Flash joule heating-induced spinel-phase surface in Ni-rich layered oxide positive electrodes to stabilise lattice oxygen

Domain: biology
Published: 2026-03-15
Editorial outcome: minor_revision

Abstract
High-nickel layered positive electrodes suffer from progressive structural degradation arising from lattice oxygen loss and inherent lattice strain. Although surface coatings are widely used to stabilize lattice-oxygen redox and mitigate electro-chemo-mechanical degradation, achieving coatings with full continuity, robust interfacial bonding, and fast Li+ conductivity remains challenging. Herein, we present a fundamentally different approach to shell formation via a self-derived subtractive strategy, departing from the conventional additive-based coating methods. By accurately applying transient thermal pulses, surface lithium is selectively extracted from layered LiNi0.8Co0.1Mn0.1O2 (x=0.8~0.9), directly converting the outer region into a coherent spinel-phase shell with tunable thickness. This nanoscale spinel-phase skin forms a robust mortise-and-tenon-like interconnection with the layered bulk, enabling isotropic, high-rate Li+ extraction/insertion while maintaining electronic conductivity throughout cycling. It effectively confines active oxygen intermediates, and suppresses interfacial side reactions and strain evolution under high-potential operation. Therefore, the spinel-phase skin-encapsulated LiNi0.8Co0.1Mn0.1O2 achieves an initial Coulombic efficiency of 95.3% and enables pouch cells with 80.1% capacity retention after 2000 cycles at 180 mA g-1. This strategy is extendable to LiNi0.9Co0.05Mn0.05O2, may open new avenues for advancing nickel-rich positive electrode technologies.

Introduction
High-nickel layered oxides, particularly LiNi x Co y Mn 1-x-y O 2 (NCM, x ≥ 0.8), are among the most promising positive electrode candidates for next-generation automotive power batteries due to their energy density[1-4]. However, rapid capacity degradation remains the most challenging issue in the commercialization of NCM. It is now widely recognized that this degradation originates primarily from lattice oxygen oxidation accompanying the H2-H3 phase transition at elevated potential (> 4.1 V)[5,6]. Lattice oxygen oxidation not only leads to gasous oxygen evolution but also generates reactive oxygen intermediates that oxidize the electrolyte, triggering irreversible phase transitions from the layered structure to spinel/rock-salt phases[7]. These structural changes[8], which initiate at the surface and propagate inward, result in a low initial Coulombic efficiency (ICE, 85~89%). The formation of an electrochemically inactive rock-salt phase at the cathode-electrolyte interface (CEI) significantly impedes Li+ and electron transport, inducing local charge heterogeneity[9-12]. Simultaneously, oxygen loss amplifies the intrinsic anisotropic lattice strain of NCM, leading to internal particle cracking[6,13,14] and mechanical pulverization[15,16]. Collectively, these degradation pathways cause the progressive loss of potential plateau and reversible capacity. Thus, suppressing lattice oxygen loss is essential for improving the electrochemical reversibility and longevity of NCM electrodes.

Given that the lattice oxygen oxidation is an intrinsic nature of highly delithiated NCM, an effective strategy must suppress O2 release and mitigate side reactions between active oxygen intermediates and the electrolyte. Surface coating[17-25] stands out as the most direct and promising solution, but to be effective, the coating must meet various stringent requirements. Ideally, coating layer should provide nearly complete coverage, support high-rate electron and Li+ transport, exhibit extraordinary chemical/ electrochemical stability, and form strong mechanical and chemical bonds with the bulk phase. Traditional coating methods, primarily based on additive processes such as dry and wet chemical deposition of metal oxides[17,18,20], fluorides[22], phosphides[23] and solid Li+ conductor[24,25], can largely reduce direct contact between the electrode and electrolyte, but often fail to achieve full coverage and sufficient interfacial bonding. Advanced thin-film deposition techniques such as physical/chemical vapor deposition and atomic layer deposition[26-28] ensure high coverage and thickness control, yet suffer from low ionic conductivity and lattice mismatch with the underlying NCM.

To overcome these limitations, a fundamentally different strategy is needed. By reverse thinking about the limitations of additive methods, a subtractive strategy[29-34] that precisely removes certain atoms from the outermost layer of NCM can enable the in situ formation of a self-derived, conformal coating shell. Previous studies have hinted at this potential: Lu[30] et al. proposed that Ar/H2 plasma treatment induced oxygen vacancies on the NCM surface, which subsequently transformed into a spinel-phase shell after electrochemical cycles. Similarly, Zhao[32] et al. utilized this approach to produce a rock-salt coating on LiCoO2, and Chen[34] et al. generated a Li-deficient thin layer on Li-rich layered oxides by adjusting lithium source amount during synthesis. However, these methods still fall short in terms of controllability, scalability and shell integrity, as reflected in only marginal improvements in the ICE.

In this work, we disclose a precisely controlled subtractive strategy using flash Joule heating to quantitatively remove surface lithium/oxygen atoms from NCM. This technique induces an in situ transformation of NCM edges into a coherent and perfectly compatible spinel-phase skin layer with tunable thickness and complete coverage. By precisely controlling the temperature and duration of transient thermal pulses, either spinel-phase or rock-salt-phase coating layers can be selectively formed (Figure 1). In particular, the nanoscale spinel-phase skin establishes a robust, mortise-and-tenon-like interconnection with the bulk phase, enabling isotropic and high-rate Li+ extraction/insertion and continuous electron transport across the interface during long-term cycling. This conformal shell tightly confines the active oxygen intermediates, eliminates local charge heterogeneity, and significantly suppresses irreversible layered-to-spinel/rock-salt phase transition both at the surface and within the bulk. This greatly contributes to enhancing reaction reversibility, improving the ICE, and mitigating strain evolution and particle cracking. Consequently, the spinel-phase skin-encapsulated LiNi0.8Co0.1Mn0.1O2 achieves an high initial Coulombic efficiency of 95.3% and retains 93.6% capacity after 500 cycles, significantly surpassing the pristine counterpart (89.6% efficiency and 65.2% retention). Moreover, the optimally engineered NCM electrode boosts the retention rate of pouch cells from 52.0% to 80.1% after 2000 cycles. More importantly, the spinel-phase skin is also applicable to higher-nickel NCM electrodes such as LiNi0.9Co0.05Mn0.05O2 and LiNi0.825Co0.115Mn0.06O2. This corrosion-resistant and crack-resistant skin paves the way for enhancing the long-lifespan of high-nickel electrodes.

Results
Self-derived growth of phase-customizable skins on NCM811

Figure 1a presents a schematic illustration of the customized modulation of the interfacial phase structure of LiNi0.8Co0.1Mn0.1O2 (NCM811) using flash Joule heating technology. Typically, commercial NCM811 powder was placed into the reaction chamber of an independently developed transient high-temperature equipment (Supplementary Fig S1a) and subjected to flash sintering under an argon atmosphere using thermal pulses with precisely-controlled temperatures and pulse widths. Notably, the NCM811 from a single processing batch can reach 5~10 g (Supplementary Fig S1b). Under transient thermal pulses, partial lithium and oxygen atoms at the surface/subsurface are sequentially stripped, leading to the formation of Li-deficient or Li&O co-deficient phases[30-32]. Therefore, with increasing transient temperature or prolonged pulse width, the outermost layer of NCM811 particles gradually transforms from a layered phase into spinel phases of varying thickness, and ultimately evolves into a rock-salt phase under severe lithium and oxygen extraction (Figure 1a).

High-resolution transmission electron microscopy (HRTEM) characterization provides direct evidence for the interfacial evolution of NCM811 under varying thermal pulse treatment. As disclosed in Figure 1b, the pristine sample shows a consistent R-3m layered phase structure (labeled as L in the box), characterized by a 0.47 nm lattice spacing, from the surface to the interior. As the temperature increases under a fixed pulse width (30 s), surface lithium atoms are gradually thermally stripped. When the temperature reaches 300 °C, the layered structure at the interface begins to transform into the spinel phase (labeled as S in the box, Figure 1c). Meanwhile, below 400 °C, the thickness of the spinel-phase skin layer increases as the depth of lithium atom extraction grows, expanding from a few nanometers to several tens of nanometers (Figure 1d and Supplementary Fig. 2). We refer to the samples treated at different temperatures (X °C) as FJX, where FJ stands for Flash Joule heating. Taking the comparison between the pristine sample and FJ350 as an example, scanning electron microscopy (SEM) images show negligible difference in their microscopic morphology (Supplementary Fig. 3). However, the electron paramagnetic resonance[33] (EPR) spectra at around 3500 G reveal that the FJ350 sample contains more oxygen vacancies. This indicates that although the layered-to-spinel transition theoretically involves only lithium stripping, it is also accompanied by a certain degree of thermal oxygen extraction during the actual flash sintering process (Supplementary Fig. 4). In addition, the O 1s X-ray photoelectron spectroscopy (XPS) spectrum of FJ350 also exhibits a more pronounced oxygen defect peak[32,33] (Odef), which can likewise be attributed to oxygen vacancies triggered by interfacial oxygen extraction (Supplementary Fig. 5). As the temperature exceeds 500 °C, further delithiation and deoxygenation triggers the transition of the interfacial layer to the rock-salt phase (labeled as R in the box, Figure 1e), with the thickness increasing as the temperature rises (e.g., 600 °C, Supplementary Fig. 6). Ex situ X-ray diffraction (XRD) patterns demonstrate that under thermal pulses with a 30 s pulse width, a distinct rock-salt phase peak only emerges when the temperature reaches 700 °C (i.e., involving changes in the bulk phase, Supplementary Fig. 7), confirming that Li and O stripping at lower temperatures predominantly occurs in the surface/subsurface regions of the material.

In addition to ex situ HRTEM, we further validated the temperature-dependent interfacial phase correlations via in situ TEM equipped with microchip-based heating devices, confirming that 300 and 450 °C are the critical temperatures for the formation of spinel phase and rock-salt phase, respectively (Supplementary Fig. 8). From another perspective, by extending the pulse width at a fixed temperature, a continuous change in the surface phase structure of NCM811 can also be observed. As displayed in Figures 1f-i, after 60 s at 350 °C, a transformation from the spinel phase to the rock-salt phase occurs. Note that although interfacial transformations occur, the XRD patterns show negligible changes even when the processing durations exceed 300 s, indicating that transient high-temperature treatments have minimal impact on the bulk crystalline phases of the material, particularly under low-temperature conditions (Supplementary Fig. 9). Interestingly, however, after 30 s of treatment at 350 °C, the calculated I(003)/I(104) intensity ratio reaches its maximum value, corresponding to the lowest degree of Li+/Ni2+ cation mixing (Supplementary Figs. 7 and 9). This demonstrates that the removal of Li and O at the interface still exerts some influence on the bulk crystalline structure of NCM811, particularly affecting the cation mixing, a parameter closely related to the ICE and long-term cycling stability.

It is crucial to emphasize that precisely controlled transient thermal pulse treatment enables the acquisition of numerous kinetically controlled metastable intermediate states, which serve as the key technical foundation for achieving customization and thickness regulation of the NCM811 interfacial phases. In stark contrast, under traditional steady-state thermal sintering conditions (typically lasting several hours at least), these fleeting interfacial phases cannot be captured or preserved, ultimately yielding only thermodynamically controlled rock-salt phase structures (e.g., products at 350 °C for 0.5 h, Supplementary Fig. 10). Undoubtedly, the unique fast heating and rapid cooling mechanisms of the Joule heating technique provide a crucial technical prerequisite for achieving nano-scale precise phase transformation at the electrode material interface.

Spinel-phase skin enables highly reversible NCM811 positive electrodes

The effect of the phase structure and thickness of the skin shell on the ICE, rate capability, and cycling performance of NCM811 is evaluated in detail through half-cell tests. Figures 2a and 2b present the charge-discharge profiles at 0.1 C (1 C=180 mA-1) and corresponding ICE of the NCM811 before and after 30 s of treatment at various temperatures. Apparently, the reversible discharge capacity and ICE of NCM811 both show a trend of first increasing and then decreasing with the rise in pulse temperature, reaching their maximum at 350 °C. Notably, the FJ350 sample discloses an high ICE of 95.3%, surpassing that of the pristine sample (89.6%). This is highly consistent with the lowest degree of Li⁺/ Ni2+ cation mixing observed in FJ350 (Supplementary Figs. 7 and 9). More effectively, the FJ350 also demonstrates both the highest rate performance and cycling stability at 1 C (Figure 2c and Supplementary Fig. 11). Moreover, the slopes obtained from cyclic voltammetry profiles at varying scan rates also reveal that FJ350 possesses superior ionic transport properties (Supplementary Fig. 12). This convincingly indicates that a spinel-phase skin with an appropriate thickness can ensure high-rate and isotropic Li+ extraction/insertion across the interface during long-term cycling. In contrast, samples with either too thick or too thin spinel-phase skin coatings (FJ300 and FJ400), although performing better than the pristine sample, are still inferior to FJ350. On one hand, the excessively thin shell lacks sufficient mechanical strength and isolation effect (e.g., incomplete coverage), only partially suppressing lattice oxygen loss and volume effects. On the other hand, the overly thick shell may hinder Li+ transport owing to the raised ionic diffusion barrier, causing adverse ion blockage and polarization. In situ differential electrochemical mass spectrometry (DEMS) was employed to monitor the oxygen evolution of different samples under high potential (e.g., 4.7 V) in order to evaluate the coverage of the spinel-phase shell. As Supplementary Fig. 13 demonstrates, the FJ300 sample with a thin and incomplete spinel-phase shell already exhibits a certain degree of suppression of oxygen evolution during the first charge compared with the pristine sample, but the reduction is limited, decreasing only from 4.6 to 3.1 nmol min-1. In sharp contrast, the oxygen-release signals of the FJ350 and FJ400 electrodes are significantly lower, with FJ350 in particular reduced to as low as 1.2 nmol min-1. This result indicates that the spinel-phase shell can effectively inhibit oxygen evolution even under the high cutoff potential of 4.7 V, directly reflecting both the completeness and the strong spatial confinement effect of the FJ350 coating layer. The suppression of oxygen release is also consistent with the increased Coulombic efficiency and reduced irreversible lithium consumption observed in the electrochemical tests, further demonstrating the critical role of an optimized coating thickness in achieving stable cycling performance. Moreover, samples subjected to higher temperatures, particularly those at 500 °C and above (e.g., FJ500 and FJ600), exhibit a significant reduction in ICE and initial discharge capacity. In particular, the FJ600 sample delivers the lowest discharge capacity of 154.5 mAh g-1, corresponding to a low ICE of 76.0%. According to the semi-quantitative estimation of FJ600 in Supplementary Fig. 14 and Table S1, this deterioration primarily arises from the thick rock-salt reconstruction layer, which blocks Li⁺ transport and exacerbates interfacial deterioration, rather than from the negligible lithium extraction during the flash Joule heating process.

Similarly, samples treated at a fixed temperature of 350 °C for different pulse widths also exhibit electrochemical performance that first increases and then decreases. As manifested in Figures 2d and 2e, 60 s is both the point at which the interfacial phase transitions from spinel to rock-salt and the turning point for the deterioration of electrochemical performance. The sample treated for 30 s, namely FJ350, possesses the most suitable spinel-phase shell thickness, stands out with the highest ICE and optimal cycling stability among all samples (Supplementary Fig. 15).

