced shrinkage of the TM ions [48, 49]. As charging proceeds from 4.5 V to the end of charge, the repulsive interaction between the oxygen layers is weakened because of electron depletion on the oxygen ions [50], resulting in a shift of the (003) peak toward higher angles and a net contraction of the c‐spacing by 1.279%. In LRMOB, a similar trend is observed upon charging to 4.5 V, with an expansion along the c‐axis (1.08%) and a contraction along the a‐axis (0.666%). However, beyond 4.5 V, the variation in c‐spacing (–0.825%) is significantly smaller than that of the LRMO. This mitigated contraction suggests that trapped O₂ within the (003) planes in the bulk hampers excessive c‐axis shrinkage in LRMOB. During discharge, continuous Li intercalation into the (003) layers leads to an expansion of the c‐axis by 0.866% in the LRMO and 0.539% in LRMOB, while the gradual reduction of TM ions results in an a‐axis expansion of 0.987% and 1.235%, respectively. These results indicate that the heterogeneous B coordination structure effectively enhances bulk structural stability, particularly in the high‐voltage regime, by suppressing excessive lattice distortion [51]. Overall, the LRMO bulk undergoes pronounced strain fluctuations during the charge‐discharge process, whereas the heterogeneous B coordination structure in LRMOB helps to alleviate strain within the bulk. To further verify this, we collected HAADF‐STEM images for both samples in the pristine state and after ten cycles at the charged state (Figure 6c–f). Compared with the well‐defined layered ordering in the pristine samples, the bulk lattice of the LRMO exhibits evident distortion after ten cycles, whereas LRMOB largely retains its original layered framework. These observations corroborate that the heterogeneous B coordination structure effectively suppresses strain accumulation in the bulk and thereby enhances the structural robustness during cycling.

FIGURE 6.

Structural stability characterized by in situ XRD and atomic‐scale imaging. (a) and (b) Initial charge–discharge curves of the in situ cell at 0.2C and the corresponding lattice parameters obtained by fitting XRD data of the LRMO and LRMOB. (c) and (d) HAADF images of pristine LRMOB and the sample after 10 cycles at 0.1C. (e) and (f) HAADF images of the pristine LRMO and the sample after 10 cycles at 0.1C.

3. Conclusion

In summary, we have developed a B‐heterogeneous coordination structure to address the characteristic “surface‐to‐bulk” degradation mechanism in Li‐rich layered cathodes. At the surface, BO₃ units with strong oxygen‐binding capability effectively suppress irreversible oxygen release and parasitic reactions between oxidized oxygen species and the electrolyte. In the bulk, BO₄ units inhibit TM migration, thereby stabilizing the layered framework. As a result, the modified LRMOB sample exhibits significantly enhanced anionic redox reversibility and structural robustness, achieving an initial Coulombic efficiency of 93.5%, a high capacity retention of 85.8% after 300 cycles at 1C, and a substantially mitigated voltage decay of only 1.07 mV per cycle (compared to 1.48 mV per cycle for the pristine sample). These findings demonstrate that the unique BO₃ structure effectively facilitates Li⁺ transport, suppresses bulk strain evolution, prevents irreversible oxygen loss, and reduces interfacial side reactions, thereby providing a robust foundation for designing high‐energy‐density cathode materials with superior electrochemical performance.

Conflicts of Interest

The authors declare no conflicts of interest.

Supporting information

Supporting File: anie73295‐sup‐0001‐SuppMat.docx.

Data Availability Statement

The data that support the findings of this study are available from the corresponding author upon reasonable request.

Accession Codes: CCDC2556720‐2556723 contain the supporting crystallographic data for this paper. These data can be obtained free of charge via www.ccdc.cam.ac.uk/data_request/cif, or by emailing data_request@ccdc.cam.ac.uk, or by contacting The Cambridge Crystallographic Data Centre, 12 Union Road, Cambridge CB2 1EZ, UK; fax: +441223336033.