Based on the TEM analysis, we performed semi-quantitative measurements and statistical analysis of the average thickness of the spinel-phase shell formed on the surface of NCM811 under different flash-sintering conditions (Supplementary Fig. 16), and correlated these results with the corresponding electrochemical performance. As summarized in Table S2, the thickness of the spinel-phase layer formed under different treatment conditions (pulse temperature and pulse width) varies systematically, generally following the trend of increasing thickness with longer pulse duration or higher treatment temperature. Meanwhile, the thickness of the spinel-phase shell, as well as the associated degree of lithium deficiency, exerts a pronounced influence on the electrochemical performance including ICEs and long-term cycling stability. These findings further demonstrate that an optimally engineered spinel layer, with well-balanced thickness and defect chemistry, enables the FJ350 sample to deliver the most superior electrochemical performance among all treated conditions. As Figure 2f demonstrates, the FJ350 electrode preserves a discharge capacity of 167.1 mAh g-1 with a retention of 93.6% after 500 cycles at 1 C, far exceeding the 115 mAh g-1 and 65.2% retention observed for the pristine sample. Charge-discharge curves convincingly demonstrate the material advantages of FJ350 in suppressing capacity fading and resisting electrode polarization (Supplementary Fig. 17). The FJ350 electrode exhibits a much slower potential plateau decay, with a drop of only 0.102 mV per cycle, in stark contrast to the pristine electrode, which presents a sharp decay of 0.589 mV per cycle (Supplementary Fig. 18). Apparently, by tightly confining the oxidized lattice oxygen and inhibiting the oxygen loss, the spinel-phase skin-encapsulated FJ350 maintains the average oxidation state of the transition metals in the bulk structure, thereby preserving the operating potential and specific capacity during extended cycling. This can be further confirmed by the differential capacity (dQ/dV vs. potential) analysis. As depicted in Figure 2g, the H2-H3 phase transition around 4.2 V, associated with c-axis contraction and lattice oxygen oxidation, is significantly diminished in the pristine sample after 500 cycles, indicating severe structural changes and lattice oxygen loss. In sharp contrast, the FJ350 sample maintains a stable H2-H3 redox peak, suggesting reinforced structural integrity and effective stabilization of lattice oxygen, thereby greatly suppressing its release.

The advantage of FJ350 over the pristine sample in reaction kinetics can be elucidated by measuring the Li-ion diffusion coefficient (DLi+). The galvanostatic intermittent titration technique (GITT) profile of the charge/discharge processes of the pristine and FJ350 electrodes and the corresponding DLi+ vs. potential curve indicate that the cycled FJ350 electrode delivers a faster Li+ transport rate (Figure 2h and Supplementary Fig. 19). Benefiting from the reaction kinetics, the FJ350 electrode retains 81.3% of its capacity after 1000 cycles at a high current density of 2 C. By contrast, the pristine sample suffers rapid capacity degradation, maintaining only 60.3% of its initial capacity after long-term cycling (Supplementary Fig. 20). More encouragingly, the FJ350 electrode demonstrates prominent performance under more harsh operating conditions, such as a potential upper limit of 4.5 V (Supplementary Fig. 21) and a high temperature of 60 °C (Supplementary Fig. 22). As a demonstration of practical application, a pouch full cell with a FJ350 positive electrode and graphite negative electrode was successfully assembled and used to power a smartphone (Figure 2i). More effectively, within the potential window of 2.7-4.2 V, the FJ350-based pouch cell maintains 80.1% capacity retention after 2000 cycles (Figure 2j), significantly surpassing that of the pristine NCM811-based pouch cell (52.0%). The charge-discharge curves further demonstrate the advantage of the FJ350 electrode in mitigating capacity fade and resisting electrode polarization, highlighting its long-term cycling stability (Supplementary Fig. 23).

Notably, half-cells and pouch cells incorporating the optimized FJ350 exhibit comparable electrochemical performance, especially in cycling stability, compared to counterparts utilizing conventional NCM electrodes reported in the reported literature (Supplementary Fig. 24 and Supplementary Table S3). To further highlight the electrochemical advantages of the self-induced coated FJ350, we conducted comparative studies on the electrochemical performance and DEMS results of FJ350, commercial LiNbO3-coated LiNi0.8Co0.1Mn0.1O2 (LiNi0.8Co0.1Mn0.1O2@LiNbO3), and commercial Li3BO3-coated LiNi0.8Co0.1Mn0.1O2 (LiNi0.8Co0.1Mn0.1O2@B). As shown in Supplementary Figs. 25 and 26, the commercial NCM811 positive electrodes coated with LiNbO3 or Li3BO3 exhibit not only lower initial coulombic efficiencies and poorer cycling stability but also inferior suppression of oxygen release under high potential (4.7 V). In addition, we further compared FJ350 with NCM811 electrodes coated using traditional wet-chemical methods and other representative coating approaches reported in the literature. As presented in Supplementary Table S4, the FJ350 electrode demonstrates comparable improvements in both cycling performance and rate capability. These results collectively highlight that our subtraction-coating strategy not only enhances electrochemical performance but also improves battery stability and extends service life, making it the most effective approach among currently available coating methods.

Samples annealed at 350 °C under an argon atmosphere using a traditional tube furnace (Supplementary Fig. 10) were compared with FJ350. The results show slight improvement compared to the pristine counterpart, but still far inferior to FJ350 (Supplementary Fig. 27). This highlights the inherent limitations of conventional annealing methods, where tube furnace processing fails to capture and preserve the beneficial spinel-phase skin, instead directly leading to rock-salt shell that degrades electrochemical performance. Undoubtedly, transient high-temperature treatment provides a unique and disruptive technological solution for the customized modulation of the interfacial phase structure of NCM811 electrodes. More importantly, this precise interfacial regulation strategy, namely, the transformation from a layered structure to a spinel shell at the interface (Supplementary Fig. 28 and Fig. 29), enhances the reversible capacity and cycling stability of various materials, including polycrystalline LiNi0.825Co0.115Mn0.06O2, polycrystalline LiNi0.9Co0.05Mn0.05O2 (NCM90), and single-crystal LiNi0.8Co0.1Mn0.1O2 (SC-NCM811). These results (Supplementary Fig. 30) highlight the broad applicability of this strategy in tailoring surface atomic composition and stabilizing interfacial phases.

Interfacial and bulk-phase stability during the first cycle

During cycling, delithiation-induced lattice distortion and stress accumulation, along with oxygen loss and cation mixing, can trigger a cascade of phase transitions in the surface layer of the pristine NCM811 electrodes from layered to disordered spinel and rock-salt, leading to continuous deterioration of electrochemical performance. This trend also gradually propagates from the interface into the bulk, resulting in continuous structural degradation of the material. As emphasized in Figure 3a, the outer layer of the pristine NCM811 begins to exhibit a NiO-like rock-salt phase[35] only after the first charge to 4.3 V. Raising the upper limit of charging potential (e.g., 4.7 V) exacerbates this process, as the increased lithium and oxygen vacancy concentrations promote the migration of transition metals, thereby thickening the rock-salt phase transformation (Supplementary Fig. 31). In stark contrast, regardless of the delithiated state at 4.3 V or 4.7 V, the spinel-phase skin on the FJ350 surface stabilizes the three-dimensional framework (Figure 3b and Supplementary Fig. 32), highlighting its high-potential compatibility and resistance to mechanical deformation. This also indicates that the highly crystalline spinel-phase shell obtained through flash sintering exhibits robust electrochemical stability and does not further transform into the rock-salt phase. This can be further confirmed by measuring the Li distribution in NCM811 after its first charge to 4.3 V. Time-of-flight secondary ion mass spectrometry[36] (TOF-SIMS) analysis reveals a heterogeneous distribution of 7Li within the pristine particle (Figure 3c), which is attributed to the formation of an interfacial inactive rock-salt phase that impedes isotropic Li+ extraction across the surface, leading to local charge heterogeneity[9,11], internal structural distortion and localized damage. By contrast, the FJ350 exhibits a uniform 7Li distribution across its entire particle, attributed to low oxygen loss, well-preserved structural integrity and isotropic high-rate Li-ion diffusion channels at the interface (Figure 3d). Given the above, the coating layer stabilizes the interface by preventing the formation of inactive phases, such as the rock-salt phase. This interfacial stability ensures efficient Li+ transport across the interface and suppresses lattice oxygen loss, thereby enhancing the overall Li⁺ diffusion within the particles. These findings highlight the critical role of the coating in optimizing both interfacial and bulk properties, leading to significantly improved electrochemical performance.

In situ XRD spectra coupled with Rietveld refinements further underscore the reinforcing effect of the spinel-phase skin on the structural stability of NCM811. During the first charge process of the pristine electrode, the (003) peak shifts towards lower angles due to the increased repulsion between adjacent oxygen planes[37,38], causing the c-axis lattice parameter to expand during charging (Figure 3e). As the potential approaches 4.3 V, the (003) peak shifts to higher angles, indicating compression of the c parameter at a high delithiation state. During the entire charge process, the (003) peak shifts by 0.42°. During the discharge process, the (003) peak cannot recover to its initial position and ultimately shifts negatively by 0.09°, indicating substantial damage to the crystal structure. In contrast, the FJ350 electrode exhibits a 0.28° shift during the charge process and fully returns to its original peak position after discharging, indicating significantly enhanced structural reversibility (Figure 3f).

The calculated cell parameters as a function of charge potential are presented in Figure 3g and Supplementary Fig. 33. The initial expansion of the c-axis is due to increased electrostatic repulsion between oxygen layers as Li⁺ ions are extracted. At higher states of charge, the contraction along the c-axis arises from electron withdrawal by highly oxidized transition-metal ions from the coordinated oxygen atoms, leading to oxygen ion oxidation and a consequent reduction in the repulsive interactions within the delithiated Li layers. Meanwhile, the lattice volume decreases continuously due to dominant in-plane contraction from the shortening of TM-O bonds as the transition metals are oxidized. Furthermore, for the FJ350 sample, the restrained H2-H3 phase mitigates lattice-induced stress accumulation, as the substantial c-axis mismatch between the H2 and H3 phases constitutes a dominant source of mechanical strain during cycling. Quantitatively, the changes in c-axis length, a-axis length and total volume of the FJ350 sample during charging are 2.12%, 1.52% and 2.85%, significantly lower than that of the pristine electrode (4.31%, 2.15% and 4.63%). This clearly indicates that the spinel-phase skin indeed helps significantly reduce the lattice anisotropic changes exacerbated by oxygen loss during the H2-H3 phase transition.

To gain deeper insight into how the spinel-phase skin influences the structural evolution and Li-ion transport behavior of NCM811 under more demanding conditions, in situ XRD analysis was additionally carried out at an upper cutoff potential of 4.7 V. As illustrated in Figure 3h, when the pristine electrode is charged from 4.3 V to 4.7 V, the intensity of the (003) peak significantly weakens, and its shape becomes asymmetrical, indicating the coexistence of H2 and H3 phases and highlighting the dramatic structural changes during the H2-H3 transition[39]. Even worse, at the fully charged state of 4.7 V, the residue of the H2 phase is still quite evident, revealing that the transition from H2 to H3 has not been fully completed (Figure 3j). In contrast, only a slight H2 phase is observed at 4.3 V in the FJ350 sample, and as the charging potential increases (e.g., 4.5 V, Figures 3i and 3k), the phase transition is fully completed[21,40]. The coexistence of the H2 and H3 phases is likely to lead to lattice mismatch within the particles, causing uneven lithium distribution, which is highly consistent with the TOF-SIMS observations (Figure 3c). Apparently, the high mechanical strength and rapid reaction kinetics of the spinel-phase skin-encapsulated FJ350 play a key role in promoting the complete transformation of H2-H3. Differential Scanning calorimetry[41] (DSC) tests further emphasize the role of the spinel-phase skin in stabilizing the FJ350 structure and suppressing oxygen evolution ( Supplementary Fig. 34), primarily reflected in a 32 °C delay in the oxygen release peak compared to the pristine sample, with corresponding heating powers of 705 J g-1 and 1013 J g-1, respectively.

Crack- and corrosion-resistant mechanisms during cycling

The microscopic mechanism by which the spinel-phase skin significantly improves the long-term cycling stability of NCM811 has also been deeply studied. As depicted in Figures 4a, the cross-sectional SEM image clearly indicates that after 500 cycles, severe cracking and damage appear inside the pristine particle in the discharged state. The enhanced bulk stability of FJ350 can be attributed to interfacial electron transfer between the spinel-phase shell and the bulk Ni-rich lattice, which mitigates c-axis variation and confines lattice oxygen under highly delithiated conditions. Consequently, this process alleviates lattice collapse induced by oxygen release from the surface to the bulk. This undoubtedly accelerates parasitic reactions with the electrolyte, promotes the dissolution and migration of transition metals, and significantly thickens the CEI (Supplementary Fig. 35). In stark contrast, only a small number of stress-induced micropores are observed inside the cycled FJ350 (Figure 4b), rather than cracks, while a dense and thin CEI film forms on the interface, and the overall structure remains well preserved The internal cracking and interface deterioration of the cycled pristine and FJ350 electrodes can be semi-quantitatively compared through the electrochemical impedance spectroscopy (EIS) characterization. As manifested in Supplementary Fig. 36 and Table S5 , the FJ350 electrode exhibits markedly lower CEI film resistance (Rf) and charge transfer resistance (Rct) at all stages of cycling[42]. The pronounced reduction in Rct confirm that the spinel-phase skin establishes a percolating electron pathway, synergizing with fast Li+ transport to achieve high rate capability and cycling capacity retention. Further insights into the surface microstructure after cycling were compared by HRTEM analysis and corresponding FFT and IFFT images. As manifested in Supplementary Fig. 37, a 15 nm thick NiO-like rock-salt phase formed on the cycled pristine electrode surface owing to the continuous side reactions with the electrolyte[15]. In contrast, the FJ350 particle perfectly preserved the spinel-phase skin-encapsulated layered structure even after 500 cycles (Figure 4c), again confirming the electrochemical stability and mechanical robustness of the flash sintering-induced highly crystalline spinel-phase shell. High angular annular dark field (HAADF) images of the cycled pristine and FJ350 electrodes suggest that the spinel-phase skin adapts well to lattice contraction and expansion and prevents the formation of microcracks (Supplementary Fig. 38), thereby suppressing the interfacial side reactions and the growth of the rock-salt phase and preserving the structural integrity of the NCM811 particles during cycling (Supplementary Fig. 39).

Depth-profile XPS analysis of the cycled electrodes reveals the CEI components arising from the decomposition of the lithium salts and solvents (Figures 4d and 4e). The F 1s XPS peaks[18,33] indicate the presence of LiF, C-F and LiPFxOy in the FJ350 sample throughout the cycling process, with a significantly higher LiF content compared to the pristine sample. The higher LiF content in the CEI can largely suppress the decomposition of the electrolyte and LiPF6, contributing to long-term stability. The O 1s spectra[43] consists of four main peaks related to C-O, C=O, P-O and lattice O, respectively. The retention of the lattice O signal at the FJ350 interface validates the effectiveness of the spinel-phase shell in suppressing outer-layer oxygen loss. Moreover, the significantly reduced intensities of the C=O and P-O peaks in the FJ350 sample indicate that the LiF-dominated CEI effectively suppresses the decomposition of solvents (ethylene carbonate and diethyl carbonate) and LiPF6 salt. The LiF-rich CEI not only passivates the interface but also exhibits moderate electronic conductivity, balancing electron blocking for side reaction suppression and sufficient charge transfer for redox reactions. This can be confirmed by the low intensities of CO3 2- and C=O peaks from ROCO2Li in the C 1s spectrum[16,18] of the FJ350 sample. The C 1s spectrum further reveals a stronger signal from the organic polymer C-C backbone in FJ350. The polymer components not only enhance the mechanical flexibility and strength of the CEI but also effectively embed the inorganic LiF, contributing to the formation of a thinner and more robust CEI layer. These results strongly indicate that the spinel-phase skin serves as a physical barrier between the highly delithiated cathode and the electrolyte, effectively inhibiting solvent and LiPF6 decomposition as well as active material dissolution, while promoting the formation of an LiF-rich CEI layer, thereby preventing severe structural degradation and interfacial side reactions.

TOF-SIMS analysis further corroborated the above results by revealing the three-dimensional structure and compositional distribution of the CEI layer[40,44] (Figures 4f, Supplementary Figs. 40 and 41). The following lists several major ion fragments, including C2H3OF- (organic species from solvent decomposition), LiF2- and PO2F2- (inorganic species from the LiPF6 decomposition), NiF3-, CoF3- and MnF3- (dissolved products of active materials). The trend of secondary ion distribution with sputtering time can be derived from the normalized intensity plot in Figure 4f. Apparently, a significantly thinner CEI layer is formed on the surface of the cycled FJ350 electrode. Moreover, the concentrations of both organic and inorganic species (e.g., C2H3OF-, MnF3-, NiF3- and CoF3-) generated from interfacial degradation are markedly reduced on the FJ350 surface, confirming that electrolyte decomposition and metal cation dissolution are greatly suppressed. This interfacial stabilization strategy demonstrates universal effectiveness across different high-Ni positive electrode materials, as evidenced by the TOF-SIMS analyses of LiNi0.9Co0.05Mn0.05O2 (Supplementary Fig. 42) and single-crystal LiNi0.8Co0.1Mn0.1O2 electrodes (Supplementary Fig. 43). Both systems exhibit a thinner CEI layer and markedly reduced fragment intensities, confirming the cross-system applicability of this approach.

Figures 4g and 4h present the X-ray absorption near-edge structure (XANES) of Ni in the pristine and FJ350 samples after 500 cycles. The Ni-K edge of cycled FJ350 nearly overlaps with that of the pristine sample, demonstrating improved stability in the oxidation state of nickel throughout the cycling process[40,45-47]. By contrast, a pronounced absorption edge shift toward higher energy is observed in the cycled pristine sample, implying incomplete reduction (i.e., irreversible Ni redox behavior) the NCM811 material. This phenomenon can be attributed to the formation of the rock-salt-type NiO phase at the interface and within the bulk during cycling, which impedes Li+/electron transport and causes localized deactivation. And the underlying cause can be traced back to lattice oxygen loss[45,46]. Fourier-transformed extended X-ray absorption fine structures (FT-EXAFS) analysis at the Ni-K edge reveals that, after prolonged cycling, the Ni-O and Ni-transition metal (e.g., Co, Mn) coordination distances are noticeably shortened in the pristine sample (Figure 4i), indicative of lattice contraction, accumulated structural stress, and oxygen vacancy formation. In contrast, the variation in interatomic distances in the cycled FJ350 electrode is nearly negligible, highlighting a highly stable local atomic coordination environment (Figure 4j). Moreover, the wavelet-transformed EXAFS analysis in Figures 4k and 4l reveals a substantial reduction in signal intensity in the high-K region for the pristine sample after 500 cycles, pointing to the degradation of Ni-transition metal coordination and increased structural disorder. In contrast, the cycled FJ350 (Figures 4m and 4n) retains strong signals in this region, suggesting that the spinel coating layer effectively mitigates coordination degradation and preserves the long-range structural order of the material.

Density functional theory calculations

Based on the structural model presented in Supplementary Fig. 44 (Supplementary Data 1), we performed a comparative analysis of the partial density of states (PDOS) for both pristine LiTMO2 (TM = Ni0.8Co0.1Mn0.1), representing the bulk layered phase, and the spinel-coated LiTMO2@LiTM2O4 under charged and discharged states[48]. The PDOS results reveal a notable increase in the electronic density near the Fermi level for the coated material, along with a narrowed band gap and a reduced electronic transition barrier (Figure 5a). These electronic structure modifications suggest a marked enhancement in electrical conductivity, which is expected to improve the rate performance relative to the uncoated sample. To further understand the redox behavior and structural stability, we analyzed the crystal orbital Hamilton population (COHP) of Ni-O bonds (Supplementary Fig. 45). The integrated COHP (-ICOHP) values indicate that in the fully lithiated state, the Ni-O bonds in LiTMO2@LiTM2O4 possess a stronger bonding strength (-ICOHP = 0.803 eV) than those in pristine LiTMO2 (-ICOHP = 0.761 eV). In the delithiated state, both TMO2 and TMO2@TM2O4 show comparable Ni-O bond strengths. Given that LiTM2O4 generally retains some Li+ during cycling, the LiTMO2@LiTM2O4 heterostructure helps maintain robust bond stability throughout charge and discharge, thereby reinforcing the host framework and suppressing oxygen loss as well as phase degradation. This effect contributes to the inhibition of irreversible oxygen and nickel redox activity. In contrast, the pristine sample exhibits a pronounced decrease in the electronic density of states near the Fermi level under high degrees of delithiation, along with lower -ICOHP values in the lithiated state. These features imply inferior electronic transport and weaker Ni-O bonding during cycling, which can trigger oxygen release and structural collapse, ultimately resulting in rapid capacity fading. These theoretical insights are consistent with experimental observations, i.e., the spinel-coated LiTMO2@LiTM2O4 maintains high structural stability even under high-potential conditions, effectively suppressing oxygen release and irreversible phase transitions.

During the lithiation/delithiation process, the system expansion rates of LiTMO2 and LiTM2O4 reach 10.6% and 9.0%, respectively. However, after composite treatment, the volume change of the material significantly reduces to 3.5% (Figure 5b). This improvement can be attributed to the strong covalent interaction of Ni-O bonds at the spinel/LiTMO2 interface, effectively suppressing the volume effects induced by lithium ions. Notably, after coating formation, the lattice of LiTMO2 expands while that of LiTM2O4 contracts (Figure 5c). This mortise-and-tenon like locking mechanism significantly reduces lattice distortion during cycling, providing reliable structural support for enhancing the cycling life of the positive material. This is consistent with our experimental observation that FJ350 exhibits significantly lower crystal volume change (Figure 3g) and higher structural stability (Figure 4b).

Based on the analysis of average lithium ion adsorption energy in LiTMO2 and LiTM2O4, the spinel phase exhibits a significantly stronger adsorption capability than LiTMO2 throughout the charge/discharge process, especially under lithium-rich conditions (Figure 5d). This enhanced adsorption behavior is closely related to Li⁺ diffusion kinetics. In line with established studies on layered oxide electrodes[49,50], we employed the climbing image nudged elastic band method (CI-NEB) method to evaluate the tetrahedral site hop (TSH) mechanism, which is the dominant Li+ diffusion pathway characterized by a relatively low activation energy. The computed diffusion barrier for TMO2 is approximately 0.3 eV, consistent with reported values. For clarity, schematic illustrations of the TSH pathway from both side and top views are provided. As shown in Figure 5e, in the high-delithiation state, the LiTMO2 region in the coated sample shows improved lithium adsorption compared to the uncoated material. However, when the spinel coating is in a lithium-rich state, this enhancement diminishes with increasing lithium content. This implies that an excessively thick LiTMO2 outer layer could impede Li+ transport into the bulk and weaken the adsorption contribution from LiTMO2. Therefore, an optimal thickness of the spinel coating is essential to enhance lithium adsorption (Figure 5d) while establishing an efficient Li+ diffusion channel (Figure 5e). Moreover, Li+ diffusion across the coating interface exhibits strongly asymmetric energy barriers (Figure 5f and 5g), suggesting that extraction of Li⁺ from LiTMO2 is energetically more favorable than insertion. Hence, rational design of the thickness and microstructure of the LiTM2O4 spinel layer is critical to improve structural stability while minimizing adverse effects on lithium-ion transport kinetics. These findings are consistent with the high ICE, rate capability, and cycling performance observed in sample FJ350, which possesses an optimally tuned spinel-phase thickness.

Another significant advantage of spinel LiTM2O4 coating on LiTMO2 lies in its enhancement of surface structural stability. As shown in Figure 5h, TM2O4 exhibits a lower Fermi level and higher work function compared to TMO2, indicating that electrons reside at lower energy levels and are less likely to be lost, thereby reducing the probability of adverse side reactions. By comparing the behavior of lithium ions at different concentrations entering LiTM2O4 and LiTMO2 (Figures 5i and 5k, and Supplementary Fig. 46a), the outer and subouter layer oxygen atoms of LiTMO2 lose stability, detach from the lattice, and form clusters with moving lithium ions, which becomes more pronounced with increasing lithium ion concentration. This may increase the risk of forming active oxygen species during charge/discharge processes, leading to crystal structure instability, oxygen evolution, phase transitions, and aggravated side reactions with the electrolyte, ultimately resulting in significant capacity reduction and cycle performance deterioration. In comparison, the surface and subsurface oxygen atoms of spinel LiTM2O4 exhibit structural stability under both low and high lithium concentrations. Particularly at low lithium concentrations, the surface oxygen atoms only display lattice distortion (Figures 5j and 5l). Thus, the spinel LiTM2O4 coating (Supplementary Fig. 46b) effectively enhances LiTMO2 stability by inhibiting oxygen release and passivating active sites. Moreover, the stable surface can avoid adverse interactions with the electrolyte, significantly improving the overall stability of the electrode material, further demonstrating its potential as a high-performance positive electrode material.

Methods
Synthesis: All NCM positive electrode materials were used directly without further pre-treatment. Polycrystalline LiNi0.8Co0.1Mn0.1O2, polycrystalline LiNi0.825Co0.115Mn0.06O2, polycrystalline LiNi0.9Co0.05Mn0.05O2 and single-crystal LiNi0.8Co0.1Mn0.1O2, were purchased directly from Xiamen XTC Intelligent Equipment Technology Inc. (>99.5%, battery grade). Flash sintering was conducted using a self-built Joule heating system equipped with a temperature control system and a circulating cooling system. Prior to sintering, 0.5~2 g of the NCM sample was enclosed in the central cavity formed by two pieces of carbon felt (100 mm × 100 mm), with both ends of the felt connected to the positive and negative poles of the heating clamp. After sealing the reaction chamber, three vacuum cycles were performed to remove residual air and ensure an oxygen-free environment for subsequent treatment. The interfacial structure of the NCM materials was tuned by adjusting the treatment temperature and pulse duration. The reaction chamber was evacuated three times to remove residual air, ensuring an oxygen-free environment for the subsequent treatment. The customized modulation of the interfacial structure of the NCM material was primarily achieved by tuning the temperature and pulse width of the transient thermal pulses. The treatment temperatures ranged from 300 °C to 800 °C, with pulse durations from 15 to 300 seconds. For comparison, samples were also prepared and tested after conventional heating in a tube furnace at 350 °C for 0.5 to 3 hours.

Characterization: The crystal structure of the samples was analyzed by X-ray diffraction (XRD, Thermo Scientific ARL EQUINOX 3500) with Cu-Kα radiation using a scanning rate of 10° min-1 and a step size of 0.02°. For in situ XRD, the slurry was coated onto carbon paper, dried in a vacuum oven, and then assembled into a specially designed Bruker chamber for testing and then the resulting data was fitted by Rietveld refinement (GSAS software). The morphology of the materials before and after cycling was characterized using scanning electron microscopy (SEM, Gemini SEM 500, Zeiss), with the samples directly mounted on conductive adhesive for testing and operated at a beam voltage of 15 kV. Oxygen vacancies were investigated by Electron Paramagnetic Resonance measurement (EPR, Bruker EMXplus-9.5/12), Measurements were performed in the X-band (9.830243 GHz) with a modulation amplitude of 1 G, microwave power of 2 mW and modulation frequency of 100 kHz. The conversion time was set to 60 ms. The interfacial microstructure of pristine, flash-sintered, charged, and cycled NCM samples was characterized using transmission electron microscopy (TEM, FEI Talos F200s) with a beam voltage of 200 kV. Additionally, in situ TEM heating was performed with a Gatan heating holder. The temperature was ramped from room temperature to 450 °C at a rate of 1 °C s-1, with measurements conducted at various intermediate temperatures. The sample was held at each temperature for approximately 10 min. In situ differential electrochemical mass spectrometry (DEMS, Hiden HPR-20) studies were investigated using a specialized cell that replaced the positive and negative button cell casings with an in situ cell. The positive and negative materials, electrolyte, gaskets, and shrapnel selected are identical to those of the button cell. To ensure gas generation during the charging process, each aluminium foil piece with a diameter of 14 mm was loaded with approximately 10~14 mg of active material. For post-cycling analyses, the cycled cells were disassembled and the electrodes were washed with dimethyl carbonate (DMC) and dried in an Ar-filled glovebox. Cross-sectional samples of the cycled particles were prepared using a triple ion beam cutting machine (Leica EM TIC 3X) and subsequently examined by SEM to examine the internal structure, as well as time-of-flight secondary ion mass spectroscopy (TOF-SIMS, IONTOF M6) to investigate the distribution of 7Li. The thermal stability of the positive electrode material was assessed using differential scanning calorimetry (DSC, STA 409 PC, NETZSCH) with a heating speed of 5 °C min-1 under an Ar atmosphere. The cycled electrodes surface chemical compositions were identified by TOF-SIMS using a Bi3+ primary ion source operated at 30 keV with a beam current of 0.5 pA. X-ray photoelectron spectroscopy with depth profiling (XPS, Thermo Fisher, Escalab Xi+) was performed using Ar⁺ ion sputtering to etch the electrode surface with a sputtering energy of 2000 eV, a sputtering rate of 0.29 nm s-1 and sputtering depths of 0, 5, 10 and 15 nm. The cycled electrodes were handled in an argon-filled glovebox due to their air sensitivity. After electrochemical cycling, the coin cells were disassembled, and the electrodes were rinsed twice with dimethyl carbonate to remove residual electrolyte. The cleaned electrodes were then dried under an inert atmosphere prior to analysis.

Electrochemical measurements: The NCM positive electrodes were fabricated using a composite material consisting of 90 wt.% active material, 5 wt.% Super P (battery grade, Shenzhen Kejing Co., Ltd.) and 5 wt.% poly(vinylidene fluoride) (PVDF, 99.9%, Sigma-Aldrich) dissolved in N-methyl-2-pyrroidone (99.9%, Canrd Technology Co., Ltd.). The obtained slurry was casted using an automatic coating machine (MSK-AFA-SC200, Shenzhen Kejing Co., Ltd.) onto carbon-coated aluminum foil (thickness of 15 μm, Shenzhen Kejing Co., Ltd.) and dried at 120 °C for 4 h in a vacuum oven. The electrodes were cut using electrode punching machine (MSK-T10, Shenzhen Kejing Co., Ltd.) into discs with a diameter of 11 mm, and the mass loading of the active material was approximately 3.5 mg cm-2. When assembling CR2032-type half cells (316 stainless steel), lithium metal foil (thickness: 1 mm, diameter: 15.8 mm, 99.9%, China Energy Lithium Co., Ltd.) was used as the negative electrode and Celgard 2500 as the separator (thickness: 25 μm, diameter: 19 mm, Canrd Technology Co., Ltd.). The spacer (thickness: 1 mm, diameter: 15.8 mm, Dodochem Technology Co., Ltd.) and spring (thickness: 1.1 mm, diameter: 15.4 mm, Dodochem Technology Co., Ltd.) are made of 316 stainless steel. The electrolyte (Dodochem Technology Co., Ltd.) was composed of 1 M LiPF6 dissolved in ethylene carbonate (EC) and diethyl carbonate (DEC) with a volume ratio of 3:7 and with 2 wt.% vinylene carbonate addition. The electrolyte amount used in the coin cells was 60 µL. Half cells were tested using the Land battery testing system (Wuhan Lanhe Electronics, CT200A) under standard conditions of 2.7-4.3 V at 26 °C, as well as under harsher conditions such as 4.5 V and 60 °C with temperature-controlled incubator. 3-5 batteries were tested in parallel, and their performances are highly consistent. For pouch-cell tests, the negative electrode was composed of 95 wt.% graphite, 1.2 wt.% Super P, 1.8 wt.% sodium carboxymethyl cellulose (99.5%, Shenzhen Kejing Co., Ltd.) and 2 wt.% styrene butadiene rubber (50%, Shenzhen Kejing Co., Ltd.) and then coated onto copper foil (thickness of 4.5 μm, Shenzhen Kejing Co., Ltd.). The negative-to-positive capacity ratio (N/P ratio) for the pouch cell is around 1.12. The electrolyte injection procedure is as follows: The electrolyte is injected with an electrolyte-to-capacity ratio (E/C) of 3.5, which refers to the ratio of the electrolyte weight to the cell capacity (g Ah-1). After injection, the pouch cell was sealed under vacuum and then soaked at 45 °C for 48 hours, followed by a formation cycle. After formation, the cell is aged at 45 °C for 24 hours. The gas in the pouch is then vented by cutting the gas bag and the cell was resealed under vacuum. Finally, the cell undergoes partial capacity measurement, followed by subsequent electrochemical testing. The pouch cell consisted of 4 layers and full cells (~200 mAh) were tested between 2.7-4.2 V at 1 C (1 C=180 mA g-1) using the Neware battery testing system (CT-4008T, Neware Technology Co., Ltd.) under an external pressure by gently clamping the cell with a pressure of 37.2 kPa. Pouch-cell cycling tests were performed under ambient conditions without temperature control. Cyclic voltammetry (CV) tests were carried out at scan rates of 0.2, 0.4, 0.6, and 0.8 mV s-1 within a potential window of 2.7-4.5 V using an electrochemical workstation (CHI 660E, Shanghai Chenhua). Galvanostatic intermittent titration technique (GITT) measurements were performed with a constant current for 600 s with data recorded every 10 s followed by 1800 s relaxation with the upper potential of 4.3 V and lower potential of 2.7 V, respectively. Electrochemical impedance spectroscopy (EIS) was conducted using an electrochemical workstation. The measurements were performed after two cycles at 0.1 C and at the 1st, 100th, 200th, 300th, 400th and 500th cycles at 1 C. Prior to EIS analysis, the cells were charged to 4.3 V under constant-potential mode for 3 h, and then impedance spectra were collected over a frequency range of 0.01 Hz to 105 Hz with a perturbation amplitude of 5 mV.

Theoretical methods: All spin-polarized calculations were conducted by using density functional theory (DFT) as implemented in the Vienna Ab initio Simulation Package[51-53] (VASP), with the projector-augmented wave method[54] (PAW) describing the ion-electron interaction. The valence electrons were taken into account using a plane wave basis set with a kinetic energy cut-off of 520 eV, which guarantees that the total energy was converged when it reached 1 meV per atom. The exchange-correlation interaction of electrons was treated within the generalized gradient approximation (GGA) with Perdew-Burke-Ernzerhof[55-57] (PBE) functional. Convergence thresholds of 10-6 eV for electronic steps and 10-2 eV Å-1 for ionic relaxation were employed. The van der Waals interaction was considered by the empirical correction of Grimme’s scheme (DFT-D3). The reciprocal space was sampled using a gamma-centered k-grid with a resolution of 0.04 in units of 2π Å-1. For single-point energy calculations, the k-grid resolution was refined to 0.02. Minimum energy pathways were calculated using the CI-NEB[58,59]. The path threshold for CI-NEB calculations was set to 0.02 eV Å-1, and the insertion points were determined by the displacement of the diffusion atom. The von Mises local shear was calculated to characterize the degree of distortion of surface atoms. Regarding our interface model, the spinel phase formed through in situ transformation maintains crystallographic continuity with the parent layered structure, particularly along the basal plane direction. We focused on the 4-coordinated facets as model surfaces since they represent vulnerable sites where phase transition likely initiates. Stabilizing these critical interfaces through our approach provides substantial improvement in overall performance. The interface model was constructed as a symmetric slab without vacuum layers[60,61], with thick atomic regions to bulk-like conditions. Full geometry optimization was performed without constraints to minimize artificial strain. Li-deficient structures were generated using the SQS method[62] to approximate random vacancy distributions, with consistent sampling applied for all comparative studies. The adsorption energy of Li+ (Eads) is defined as: Eads = (Etotal - ETMO2/TM2O4 - nELi) / n where Etotal, ETMO2/TM2O4, and ELi denote the total energy of TMO2 or TM2O4 with n Li+, the energy of TMO2 or TM2O4, and the energy of one Li atom, respectively.

Peer review rounds
[{"round": 1, "reviewer_comments": {"reviewer_1": "In this work, Yang et al proposes a generalizable and scalable strategy for constructing a coherent spinel-phase shell on the surface of high-nickel NCM cathodes. This is achieved through a disruptive “subtractive” interface engineering approach, in which flash Joule heating is employed to quantitatively extract surface Li atoms from the NCM. The nanoscale spinel-phase “skin” forms a robust interconnection with the bulk phase, effectively shielding the NCM cathode from reacting with the electrolyte, suppressing the oxygen loss, accommodating isotropic and high-rate Li+ extraction/insertion, and thus facilitating its cycling performance in both half-cells and full pouch cells. Overall, the manuscript is well-organized and presents a meticulous investigation of the plausible mechanism, supported by both computational modeling and experimental evidence. However, I still have a few questions before I can recommend it to be published in Nature Communications.\n(1) DFT calculations indicate that the coating layer can enhance electronic conductivity. However, it remains unclear why such a thin coating layer can distinctly improve Li-ion diffusion in the bulk material. In addition, the mechanism by which this thin layer mitigates the c-axis contraction during the first charge needs to be further clarified.\n(2) The manuscript states that conventional annealing fails to generate the spinel-phase coatings, whereas flash Joule heating enables the formation of metastable structures. Please provide a more detailed analysis of the fundamental differences between these two approaches that result in such contrasting outcomes.\n(3) Please specify the amount of electrolyte used in both the coin cells and the final pouch cells. Although the pouch cells employing graphite as the counter electrode exhibit excellent electrochemical performance, it is recommended that the charge–discharge profiles be included to provide a more complete assessment.\n(4) In the electrochemical performance section, the coated cathodes are reported to exhibit lower surface film resistance and reduced polarization. However, the authors do not provide a clear explanation for the observed improvement in kinetic properties. Typically, coating tends to increase the impedance of cathodes. Please clarify the underlying mechanism responsible for the contrary behavior observed here.\n(5) The authors report enhanced stability of the FJ350 sample based on both interfacial and bulk characterizations, despite the material undergoing only a surface modification. It would be valuable if the authors could discuss and emphasize how improvements in interfacial stability might translate into enhanced bulk structural stability.\n(6) In addition to temperature and duration, what other factors may influence the stripping of lithium and oxygen atoms from the surface? Is this “subtraction” strategy applicable to other cathode systems, such as lithium-rich manganese-based materials and lithium cobalt oxide？\n(7) The EPR results indicate that FJ350 contains a higher concentration of oxygen vacancies. Typically, such defects are expected to deteriorate the initial Coulombic efficiency (ICE). However, in this case, FJ350 exhibits the highest ICE among the samples. Could the authors clarify whether this observation is consistent with current understanding, and provide a plausible explanation for this unexpected behavior?\n(8) Figure 3g presents the evolution of c-axis length and lattice volume during cathode delithiation. Notably, the c-axis initially expands before contracting, whereas the overall lattice volume decreases continuously throughout the process. This apparent inconsistency warrants clarification.\n(9) Some small questions:\n1. The scale bar for Fig. 2i, and Fig. S1 should be provided.\n2. In Fig. S4, the g value from the EPR spectra should be labeled.", "reviewer_2": "This manuscript presents a novel and impactful strategy to suppress oxygen release and enhance structural stability in Ni-rich layered cathodes by inducing a self-derived spinel-phase “skin” through flash Joule heating. Unlike traditional additive coating methods, the authors propose a subtractive strategy that transforms the surface of LiNi0.8Co0.1Mn0.1O2 into a robust, conformal spinel layer. The engineered spinel-phase “skin” effectively confines lattice oxygen, reduces interfacial reactivity, mitigates structural distortion, and substantially improves electrochemical performance, including a remarkable Coulombic efficiency of 95.3% and capacity retention of 80.2% after 2000 cycles in pouch cells.\nThe work is well-conceived, methodically executed, and thoroughly analyzed using a wide range of techniques including in situ XRD, TOF-SIMS, EIS, DSC, and DFT calculations. The concept is novel and highly relevant to the ongoing challenges in Ni-rich cathode development. However, several major points should be clarified or strengthened before the manuscript can be considered for publication.\nComment 1:\nThe authors briefly claim the extension of their strategy to NCM90 and single-crystalline NCM811 (Supplementary Fig. 21), yet no mechanistic validation is provided. Key characterizations—such as interfacial phase confirmation, ICE measurement, or CEI composition analysis—are completely absent. Without these, the claim of general applicability remains speculative. These materials must be investigated with comparable rigor to support the universality of the approach.\nComment 2:\nThe manuscript repeatedly refers to \"surface lithium stripping\" during flash Joule heating but fails to clarify whether this extracted lithium is reintegrated, externally compensated, or permanently lost. This ambiguity undermines the understanding of the system’s lithium inventory and its impact on ICE. Particularly for samples like FJ600, it is unclear whether the reduced ICE arises from irreversible lithium loss or other structural factors. A quantitative analysis of lithium loss and its correlation with electrochemical capacity is essential\nComment 3:\nContrary to the coin cell cycle evaluation results, the discharge capacity of FJ350 in pouch cells is less uniform than that of pristine. Is this due to the inhomogeneity caused by the FJ treatment, which processes powder in 5-10 g quantities? Comments are needed on the inhomogeneity in these data.\nComment 4:\nIn Fig. S27, the authors present the charge transfer resistance of each electrode through EIS analysis. The Nyquist plot used to represent these Rct values is also shown, and it is recommended that all other resistance values be displayed in addition to the Rct values.\nComment 5:\nIn Figure 5a, the authors note that in the highly delithiated state of the pristine sample, the electron density near the Fermi level decreases sharply, suppressing the redox activity of oxygen and nickel, leading to oxygen release and structural collapse. The authors should provide specific references supporting these claims.\nComment 6:\nIt is stated that the in-situ XRD results support the claim that the spinel-phase “skin” reduces the coexistence of the H2 and H3 phases, facilitates reversible structural formation, and promotes uniform lithium distribution. However, the rationale for comparing these results with TOF-SIMS at the 4.3 V cutoff is unclear. It is recommended to modify Fig3. (e, f) to correspond with Fig3. (h, i), enabling a direct comparison of results obtained at the same cutoff voltage\nComment 7:\nTo support the claim of spinel phases of suitable thickness, it would be helpful to add numerical values of spinel layer thickness in addition to cycle data and lattice spacing.", "reviewer_3": "The manuscript introduces a subtractive strategy to engineer a conformal spinel-phase “skin” on Ni-rich layered cathode materials (e.g., NCM811, NCM90) via flash Joule heating (FJH). Unlike conventional additive coatings, this approach selectively strips surface Li and O atoms, forming a coherent spinel layer with tunable thickness. The authors demonstrate that this spinel “skin” provides robust interfacial bonding, improved ionic/electronic conductivity, and strong suppression of lattice oxygen loss. However, the main novel point associated with the DFT calculation should be majorly revised.\n1.“In contrast, samples with either too thick or too thin spinel-phase ‘skin’ coatings (FJ300 and FJ400), although performing better than the pristine sample, are still inferior to FJ350. On one hand, the excessively thin shell lacks sufficient mechanical strength and isolation effect”\nHow do the authors measure the coverage of Li depletion in comparison to the coating method, and then conclude its supremacy? Data needs to be presented.\n2.“The advantage of FJ350 over the pristine sample in reaction kinetics can be elucidated by measuring the Li-ion diffusion coefficient (DLi +)”\nHere the authors are comparing their novel processed cathode with the pristine one. However, the novelty of this point should be a better coating method. Thus, the authors should compare properties between coated NCM and self-induced coated NCM, not just with the pristine one.\n3.“This phenomenon indicates a substantial improvement in the electrical conductivity of material, suggesting a superior rate capability compared to the pristine sample. In contrast, the pristine sample demonstrates a sharp decrease in electronic density of states near the Fermi level at high delithiation states, indicating suppressed redox activity of oxygen and nickel.”\nI do not agree with this statement. A sharp decrease in electronic density of states near the Fermi level cannot alone indicate suppressed redox activity of oxygen and nickel. The oxygen redox should be considered not only from the electrons right at the Fermi level but also from the region close to it. In this sense, I cannot see much difference between Fig. 5 (a) top and bottom. The authors may use LOBSTER to accurately calculate this.\n3.\nI cannot agree with the NEB calculation. Only the ODH migration path is considered. In layered oxide cathode materials, the TSH is regarded as the common diffusion pathway as it has ~1/3 Ea of ODH. This is well studied and reported by Ceder et al. Please refer to:\nhttps://doi.org/10.1103/PhysRevB.79.014305\nhttps://doi.org/10.1021/acsami.3c14444\n“Theoretical methods: All spin-polarized xxxx”\n4. For the DFT calculation, the authors should provide a comprehensive picture and data of their structure. If the structure is as short as given in Fig. 5, the bottom or top vacuum layer would impact the calculation at the interface. The authors are recommended to refer to DFT interface modeling papers:\nhttps://doi.org/10.1021/acsami.2c07560\nhttps://doi.org/10.1021/acsami.4c05737\nMoreover, how did the authors sample such Li-deficient structures? If they simply picked one structure without careful sampling, and if that structure is not the ground state energy, then such DFT calculation data would be totally meaningless."}, "author_response": "Point-by-point Response to Review Comments\nWe express our sincere gratitude to the reviewers for their insightful comments and valuable\nsuggestions, which have greatly contributed to the improvement of this paper. We have\nthoroughly revised the manuscript based on these comments, making it more comprehensive,\nconcise, and coherent. The reviewers’ suggestions have also served as a valuable reminder for\nus to be meticulous in our scientific research. All the revisions have been incorporated into the\nrevised manuscript. In order to address the reviewers’ comments more clearly, we have\ncategorized them into specific question areas and provided detailed responses as follows.\nReviewers’ comments and our Response:\nReviewer #1 (Remarks to the Author): ================================\nIn this work, Yang et al proposes a generalizable and scalable strategy for constructing a\ncoherent spinel-phase shell on the surface of high-nickel NCM cathodes. This is achieved\nthrough a disruptive “subtractive” interface engineering approach, in which flash Joule\nheating is employed to quantitatively extract surface Li atoms from the NCM. The nanoscale\nspinel-phase “skin” forms a robust interconnection with the bulk phase, effectively shielding\nthe NCM cathode from reacting with the electrolyte, suppressing the oxygen loss,\naccommodating isotropic and high-rate Li+ extraction/insertion, and thus facilitating its\ncycling performance in both half-cells and full pouch cells. Overall, the manuscript is well-\norganized and presents a meticulous investigation of the plausible mechanism, supported by\nboth computational modeling and experimental evidence. However, I still have a few questions\nbefore I can recommend it to be published in Nature Communications.\nOur response: Thank you very much for your affirmation of our work and constructive\ncomments on the manuscript. According to the reviewer’s comments, we have carefully revised\nthe manuscript and made the point-to-point responses as below.\nQ1: DFT calculations indicate that the coating layer can enhance electronic conductivity.\nHowever, it remains unclear why such a thin coating layer can distinctly improve Li-ion\ndiffusion in the bulk material. In addition, the mechanism by which this thin layer mitigates the\nc-axis contraction during the first charge needs to be further clarified.\nOur response: Thank you for raising these important points regarding the Li+ diffusion\nenhancement and c-axis contraction mitigation mechanisms enabled by the thin spinel-phase\nskin. The nanoscale spinel-phase “skin” formed via flash Joule heating acts as a coherent\ninterfacial layer that establishes isotropic and high-rate Li+ transport pathways between the\nbulk layered phase and the electrolyte. While DFT calculations confirm that the spinel coating\nitself does not significantly reduce the intrinsic Li+ diffusion barrier compared to the bulk\nLiTMO (Figure 5e), the key advantage lies in its mortise-and-tenon-like interconnection with\n2\nthe bulk phase, which minimizes interfacial resistance and suppresses detrimental phase\ntransitions (e.g., layered-to-rock-salt) that typically block Li+ pathways. Specifically, the high\ncrystallinity of spinel skin and compatibility with the layered structure facilitate continuous Li+\nflux across the interface, as evidenced by galvanostatic intermittent titration technique (GITT)\nmeasurements showing a higher Li+ diffusion coefficient for the modified FJ350 electrode than\nthe pristine sample (Figure 2h).\nMoreover, the thin layer avoids excessive thickness-induced Li+ blockage, i.e., thicker\ncoatings (e.g., FJ500) impede ion transport due to increased diffusion barriers, while\ninsufficiently thin layers (e.g., FJ300) fail to provide full coverage. Therefore, the optimized\nthickness strikes a balance between enabling rapid interfacial Li+ exchange and maintaining\nbulk diffusion kinetics, thereby enhancing the rate capability and cycling stability (Figures 2c\nand 2f). More importantly, the spinel-phase skin can suppress the c-axis contraction during the\nH2-H3 phase transition by mitigating lattice oxygen loss and reinforcing structural integrity. In\nsitu XRD analyses reveal that the pristine electrode undergoes pronounced c-axis lattice\ncontraction at high delithiation states due to oxygen evolution and anisotropic strain, whereas\nthe FJ350 electrode exhibits significantly suppressed lattice variation and fully reversible\nstructural recovery upon discharge (Figures 3e-3g). In addition, we demonstrate that the high\nmechanical strength and rapid reaction kinetics of the spinel-phase “skin”-encapsulated FJ350\nplay a key role in promoting the complete transformation of H2-H3. Otherwise, the coexistence\nof partially converted H2 and H3 phases tends to induce lattice mismatch within the particles,\nleading to inhomogeneous Li distribution and harmful structural distortion (Figures 3h-3k).\nGiven the above, the completeness of the H2-H3 phase transition, along with the suppression\nof the loss of oxidized lattice oxygen, greatly enhances the reversibility of the H2-H3\ntransformation, which is crucial for mitigating c-axis contraction and maintaining structural\nintegrity. As depicted in Figure 2g, the H2-H3 phase transition around 4.2 V is significantly\ndiminished in the pristine sample after 500 cycles, indicating severe structural changes and\nlattice oxygen loss. In stark contrast, the FJ350 sample maintains a stable H2-H3 redox peak,\nsuggesting reinforced structural integrity and effective stabilization of lattice oxygen, thereby\ngreatly suppressing its release. This also directly results in the more intact particle structure\nand minimal cracking of FJ350 after long-term cycling (Figure 4b).\nAs evidenced by DFT, the skin’s lower Fermi level and higher work function also reduce\nelectron loss from the bulk, minimizing oxidative damage at the surface. Consequently, the\nskin promotes a complete H2-H3 transition without phase coexistence, alleviating internal\nstrain and crack formation. In summary, the thin spinel-phase layer enhances Li+ diffusion by\nproviding a kinetically favorable interface and mitigates c-axis contraction through structural\nstabilization, both of which are critically dependent on precise thickness control.\nAccording to the reviewer's constructive suggestions, we have revised the relevant parts on\npage 14 of the revised manuscript. We copy the relevant text here for your check:\n“Given the above, the coating layer stabilizes the interface by preventing the formation of\ninactive phases, such as the rock-salt phase. This interfacial stability ensures efficient Li+\ntransport across the interface and suppresses lattice oxygen loss, thereby enhancing the overall\nLi⁺ diffusion within the particles. These findings highlight the critical role of the coating in\noptimizing both interfacial and bulk properties, leading to significantly improved\nelectrochemical performance.”\nQ2: The manuscript states that conventional annealing fails to generate the spinel-phase\ncoatings, whereas flash Joule heating enables the formation of metastable structures. Please\nprovide a more detailed analysis of the fundamental differences between these two approaches\nthat result in such contrasting outcomes.\nOur response: We are grateful for the very professinal and insightful comment. Flash Joule\nheating, due to its rapid heating and cooling rates, primarily affects the surface and subsurface\nof the cathode material when the temperature is precisely controlled. When the processing time\nis limited to a few to several tens of seconds, trace lithium extraction can be induced, thereby\npromoting a surface phase transition from the layered structure to a spinel phase in the\nNCM811 cathode material. Notably, this corresponds to a surface phase transition occurring\nonly within a few nanometers in thickness for the cathode. In stark contrast, the tube furnace\nannealing method, with slow heating rate (around 5 °C min-1) and long dwell time (several\nhours), leads to severe lithium and oxygen loss, which often involves degradation of the bulk\nphase. Therefore, the tube-furnace treatment fails to produce the desired spinel-phase coating\nand instead results in the formation of a thicker rock-salt shell. The fundamental difference\nbetween the two methods lies in the heating rate, only under transient thermal pulses within an\nextremely short duration can a nanometer-scale spinel shell be formed and preserved.\nAccording to the reviewer’s constructive suggestions, we have revised the relevant part on\npage 8 of the supplementary information. We copy the relevant text here for your check:\n“Undoubtedly, the unique ultrafast heating and rapid cooling mechanisms of the Joule heating\ntechnique provide a crucial technical prerequisite for achieving nano-scale precise phase\ntransformation at the cathode material interface.”\nQ3: Please specify the amount of electrolyte used in both the coin cells and the final pouch\ncells. Although the pouch cells employing graphite as the counter electrode exhibit excellent\nelectrochemical performance, it is recommended that the charge-discharge profiles be included\nto provide a more complete assessment.\nOur response: Thank you for your insightful comments. In response to the inquiry regarding\nthe electrolyte amount, the electrolyte used in the coin cells was 60 µL. For the pouch cells,\nwe employed an electrolyte injection factor of 5, which refers to the ratio of the electrolyte weight\nto the cell capacity (g/Ah). This choice was made to strike a balance between ensuring high\nenergy density and maintaining sufficient ionic conductivity, thereby enabling long-cycle\nstability.\nAs highlighted by the reviewer, the inclusion of charge-discharge profiles is crucial for a\nmore comprehensive evaluation of the electrochemical performance. In accordance with this\nsuggestion, we have added the charge-discharge curves in new Supplementary Fig. 21. These\ncurves clearly illustrate the advantages of the FJ350 electrode, particularly its remarkable\nability to mitigate capacity fading and resist electrode polarization during cycling. This\ndemonstrates the material’s potential for long-term cycling stability and efficiency, positioning\nit as a promising candidate for practical applications.\nNew Supplementary Figure 21\nSupplementary Fig. 21 | Charge-discharge profiles of (a) the pristine NCM811-based pouch\ncell and (b) the FJ350-based pouch cell within the voltage range of 2.7-4.2 V at 1 C rate for\nover 2000 cycles.\nAs the reviewer suggested, relevant information have been provided on page 12 the revised\nmanuscript and pages 3 and 24 of the supplementary information. We copy the relevant text\nhere for your check:\n“The electrolyte was composed of 1 M LiPF dissolved in ethylene carbonate (EC) and diethyl\n6\ncarbonate (DEC) with a volume ratio of 3:7 and 2 wt.% vinylene carbonate addition. The\nelectrolyte amount used in the coin cells was 60 µL.”\n“The negative-to-positive capacity ratio (N/P ratio) for the pouch cell is around 1.12 and the\nelectrolyte injection factor was 5, which refers to the ratio of the electrolyte weight to the cell\ncapacity (g/Ah).”\n“The charge-discharge curves further demonstrate the advantage of the FJ350 electrode in\nmitigating capacity fade and resisting electrode polarization, highlighting its outstanding long-\nterm cycling stability (Supplementary Fig. 21).”\nQ4: In the electrochemical performance section, the coated cathodes are reported to exhibit\nlower surface film resistance and reduced polarization. However, the authors do not provide a\nclear explanation for the observed improvement in kinetic properties. Typically, coating tends\nto increase the impedance of cathodes. Please clarify the underlying mechanism responsible\nfor the contrary behavior observed here.\nOur response: Thank you for your constructive recommendations. We apologize for the lack\nof a clear explanation. The improved kinetic properties after surface engineering are primarily\nattributed to the spinel-phase coating layer. Density functional theory (DFT) calculations\nindicate that this coating layer can significantly improves the electronic conductivity of the\nelectrode, creating efficient charge transfer pathways that enhance overall performance. This\ntheoretical insight accords well to the experimental observations of improved electrochemical\nproperties. Based on these enhanced electronic properties, galvanostatic intermittent titration\ntechnique (GITT) measurements further demonstrate that the FJ350 electrode exhibits a faster\nLi⁺ diffusion rate compared to the pristine electrode (Figure 2h), confirming that the coating\nimproves the ion transport efficiency within the electrode. Additionally, the surface coating\nstabilizes the cathode-electrolyte interface (CEI) by suppressing side reactions, resulting in a\nthinner and more stable CEI film (Figures 4f and Supplementary Fig. 38-41). Moreover,\nelectrochemical impedance spectroscopy (EIS) measurements reveal that both the R (CEI film\nf\nresistance) and R (charge-transfer resistance) of J350 remain significantly lower than those of\nct\nthe pristine cathode, indicating that the coating layer effectively reduces the interfacial\nimpedance and improved the reaction kinetics (as manifested in the new Supplementary Fig.\n34). The decrease in interfacial resistance further confirms the reduction in electrochemical\npolarization and improvement in overall performance the FJ350 electrode. Given the above,\nthe spinel-phase skin improves electronic conductivity, accelerates Li⁺ diffusion, stabilizes the\nCEI, and reduces electrochemical impedance, thereby significantly enhancing electrochemical\nperformance and minimizing polarization.\nAccording to the reviewer’s constructive suggestions, we have revised the relevant part on\npage 17 of the revised manuscript and page 37 of the supplementary information. We copy the\nrelevant text here for your check:\n“As manifested in Supplementary Fig. 34, the FJ350 electrode exhibits markedly lower CEI\nfilm resistance (R) and charge transfer resistance (R ) at all stages of cycling[42]. The\nf ct\npronounced reduction in R confirm that the spinel-phase skin establishes a percolating\nct\nelectron pathway, synergizing with fast Li+ transport to achieve high rate capability and cycling\ncapacity retention.”\nNew Supplementary Figure 34\nSupplementary Fig. 34 | EIS spectra at selected cycle numbers for the (a) pristine and (b)\nFJ350 electrodes. (c) Comparison of the fitted R values of the pristine and FJ350 electrodes\nf\nduring cycling. (d) Comparison of the fitted R values of the pristine and FJ350 electrodes\nct\nduring cycling.\nQ5: The authors report enhanced stability of the FJ350 sample based on both interfacial and\nbulk characterizations, despite the material undergoing only a surface modification. It would\nbe valuable if the authors could discuss and emphasize how improvements in interfacial\nstability might translate into enhanced bulk structural stability.\nOur response: Thanks for your very professional suggestion. The performance deterioration\nof Ni-rich cathodes is generally recognized as the result of a cascade of interlinked degradation\npathways. Among these, collapse along the c-axis at elevated voltages has been regarded as a\ncritical source of bulk structural instability. As extensively documented (e.g., Peter et al., Nat.\nChem., 2016, 8, 684; Meng et al., Nat. Rev. Mater., 2022, 7, 522), the contraction of the lattice\nc parameter under deep delithiation arises from the rehybridization between high-valence\ntransition metal ions, particularly Ni4+ in Ni-rich compositions, and oxygen. In this state, the\nstrongly oxidizing metal ions tend to withdraw electrons from coordinated oxygen, thereby\nraising the oxygen p orbitals closer to the Fermi level (E). This enhances the overlap between\nf\nTM 3d and O 2p orbitals, ultimately destabilizing the oxygen. The subsequent release of\noxygen reduces the interlayer Coulomb repulsion within the Li-O slabs, driving contraction\nalong the c-axis. More importantly, oxygen loss is thought to originate at the particle surface\nand propagate inward (e.g., Sun et al., Adv. Energy Mater. 2022, 2200569; Wang et al. Nat.\nCommun. 2018, 9, 2437). Consequently, strategies that mitigate oxygen release at the surface\nare considered highly effective for reinforcing the bulk stability of Ni-rich layered oxides.\nIn this work, the enhanced bulk structural stability of FJ350 can be attributed to multiple\nfactors. First, DFT simulations (Figure 5h) reveal that, in the highly deintercalated state,\nstronger TM-O bonds are formed at the interface, accompanied by a lower Fermi level and a\nhigher work function, which collectively contribute to improved structural robustness during\nthe redox process. Second, electrochemical evidence further supports this conclusion. After\n500 cycles, the H2-H3 redox peak around 4.2 V in Figure 2g, which is associated with c-axis\ncontraction and lattice oxygen oxidation, becomes significantly diminished in the pristine\nsample, indicating severe structural degradation and oxygen loss. In sharp contrast, the FJ350\nsample retains a stable H2-H3 redox peak, demonstrating reinforced structural integrity and\neffective stabilization of lattice oxygen. Notably, the c-axis variation is reduced from 4.31% in\nthe pristine sample to 2.12% in FJ350 (Figure 3g), thereby mitigating volume expansion and\nfurther enhancing the overall structural stability. Finally, the surface spinel layer not only\nsuppresses interfacial failure caused by side reactions through a physical barrier effect, but also\nalleviates c-axis contraction via electron transport between the modified layer and the bulk Ni-\nrich material, thus effectively enhancing bulk stability.\nAccording to the reviewer’s constructive suggestions, we have revised the relevant part on\npage 17 of the revised manuscript. We copy the relevant text here for your check:\n“The enhanced bulk stability of FJ350 can be attributed to interfacial electron transfer between\nthe spinel-phase shell and the bulk Ni-rich lattice, which mitigates c-axis variation and confines\nlattice oxygen under highly delithiated conditions. Consequently, this process alleviates lattice\ncollapse induced by oxygen release from the surface to the bulk.”\nQ6: In addition to temperature and duration, what other factors may influence the stripping of\nlithium and oxygen atoms from the surface? Is this “subtraction” strategy applicable to other\ncathode systems, such as lithium-rich manganese-based materials and lithium cobalt oxide？\nOur response: We sincerely appreciate the professional and insightful comment. Indeed, this\nis a critical question that we carefully considered during the preparation of the initial\nmanuscript. Beyond temperature and reaction duration, several additional parameters can\nsubstantially influence the stripping behavior of lithium and oxygen atoms from the surface. In\nparticular, the ambient atmosphere plays a decisive role, whether the process occurs in air, an\ninert gas such as argon, or a reducing environment like hydrogen can significantly alter the\nreaction pathways. Similarly, pressure conditions (positive pressure, negative pressure, or\nambient pressure) can modulate diffusion kinetics and surface thermodynamics, thereby\naffecting both the rate and the extent of atomic stripping. Although our current experiments\nhave not yet systematically investigated these factors, they are the focus of our ongoing\ntheoretical analysis and planned experimental studies. We believe that incorporating these\nvariables will yield a more comprehensive understanding of the mechanisms governing surface\natom stripping. We sincerely appreciate the reviewer’s attention to this important aspect and\nhope they will continue to follow our future progress in this area.\nFurthermore, in addition to the ternary cathode materials (e.g., NCM811), we have also\nconducted Joule heating treatment on lithium-rich manganese materials (Li-rich) and lithium\ncobalt oxide (LCO) materials and sodium nickel manganese oxide (NFM). This represents a\nkey direction of our ongoing and future research efforts. As shown in the Figures R1-R3 (Note\nthat Figure RX represents the figures used in the response letter and will not appear in the\nmanuscript and supporting information), the Joule heating process significantly enhances the\nelectrochemical performance of all three systems, underscoring the broad applicability of our\n“subtraction” strategy.\nFigure R1 | Initial charge-discharge profiles of the lithium-rich manganese electrodes before\nand after flash-sintering treatment at 0.05 C.\nFor Li-rich cathode materials, the samples subjected to Joule heating exhibit remarkable\nimprovements in electrochemical performance. As presented in Figure R1, the initial\ncoulombic efficiency significantly increases from 72.4% for the pristine sample to 86.5% after\nJoule heating, indicating a substantial reduction in irreversible capacity loss during the first\ncharge-discharge cycle. The treated electrode also delivers a higher discharge capacity and a\nmore stable voltage plateau.\nFigure R2 | (a) Initial charge-discharge profiles at 0.1 C of the lithium cobalt oxide (LCO)\nelectrodes before and after flash-sintering treatment. (b) Cycling performance of the pristine\nand treated LCO electrodes for 200 cycles at 0.5 C within 3-4.65 V.\nIn the case of LCO cathode, the treated samples exhibit superior electrochemical properties,\nwith the initial Coulombic efficiency increasing from 89.4% to 94.1% (Figure R2a). More\nimportantly, the capacity retention after 200 cycles improves dramatically from 65.3% for the\npristine sample to 95.1% within 3-4.65 V (Figure R2b). These improvements convincingly\ndemonstrate that Joule heating promotes the formation of a more robust surface and a stabilized\nelectrode-electrolyte interface under high-voltage conditions.\nFigure R3 | (a) Initial charge-discharge profiles at 0.1 C of the sodium nickel manganese oxide\n(NFM) electrodes before and after flash-sintering treatment. (b) Cycling performance of the\npristine and treated NFM electrodes for 140 cycles at 0.5 C within 2-4.2 V.\nSimilarly, for the NFM cathode, the samples subjected to Joule heating deliver an increased\ninitial Coulombic efficiency (from 87.3% to 91.4%) and an improved capacity retention (from\n20.1% to 55.1% after 100 cycles, Figure R3). These results clearly demonstrate that the rapid\nthermal activation can also reinforce the structural integrity and interfacial stability of sodium-\nion cathodes.\nGiven the above, these results confirm that the Joule heating method possesses broad\napplicability, versatility, and scalability across different cathode chemistries. The rapid and\ncontrollable thermal treatment not only enhances the initial coulombic efficiency and long-term\ncycling stability but also provides a general strategy for regulating interfacial stability and\nredox reversibility in layered oxide systems. Of course, the microscopic mechanisms\nunderlying the performance enhancement of the above cathode systems need to be further\nelucidated in our future work.\nQ7: The EPR results indicate that FJ350 contains a higher concentration of oxygen vacancies.\nTypically, such defects are expected to deteriorate the initial Coulombic efficiency (ICE).\nHowever, in this case, FJ350 exhibits the highest ICE among the samples. Could the authors\nclarify whether this observation is consistent with current understanding, and provide a\nplausible explanation for this unexpected behavior?\nOur response: Thank you for this insightful comment. In oxide-based cathode materials, a\nhigh concentration of oxygen vacancies can indeed lead to a reduction in the initial coulombic\nefficiency (ICE), which is typically attributed to enhanced electrolyte decomposition at defect\nsites. In contrast, the oxygen vacancies in the FJ350 sample of this study are associated with\nthe reconstructed surface layer, which fundamentally alters their electrochemical functionality.\nThe presence of surface and subsurface oxygen vacancies in FJ350 lowers the average\noxidation state of transition metals, thereby mitigating the extent of lattice oxygen oxidation\nduring high-voltage charging. This suppresses oxygen release and reduces the hysteresis\nassociated with anionic redox activity, consistent with the in situ DEMS results (new\nSupplementary Fig. 13). Consequently, the irreversible capacity loss related to oxygen\nevolution reactions is significantly reduced, leading to an improved initial coulombic efficiency\n(ICE). More importantly, the reconstructed surface layer in FJ350 effectively passivates the\nelectrode-electrolyte interface, thereby inhibiting deleterious side reactions at high voltages.\nAs illustrated in Figures 3a, 3b, the pristine material undergoes significant surface degradation,\nforming a rock-salt phase outer layer during the initial cycle, which blocks Li⁺ transport and\nleads to irreversible capacity loss. In contrast, FJ350 maintains a spinel-phase surface structure\nunder the same conditions, thereby preserving facile Li⁺ (de)intercalation pathways and\nreducing side reactions.\nTaken together, the synergistic effect of suppressing bulk oxygen loss and stabilizing the\nsurface structure significantly improves the initial coulombic efficiency (ICE), increasing it\nfrom 85% for the pristine material to 95.3% for FJ350. These findings indicate that the\ninfluence of oxygen vacancies largely depends on their local coordination environment and the\naccompanying interfacial chemistry. In FJ350, the coexistence of surface/subsurface oxygen\nvacancies and the reconstructed surface layer jointly contribute to enhanced structural and\ninterfacial stability, thereby accounting for the observed improvement in the initial coulombic\nefficiency.\nNew Supplementary Figure 13\nSupplementary Figure 13 | In situ DEMS data of (a) Pristine, (b) FJ300, (c) FJ350, (d) FJ400\nduring the first charge within 2.7-4.7 V.\nAccording to the reviewer’s constructive suggestions, relevant information have been\nprovided on page 9 of the revised manuscript and page 16 the supplementary information. We\ncopy the relevant text here for your check:\n“In situ differential electrochemical mass spectrometry (DEMS) was employed to monitor the\noxygen evolution of different samples under high voltages (e.g., 4.7 V) in order to evaluate the\ncoverage of the spinel-phase shell. As Supplementary Fig. 13 demonstrates, the FJ300 sample\nwith a thin and incomplete spinel-phase shell already exhibits a certain degree of suppression\nof oxygen evolution during the first charge compared with the pristine sample, but the reduction\nis limited, decreasing only from 4.6 to 3.1 nmol min-1. In sharp contrast, the oxygen-release\nsignals of the FJ350 and FJ400 electrodes are significantly lower, with FJ350 in particular\nreduced to as low as 1.2 nmol min-1. This result indicates that the spinel-phase shell can\neffectively inhibit oxygen evolution even under the extremely high cutoff voltage of 4.7 V,\ndirectly reflecting both the completeness and the strong spatial confinement effect of the FJ350\ncoating layer. The suppression of oxygen release is also consistent with the increased\nCoulombic efficiency and reduced irreversible lithium consumption observed in the\nelectrochemical tests, further demonstrating the critical role of an optimized coating thickness\nin achieving stable cycling performance.”\nQ8: Figure 3g presents the evolution of c-axis length and lattice volume during cathode\ndelithiation. Notably, the c-axis initially expands before contracting, whereas the overall lattice\nvolume decreases continuously throughout the process. This apparent inconsistency warrants\nclarification.\nOur response: We sincerely appreciate the insightful comment. The observed evolution of the\nlattice parameters during delithiation is indeed characteristic of layered oxide cathode materials.\nSpecifically, in ternary cathodes, the initial expansion of the c-axis upon charging arises\nprimarily from the electrostatic repulsion between adjacent oxygen layers. As Li⁺ ions are\nextracted from the interslab region, the reduced screening effect increases the O-O repulsion,\nleading to a temporary elongation of the c-axis. At higher states of delithiation, however, the c-\naxis contraction occurs due to the H2-H3 phase transition, which is associated with a reduction\nof the interlayer spacing.\nMeanwhile, the overall lattice volume decreases monotonically throughout the delithiation\nprocess. This behavior can be rationalized by the dominant contraction, which outweighs the\ntransient expansion along the c-axis. Upon oxidation of the transition-metal ions, the TM-O\nbonds shorten significantly, inducing a strong in-plane contraction. The cumulative effect of\nthis bond shortening exceeds the temporary c-axis elongation caused by electrostatic repulsion,\nultimately resulting in a net reduction in the total lattice volume.\nAccording to the reviewer’s constructive suggestions, we have revised the relevant part on\npage 16 of the revised manuscript. We copy the relevant text here for your check:\n“The initial expansion of the c-axis is due to increased electrostatic repulsion between oxygen\nlayers as Li⁺ ions are extracted. At higher states of charge, the contraction along the c-axis\narises from electron withdrawal by highly oxidized transition-metal ions from the coordinated\noxygen atoms, leading to oxygen ion oxidation and a consequent reduction in the repulsive\ninteractions within the delithiated Li layers. Meanwhile, the lattice volume decreases\ncontinuously due to dominant in-plane contraction from the shortening of TM-O bonds as the\ntransition metals are oxidized.”\nQ9: Some small questions:\n1. The scale bar for Fig. 2i, and Fig. S1 should be provided.\n2. In Fig. S4, the g value from the EPR spectra should be labeled.\nOur response: We sincerely apologize for not noticing these details earlier. We have made the\nnecessary revisions and added the missing information. The updated version now includes scale\nbars for Figure 2i and Supplementary Fig. 1. In addition, the g-values have been indicated in\nthe EPR spectra of Supplementary Fig. 4. We hope these updates meet your expectations.\nAccording to the reviewer’s constructive suggestions, we have revised the relevant parts on\npage 10 of the revised manuscript and pages 4 and 7 of the supplementary information. We\ncopy the relevant text here for your check:\nRevised Figure 2i\nFigure 2i | Digital photograph of a smartphone being charged by a pouch full cell configured\nwith a FJ350 cathode and graphite anode. Scale bar: 20 mm.\nRevised Supplementary Figure 1\nSupplementary Fig. 1 | (a) Digital photograph of the self-built Joule heating setup equipped\nwith a temperature control system and a circulating cooling system. (b) A real photograph of a\n10 cm*10 cm flash Joule heating clamp in working condition. The inset image shows the\noperating condition during a single loading of a 5.0 g NCM811 sample.\nNew Supplementary Figure 4\nSupplementary Fig. 4 | Electron paramagnetic resonance (EPR) spectra of the pristine and\nFJ350 samples.\nReviewer #2 (Remarks to the Author): ================================\nThis manuscript presents a novel and impactful strategy to suppress oxygen release and\nenhance structural stability in Ni-rich layered cathodes by inducing a self-derived spinel-phase\n“skin” through flash Joule heating. Unlike traditional additive coating methods, the authors\npropose a subtractive strategy that transforms the surface of LiNi Co Mn O into a robust,\n0.8 0.1 0.1 2\nconformal spinel layer. The engineered spinel-phase “skin” effectively confines lattice oxygen,\nreduces interfacial reactivity, mitigates structural distortion, and substantially improves\nelectrochemical performance, including a remarkable Coulombic efficiency of 95.3% and\ncapacity retention of 80.2% after 2000 cycles in pouch cells. The work is well-conceived,\nmethodically executed, and thoroughly analyzed using a wide range of techniques including\nin situ XRD, TOF-SIMS, EIS, DSC, and DFT calculations. The concept is novel and highly\nrelevant to the ongoing challenges in Ni-rich cathode development. However, several major\npoints should be clarified or strengthened before the manuscript can be considered for\npublication.\nOur response: We are grateful to the reviewer for the positive and professional comments. All\nrequested revisions based on the constructive comments from the reviewer have been fully\nconsidered and thoroughly addressed, as shown below.\nQ1: The authors briefly claim the extension of their strategy to NCM90 and single-crystalline\nNCM811 (Supplementary Fig. 21), yet no mechanistic validation is provided. Key\ncharacterizations-such as interfacial phase confirmation, ICE measurement, or CEI\ncomposition analysis-are completely absent. Without these, the claim of general applicability\nremains speculative. These materials must be investigated with comparable rigor to support\nthe universality of the approach.\nOur response: We sincerely appreciate the reviewer’s valuable and insightful comments.\nYour feedback is highly appreciated and has helped us identify key areas where our study could be\nfurther strengthened. As the reviewer rightly pointed out, the current manuscript lacks\nsufficient characterization data to substantiate the applicability of our strategy to NCM90 and\nsingle-crystal NCM811 cathodes, particularly regarding interfacial phase analysis, initial\ncoulombic efficiency (ICE), and the composition of the cathode-electrolyte interphase (CEI).\nNew Supplementary Figure 26\nSupplementary Fig. 26 | HRTEM and FFT images of (a) the pristine LiNi Co Mn O\n0.9 0.05 0.05 2\n(NCM90) and (b) Treated NCM90 sample. Scale bars: 2 nm.\nNew Supplementary Figure 27\nSupplementary Fig. 27 | HRTEM and FFT images of (a) the pristine single-crystal NCM811\n(SC-NCM811) and (b) Treated SC-NCM811 sample. Scale bars: 2 nm.\nIn response to these suggestions, we have conducted additional experiments and provided\nmore data and analyses to comprehensively verify the general applicability of our approach to\nthese materials. First, we performed high-resolution transmission electron microscopy\n(HRTEM) analyses to investigate the interfacial phases of NCM90 and single-crystal NCM811\nbefore and after Joule heating treatment. As disclosed in Supplementary Figs. 26 and 27, a\npronounced phase transformation was observed on the surfaces of both NCM90 and NCM811,\nwhere the layered structure converted into a spinel-phase shell after flash Joule heating. These\nfindings indicate that our flash sintering strategy can effectively modify the surface structure\nof both materials, which is expected to improve cycling stability and suppress interfacial side\nreactions. Subsequently, to further validate the effectiveness of this strategy, we measured the\nICEs of the NCM90 and single-crystal NCM811 cathodes.\nSupplementary Fig. 28 | (a) Initial charge-discharge profiles and (b) long-term cycling\nperformance of the LiNi Co Mn O before and after flash sintering at 350 °C for 30 s.\n0.825 0.115 0.06 2\n(c) Initial charge-discharge profiles and (d) long-term cycling performance of\nLiNi Co Mn O (NCM90) before and after flash sintering at 350 °C for 30 s.\n0.9 0.05 0.05 2\n(e) Initial\ncharge-discharge profiles and (f) long-term cycling performance of the single-crystal\nLiNi Co Mn O (SC-NCM811) before and after flash sintering at 350 °C for 30 s.\n0.8 0.1 0.1 2\nAs depicted in Supplementary Fig. 28, the ICE of NCM90 increased from 91.0% to\n94.1%, while that of single-crystal NCM811 improved from 85.2% to 88.3%. This\nimprovement can be attributed to the significant suppression of interfacial side reactions and\nlattice oxygen loss, especially after treatment, which enhanced the stability of the interface.\nThese results further suggest that the treated materials effectively minimize irreversible\ncapacity loss during the first cycle, thereby improving the overall battery performance. In\naddition, TOF-SIMS analysis (Supplementary Fig. 40 and 41) of the materials after cycling\nfurther revealed that the flash treated NCM90 and NCM811 cathodes exhibited reduced metal\ndissolution and enhanced chemical stability over extended cycling.\nNew Supplementary Figure 40\nSupplementary Fig. 40 | TOF-SIMS secondary depth profiling analysis of selected species in\nLiNi Co Mn O electrode after cycling, before and after flash sintering treatment. The\n0.9 0.05 0.05 2\ninset shows the TOF-SIMS 3D reconstructions of the C H OF-, NiF -, CoF - , MnF - species.\n2 3 3 3 3\nNew Supplementary Figure 41\nSupplementary Fig. 41 | TOF-SIMS secondary depth profiling analysis of selected species in\nsingle-crystal LiNi Co Mn O electrode after cycling, before and after flash sintering\n0.8 0.1 0.1 2\ntreatment. The inset shows the TOF-SIMS 3D reconstructions of the C H OF-, NiF -, CoF - ,\n2 3 3 3\nMnF - species.\npages 13 and 20 of the revised manuscript and pages 29, 30, 43, 44 of the supplementary\ninformation. We copy the relevant text here for your check:\n“More importantly, this precise interfacial regulation strategy, namely, the transformation from\na layered structure to a spinel shell at the interface (Supplementary Fig. 26 and Fig. 27),\nenhances the reversible capacity and cycling stability of various cathode materials, including\npolycrystalline LiNi Co Mn O , polycrystalline LiNi Co Mn O (NCM90), and\n0.825 0.115 0.06 2 0.9 0.05 0.05 2\nsingle-crystal LiNi Co Mn O (SC-NCM811). These results highlight the broad\n0.8 0.1 0.1 2\napplicability of this strategy in tailoring surface atomic composition and stabilizing interfacial\nphases (Supplementary Fig. 28).”\n“This interfacial stabilization strategy demonstrates universal effectiveness across different\nhigh-Ni cathode materials, as evidenced by the TOF-SIMS analyses of LiNi Co Mn O\n0.9 0.05 0.05 2\n(Supplementary Fig. 40) and single-crystal LiNi Co Mn O electrodes (Supplementary\n0.8 0.1 0.1 2\nFig. 41). Both systems exhibit a thinner CEI layer and markedly reduced fragment intensities,\nconfirming the cross-system applicability of this approach.”\nQ2: The manuscript repeatedly refers to \"surface lithium stripping\" during flash Joule heating\nbut fails to clarify whether this extracted lithium is reintegrated, externally compensated, or\npermanently lost. This ambiguity undermines the understanding of the system’s lithium\ninventory and its impact on ICE. Particularly for samples like FJ600, it is unclear whether the\nreduced ICE arises from irreversible lithium loss or other structural factors. A quantitative\nanalysis of lithium loss and its correlation with electrochemical capacity is essential.\nOur response: We thank the reviewer for raising this important question regarding the fate of\nlithium during the flash Joule heating process and its potential impact on the initial coulombic\nefficiency (ICE). First, we can state with high confidence that the extracted lithium is\npermanently lost during the flash treatment process, as lithium loss is a common\nphenomenon during the thermal treatment of cathode materials (e.g., Lee et al., Adv. Energy\nMater. 2024, 14, 2400662; Choe et al., Energy Storage Mater. 2025, 74, 103884). However, it\nis important to note that under conventional concepts and experimental conditions, such lithium\nloss cannot be precisely controlled, which is exactly what has been achieved in this work.\nAlthough the lithium loss during the flash Joule heating process is irreversible, its\nimpact on the ICE is negligible, primarily because the loss occurs only within a few\nnanometers at the surface and subsurface regions. In contrast, as demonstrated in this work,\nthe formation of a spinel-phase shell induced by controlled lithium extraction effectively\nsuppresses lattice oxygen loss and interfacial side reactions, thereby significantly enhancing\nthe ICE. In the following section, we provide a simple mathematical estimation to illustrate the\nnegligible influence of trace interfacial lithium loss on the overall ICE.\nAs presented in the new Supplementary Fig. 45, we established a simplified mathematical\nmodel to quantify the lithium loss during flash Joule treatment. Based on our experimental\nresults, the key parameters used in the calculations are set as follows:\nSupplementary Fig. 45 | Schematic illustration of a typical spherical NCM particle model\n(radius set to 5 μm), in which a representative 10 nm-thick surface shell region affected by the\nflash Joule heating treatment is delineated based on our experimental observations. We defined\nthree different compositions for the surface shell including the original layered structure, spinel\nphase and rock-salt phase, and for simplicity, the coexistence of multiple phases was not\nconsidered in the calculation.\nThe above model is used to semi-quantitatively estimate lithium loss during flash Joule\nprocessing.\nStep 1: The key parameters used in the calculations are as follows:\n Radius of particle: r=5 μm=5000 nm\n Thickness of Li-deficient shell: t=10 nm\n Volume of the particle:\n4 4\n𝑉 = 𝜋𝑅3 = ×3.14×(5000)3 = 5.236×1011 nm3\n𝑡𝑜𝑡𝑎𝑙 3 3\n Volume of the Li-deficient shell:\n4 4\n𝑉 = 𝜋𝑅3 − 𝜋(𝑅 −𝑡)3 = 3.136×109 nm3\n𝑠ℎ𝑒𝑙𝑙 3 3\n The proportion of lithium in the 10 nm surface layer:\n𝑉 3.136×109\n𝑠ℎ𝑒𝑙𝑙\n𝑓 = = ≈ 0.6%\n𝑉 5.236×1011\n𝑡𝑜𝑡𝑎𝑙\n The molar masses:\nM (LiNi Co Mn O ) = 97.28 g/mol\nLayered 0.8 0.1 0.1 2\nM (Li Ni Co Mn O , x=0.5) =93.81 g/mol\nSpinel x 0.8 0.1 0.1 2\nM (Li TMO, the phase occurred in XRD) =75.80 g/mol\nRock-salt 0.21\n The typical densities (obtained by the MDI Jade database):\nρ (LiNi Co Mn O ) = 4.74 g/cm³\nLayered 0.8 0.1 0.1 2\nρ (Li Ni Co Mn O , x=0.5) = 4.79 g/cm³\nSpinel x 0.8 0.1 0.1 2\nρ (Li TMO, the phase occurred in FJ700 and FJ800) = 5.85 g/cm³\nRock salt 0.21\n The number of lithium atoms per unit volume of different samples:\nρ 4.74\n𝑛 = NCM811 𝑁 = ×6.022×1023 ×10−21 = 29.35 atoms/nm3\nLi,Layered M 𝐴 97.28\nNCM811\nρ 4.79\n𝑛 = spinel 𝑁 = ×6.022×1023 ×10−21 ×0.5\nLi,Spinel M 𝐴 93.81\nspinel1\n= 15.37 atoms/nm3\nρ 5.85\n𝑛 = rock salt 𝑁 = ×6.022×1023 ×10−21 ×0.21\nLi,Rock−salt M 𝐴 75.80\nrock salt\n= 9.76 atoms/nm3\nStep 2: When the surface undergoes a layered-to-spinel phase transformation\n Li atoms in the layered surface:\n𝑁 = 𝑛 ×𝑉 = 29.35×3.136×109 = 9.20×1010\nLi,shell,Layered Li,NCM811 𝑠ℎ𝑒𝑙𝑙\n Li atoms in the spinel surface:\n𝑁 = 𝑛 ×𝑉 = 15.37×3.136×109 = 4.82×1010\nLi,shell,spinel Li,spinel 𝑠ℎ𝑒𝑙𝑙\n Relative lithium loss:\n9.20−4.82\n×0.6% ≈ 0.286%\n9.20\n The reduced capacity caused by lithium loss in spinel phase:\n0.00286×275.7 (Theoretical capacity of NCM811) ≈ 0.79 mAh g−1\nStep 3: When the surface undergoes a layered-to-rock-salt phase transformation\nLi atoms in the layered surface:\n𝑁 = 𝑛 ×𝑉 = 29.35×3.136×109 = 9.20×1010\nLi,shell,NCM811 Li,NCM811 𝑠ℎ𝑒𝑙𝑙\nLi atoms in the rock-salt surface:\n𝑁 = 𝑛 ×𝑉 = 9.76×3.136×109 = 3.06×1010\nLi,shell,rock−salt Li,rock−salt 𝑠ℎ𝑒𝑙𝑙\nRelative lithium loss:\n9.20−3.06\n×0.6% ≈ 0.401%\n9.20\nThe reduced capacity caused by lithium loss in rock-salt phase:\n0.00401×275.7 ≈ 1.11 mAh g−1\nAccording to the reviewer’s constructive suggestions, we have revised the relevant parts on\npage 9 of the revised manuscript and pages 48-50 of the revised supplementary information.\nWe copy the relevant text here for your check:\nNew Supplementary Figure 45\nSupplementary Fig. 45 | Schematic illustration of a typical spherical NCM particle model\n(radius set to 5 μm), in which a representative 10 nm-thick surface shell region affected by the\nflash Joule heating treatment is delineated based on our experimental observations. We defined\nthree different compositions for the surface shell including the original layered structure, spinel\nphase and rock-salt phase, and for simplicity, the coexistence of multiple phases was not\nconsidered in the calculation.\nThe above model is used to semi-quantitatively estimate lithium loss during flash Joule\nprocessing.\nStep 1: The key parameters used in the calculations are as follows:\n Radius of particle: r=5 μm=5000 nm\n Thickness of Li-deficient shell: t=10 nm\n Volume of the particle:\n4 4\n𝑉 = 𝜋𝑅3 = ×3.14×(5000)3 = 5.236×1011 nm3\n𝑡𝑜𝑡𝑎𝑙 3 3\n Volume of the Li-deficient shell:\n4 4\n𝑉 = 𝜋𝑅3 − 𝜋(𝑅 −𝑡)3 = 3.136×109 nm3\n𝑠ℎ𝑒𝑙𝑙 3 3\n The proportion of lithium in the 10 nm surface layer:\n𝑉 3.136×109\n𝑠ℎ𝑒𝑙𝑙\n𝑓 = = ≈ 0.6%\n𝑉 5.236×1011\n𝑡𝑜𝑡𝑎𝑙\n The molar masses:\nM (LiNi Co Mn O ) = 97.28 g/mol\nLayered 0.8 0.1 0.1 2\nM (Li Ni Co Mn O , x=0.5) =93.81 g/mol\nSpinel x 0.8 0.1 0.1 2\nM (Li TMO, the phase occurred in XRD) =75.80 g/mol\nRock-salt 0.21\n The typical densities (obtained by the MDI Jade database):\nρ (LiNi Co Mn O ) = 4.74 g/cm³\nLayered 0.8 0.1 0.1 2\nρ (Li Ni Co Mn O , x=0.5) = 4.79 g/cm³\nSpinel x 0.8 0.1 0.1 2\nρ (Li TMO, the phase occurred in FJ700 and FJ800) = 5.85 g/cm³\nRock salt 0.21\n The number of lithium atoms per unit volume of different samples:\nρ 4.74\n𝑛 = NCM811 𝑁 = ×6.022×1023 ×10−21 = 29.35 atoms/nm3\nLi,Layered M 𝐴 97.28\nNCM811\nρ 4.79\n𝑛 = spinel 𝑁 = ×6.022×1023 ×10−21 ×0.5\nLi,Spinel M 𝐴 93.81\nspinel1\n= 15.37 atoms/nm3\nρ 5.85\n𝑛 = rock salt 𝑁 = ×6.022×1023 ×10−21 ×0.21\nLi,Rock−salt M 𝐴 75.80\nrock salt\n= 9.76 atoms/nm3\nStep 2: When the surface undergoes a layered-to-spinel phase transformation\n Li atoms in the layered surface:\n𝑁 = 𝑛 ×𝑉 = 29.35×3.136×109 = 9.20×1010\nLi,shell,Layered Li,NCM811 𝑠ℎ𝑒𝑙𝑙\n Li atoms in the spinel surface:\n𝑁 = 𝑛 ×𝑉 = 15.37×3.136×109 = 4.82×1010\nLi,shell,spinel Li,spinel 𝑠ℎ𝑒𝑙𝑙\n Relative lithium loss:\n9.20−4.82\n×0.6% ≈ 0.286%\n9.20\n The reduced capacity caused by lithium loss in spinel phase:\n0.00286×275.7 (Theoretical capacity of NCM811) ≈ 0.79 mAh g−1\nStep 3: When the surface undergoes a layered-to-rock-salt phase transformation\nLi atoms in the layered surface:\n𝑁 = 𝑛 ×𝑉 = 29.35×3.136×109 = 9.20×1010\nLi,shell,NCM811 Li,NCM811 𝑠ℎ𝑒𝑙𝑙\nLi atoms in the rock-salt surface:\n𝑁 = 𝑛 ×𝑉 = 9.76×3.136×109 = 3.06×1010\nLi,shell,rock−salt Li,rock−salt 𝑠ℎ𝑒𝑙𝑙\nRelative lithium loss:\n9.20−3.06\n×0.6% ≈ 0.401%\n9.20\nThe reduced capacity caused by lithium loss in rock-salt phase:\n0.00401×275.7 ≈ 1.11 mAh g−1\nSupplementary Table S1 Semi-Quantitative calculation of Li inventory loss and its correlation\nwith ICE for flash Joule-heated LiNi Co Mn O samples.\n0.8 0.1 0.1 2\n10 nm shell Estimated Li Equivalent capacity\nICE Capacity\nstructure inventory loss (%) loss (mAh g-1)\n95.3%\nSpinel phase 214.5 0.286 0.79\n(FJ350)\nRock-salt 76.0%\n154.5 0.401 1.11\nphase (FJ600)\n“According to the semi-quantitative estimation of FJ600 in Table S1, this deterioration\nprimarily arises from the thick rock-salt reconstruction layer, which blocks Li⁺ transport and\nexacerbates interfacial deterioration, rather than from the negligible lithium extraction during\nthe flash Joule heating process.”\n“Given the above, even if the outermost 10 nm-thick surface region of the spherical NCM811\nparticles were to fully reconstruct into Li-deficient phases (such as spinel or rock-salt),\\ the\nresulting loss of total lithium inventory would be only 0.286~0.401% (≈ 0.79~1.11 mAh g-1).\nConversely, instead of exhibiting a capacity decrease, our FJ350 sample delivers higher\nreversible capacity (214.5 mAh g-1) and ICE (95.3%), owing to the suppression of lattice\noxygen loss and interfacial side reactions.\nIn stark contrast, the FJ600 electrode delivers a very low reversible capacity of 154.5 mAh\ng-1, corresponding to a low ICE of 76.0%. Clearly, the substantial irreversible capacity loss of\nup to 60 mAh g-1 cannot be explained by the trace lithium extraction that occurs during the\nflash Joule heating process, as there is a significant discrepancy between the two. Therefore,\nthe drastic decrease in the ICE of the FJ600 electrode does not originate from the limited\nlithium inventory loss during the flash Joule heating process, but rather from structural and\ninterfacial effects, most notably, the formation of a thicker rock-salt reconstruction layer, which\nincreases charge-transfer resistance and hinders Li⁺ transport. These factors collectively\nconsume the cyclable lithium and intensify polarization, thereby reducing both the reversible\ncapacity and the initial coulombic efficiency of FJ600.”\nQ3: Contrary to the coin cell cycle evaluation results, the discharge capacity of FJ350 in pouch\ncells is less uniform than that of pristine. Is this due to the inhomogeneity caused by the FJ\ntreatment, which processes powder in 5-10 g quantities? Comments are needed on the\ninhomogeneity in these data.\nOur response: We sincerely appreciate the reviewer’s insightful comment. Regarding the\nfluctuations in the discharge capacity of the FJ350 samples in the pouch cell, we would like to\nclarify that these fluctuations are not due to material inhomogeneities introduced by the flash\nJoule processing. The observed fluctuations are associated with uncontrollable variations in\nambient temperature, as the long-term cycling tests of the pouch cells were conducted without\ntemperature regulation (Figure R4a). Fluctuations in environmental temperature may\ninfluence the viscosity of the electrolyte and the reaction kinetics, thereby leading to the\nobserved capacity variations.\nTo confirm the effect of ambient temperature changes on capacity fluctuation, we further\nevaluated the electrochemical performance of the pouch cells under controlled temperature\nconditions (Figure R4b, 26 °C). As Figure R4c demonstrates, the pouch cell exhibited no\nnoticeable capacity fluctuation, and a capacity retention of 99.3% was achieved after 200 cycles.\nThese results underscore the excellent electrochemical stability and uniformity of the prepared\nFJ350.\nFigure R4 | (a) Photograph of the pouch cell testing setup under ambient temperature\nconditions. (b) Temperature-controlled chamber built for temperature-regulated testing of the\npouch cell. (c) Cycling performance of the FJ350 pouch cell at 26 °C at 1 C within a voltage\nwindow of 2.7-4.3 V.\nIt is also worth emphasizing that both the pulse temperature and pulse width in our\nself-made setup are precisely controlled. Moreover, the uniform thermal field can be\nachieved over areas of 10 × 10 cm2 (Supplementary Fig. 1) and even up to 50 × 50 cm2.\nWithin this uniform temperature field, the reproducibility and consistency of samples treated\nin a single batch are excellent. Therefore, a batch processing amount of 5~10 g does not lead\nto any perceptible inhomogeneity.\nQ4: In Fig. S27, the authors present the charge transfer resistance of each electrode through\nEIS analysis. The Nyquist plot used to represent these Rct values is also shown, and it is\nrecommended that all other resistance values be displayed in addition to the Rct values.\nOur response: We sincerely appreciate the reviewer’s valuable suggestion. According to the\nreviewer’s comments, we have additionally analyzed the CEI film resistance (R) of each\nf\nelectrode, and the results have been added to Supplementary Figure 34. As shown in the\nupdated figure, the R value of the treated sample is significantly lower than that of the pristine\nf\none, indicating that the spinel coating layer effectively reduces the surface film resistance.\nThis improvement can be attributed to the enhanced interfacial stability and faster ion/electron\ntransport pathways provided by the spinel-phase layer, which together facilitate charge transfer\nand overall electrochemical performance.\nAccording to the reviewer’s constructive suggestions, we have revised the relevant part on\npage 17 of the revised manuscript. We copy the relevant text here for your check:\n“As manifested in Supplementary Fig. 34, the FJ350 electrode exhibits markedly lower CEI\nfilm resistance (R) and charge transfer resistance (R ) at all stages of cycling[42]. The\nf ct\npronounced reduction in R confirm that the spinel-phase skin establishes a percolating\nct\nelectron pathway, synergizing with fast Li+ transport to achieve high rate capability and cycling\ncapacity retention.”\nNew Supplementary Figure 34\nSupplementary Fig. 34 | EIS spectra at selected cycle numbers for the (a) pristine and (b)\nFJ350 electrodes. (c) Comparison of the fitted R values of the pristine and FJ350 electrodes\nf\nduring cycling. (d) Comparison of the fitted R values of the pristine and FJ350 electrodes\nct\nduring cycling.\nQ5: In Figure 5a, the authors note that in the highly delithiated state of the pristine sample,\nthe electron density near the Fermi level decreases sharply, suppressing the redox activity of\noxygen and nickel, leading to oxygen release and structural collapse. The authors should\nprovide specific references supporting these claims.\nOur response: Thank you for your valuable comment on the connection between the electronic\ndensity of states near the Fermi level and the redox behavior of oxygen and nickel. We\nacknowledge that relying solely on DOS changes is insufficient to fully explain redox\nsuppression, and we have revised our discussion accordingly.\nAs recommended, we conducted additional LOBSTER calculations to evaluate the crystal\norbital Hamilton population (COHP) and integrated COHP (ICOHP) for Ni-O bonds. As\nrevealed in Supplementary Fig. 43, the results indicate that in the fully lithiated state, the Ni-\nO bonds in the LiTMO @LiTM O heterostructure show stronger bonding interactions\n2 2 4\n(higher-ICOHP values) than those in bare LiTMO . Given that LiTM O generally retains\n2 2 4\npartial lithium during cycling, the spinel-phase skin contributes to improved framework\nstability. This reinforced bonding helps suppress oxygen release and irreversible phase\ntransitions under high-voltage conditions, thereby enhancing structural integrity and redox\nreversibility.\nConcerning electronic structure, the decreased DOS near the Fermi level in the highly\ndelithiated pristine sample does reflect reduced electronic conductivity. However, as rightly\nnoted, redox activity particularly anionic redox depends not only on states at the Fermi level\nbut also on the bonding/anti-bonding characteristics across a broader energy range. Our COHP\nanalysis confirms that the pristine sample under deep delithiation exhibits stronger Ni-O anti-\nbonding interactions near the Fermi level, which aligns with greater oxygen instability. In\ncontrast, the LiTMO @LiTM O heterostructure maintains a more stable electronic\n2 2 4\nenvironment with moderated anti-bonding character, effectively inhibiting oxygen oxidation\nand loss.\nNew Supplementary Figure 43\nSupplementary Fig. 43 | The crystal orbital Hamilton population (COHP) and integrated\nCOHP (ICOHP) for TMO , LiTMO , TMO @TM O , and LiTMO @LiTM O .\n2 2 2 2 4 2 2 4\nAccording to the reviewer’s constructive suggestions, relevant information have been\nprovided on page 21 of the revised manuscript and page 46 of the revised supplementary\ninformation. We copy the relevant text here for your check:\n“To further understand the redox behavior and structural stability, we analyzed the crystal\norbital Hamilton population (COHP) of Ni-O bonds (Supplementary Fig. 43). The integrated\nCOHP (−ICOHP) values indicate that in the fully lithiated state, the Ni-O bonds in\nLiTMO @LiTM O possess a stronger bonding strength (−ICOHP = 0.803 eV) than those in\n2 2 4\npristine LiTMO (−ICOHP = 0.761 eV). In the delithiated state, both TMO and\n2 2\nTMO @TM O show comparable Ni-O bond strengths. Given that LiTM O generally retains\n2 2 4 2 4\nsome Li+ during cycling, the LiTMO @LiTM O heterostructure helps maintain robust bond\n2 2 4\nstability throughout charge and discharge, thereby reinforcing the host framework and\nsuppressing oxygen loss as well as phase degradation. This effect contributes to the inhibition\nof irreversible oxygen and nickel redox activity. In contrast, the pristine sample exhibits a\npronounced decrease in the electronic density of states near the Fermi level under high degrees\nof delithiation, along with lower −ICOHP values in the lithiated state. These features imply\ninferior electronic transport and weaker Ni-O bonding during cycling, which can trigger\noxygen release and structural collapse, ultimately resulting in rapid capacity fading. These\ntheoretical insights are consistent with experimental observations, i.e., the spinel-coated\nLiTMO @LiTM O maintains high structural stability even under high-voltage conditions,\n2 2 4\neffectively suppressing oxygen release and irreversible phase transitions.”\nQ6: It is stated that the in-situ XRD results support the claim that the spinel-phase “skin”\nreduces the coexistence of the H2 and H3 phases, facilitates reversible structural formation,\nand promotes uniform lithium distribution. However, the rationale for comparing these results\nwith TOF-SIMS at the 4.3 V cutoff is unclear. It is recommended to modify Fig3. (e, f) to\ncorrespond with Fig3. (h, i), enabling a direct comparison of results obtained at the same cutoff\nvoltage.\nOur response: We sincerely appreciate the reviewer’s insightful comment. We fully\nunderstand the reviewer’s concern regarding the comparison between Figures 3e, 3f and\nFigures 3h, 3i, and we would like to provide a more detailed clarification of our experimental\nrationale and data presentation.\nIn this study, the majority of the comparative analyses, including TEM, TOF-SIMS were\nconducted at an upper cutoff voltage of 4.3 V to ensure a controlled and direct comparison\nunder standard conditions, as the reviewer rightly pointed out. However, for the in situ XRD\nexperiment, we intentionally extended the upper cutoff voltage to 4.7 V. The primary\nrationale was to subject the materials to a more stringent condition. We hypothesized that the\nbeneficial effects of the spinel-phase “skin”, particularly in suppressing phase coexistence and\nfacilitating structural reversibility, would be more pronounced and discernible under such harsh\nconditions. This approach allows for a more rigorous validation of the structural stability\nafforded by the surface modification. The in situ XRD results convincingly demonstrate that\nthe spinel-phase “skin” effectively mitigates these detrimental structural changes, which is a\ncentral claim of our work.\nAccording to the reviewer's constructive suggestions, we have revised the relevant part on\npage 16 of the revised manuscript. We copy the relevant text here for your check:\n“To gain deeper insight into how the spinel-phase “skin” influences the structural evolution\nand Li-ion transport behavior of NCM811 under more demanding conditions, in situ XRD\nanalysis was additionally carried out at an upper cutoff voltage of 4.7 V.”\nQ7: To support the claim of spinel phases of suitable thickness, it would be helpful to add\nnumerical values of spinel layer thickness in addition to cycle data and lattice spacing.\nOur response: We sincerely appreciate the reviewer’s valuable suggestion. In response, the\nthickness of the spinel layer for each treated sample has been carefully measured and\nsummarized in the table below for clarity. The numerical values represent the average thickness\nobtained from TEM observations. Note that this represents only a semi-quantitative statistical\naverage.\nNotably, the thickness of the spinel-phase layer formed under different treatment conditions\n(pulse temperature and pulse width) varies systematically, generally following the trend of\nincreasing thickness with longer pulse duration or higher treatment temperature (as shown in\nthe new Table S2 and Supplementary Fig. 15). Meanwhile, the thickness of the spinel-phase\nshell, as well as the associated degree of lithium deficiency, exerts a pronounced influence on\nthe electrochemical performance including ICEs and cycling stability. These observations\nfurther support that an optimally formed spinel layer, both in thickness and defect chemistry,\ncontributes to the superior cycling stability of the FJ350 sample.\nNew Supplementary Figure 15\nSupplementary Figure 15 | Average thickness of the spinel-phase shell formed on the surface\nof NCM811 under different treatment conditions.\nNew Table S2\nSupplementary Table S2 Summary of the spinel-phase layer thickness and electrochemical\nperformance of NCM811 cathodes treated under different flash sintering conditions.\nSpinel-phase\nConditions Treatment Cathodes ICEs Cyclability\nlayer thickness\n/ / Pristine 0 nm 89.6% 80.3%, 200th\n300 °C-30 s FJ300 6.5±0.5 nm 91.2% 88.3%, 200th\nVarying the\ntemperature at 350 °C-30 s FJ350 7.5±0.5 nm 95.3% 99.2%, 200th\na fixed time\n400 °C-30 s FJ400 8.5±0.5 nm 92.3% 92.3%, 200th\nVarying the 350 °C-15 s 350-15 s 5.5±0.5 nm 93.6% 89.2%, 200th\ntime at a fixed\ntemperature 350 °C-60 s 350-60 s 10±0.5 nm 90.8% 90.0%, 200th\nAccording to the reviewer’s constructive suggestions, relevant information have been\nprovided on page 11 of the revised manuscript and pages 18 and 51 of the supplementary\ninformation. We copy the relevant text here for your check:\n“Based on the TEM analysis, we performed semi-quantitative measurements and statistical\nanalysis of the average thickness of the spinel-phase shell formed on the surface of NCM811\nunder different flash-sintering conditions (Supplementary Fig. 15), and correlated these\nresults with the corresponding electrochemical performance. As summarized in Table S2, the\nthickness of the spinel-phase layer formed under different treatment conditions (pulse\ntemperature and pulse width) varies systematically, generally following the trend of increasing\nthickness with longer pulse duration or higher treatment temperature. Meanwhile, the thickness\nof the spinel-phase shell, as well as the associated degree of lithium deficiency, exerts a\npronounced influence on the electrochemical performance including ICEs and long-term\ncycling stability. These findings further demonstrate that an optimally engineered spinel layer,\nwith well-balanced thickness and defect chemistry, enables the FJ350 sample to deliver the\nmost superior electrochemical performance among all treated conditions.”\nReviewer #3 (Remarks to the Author): ================================\nThe manuscript introduces a subtractive strategy to engineer a conformal spinel-phase “skin”\non Ni-rich layered cathode materials (e.g., NCM811, NCM90) via flash Joule heating (FJH).\nUnlike conventional additive coatings, this approach selectively strips surface Li and O atoms,\nforming a coherent spinel layer with tunable thickness. The authors demonstrate that this spinel\n“skin” provides robust interfacial bonding, improved ionic/electronic conductivity, and strong\nsuppression of lattice oxygen loss. However, the main novel point associated with the DFT\ncalculation should be majorly revised.\nOur response: We are grateful to the reviewer for the positive and professional comments. All\nrequested revisions based on the constructive comments from the reviewer have been fully\nconsidered and thoroughly addressed, as shown below.\nQ1: “In contrast, samples with either too thick or too thin spinel-phase ‘skin’ coatings (FJ300\nand FJ400), although performing better than the pristine sample, are still inferior to FJ350.\nOn one hand, the excessively thin shell lacks sufficient mechanical strength and isolation\neffect”. How do the authors measure the coverage of Li depletion in comparison to the coating\nmethod, and then conclude its supremacy? Data needs to be presented.\nOur response: Thank you for your valuable comments on our manuscript. In response to your\nquery regarding the impact of coating thickness on irreversible lithium depletion and coating\neffectiveness, we have conducted a detailed analysis supported by theoretical calculations and\nexperimental data.\nFirst, from a theoretical perspective, as displayed in Figure 5e, we observe that when the\nspinel material is in a high delithiation state, the coating layer significantly enhances the lithium\nadsorption capability of LiNi Co Mn O . However, as the lithium content increases, the\n0.8 0.1 0.1 2\nlithium adsorption capacity of the coating gradually decreases. Specifically, in a lithium-rich\nstate, an excessively thick coating may hinder lithium-ion diffusion, thereby limit effective\nutilization of lithium ions and ultimately exacerbate irreversible lithium depletion. Therefore,\nan optimal coating thickness helps to maintain smooth lithium-ion diffusion and prevents\nexcessive lithium depletion, optimizing battery performance.\nFrom an experimental perspective, the initial Coulombic efficiency (ICE) results reveal that\nthe pristine sample exhibits an ICE of 89.6%, while the thinly coated FJ300 sample reaches\n91.2%. In contrast, FJ350 achieves a markedly higher ICE of 95.3%, whereas FJ400 records a\nslightly lower value of 92.3%. As stated in our manuscript, the improvement in ICE provided\nby the spinel-phase shell primarily originates from two factors: 1) the spinel-phase shell\neffectively isolates the electrode surface from the electrolyte, thereby greatly suppressing side\nreactions; and 2) it mitigates oxygen loss and structural degradation during cycling. Notably,\nthe surface coverage of the spinel-phase shell has a direct influence on the ICE. Therefore, the\nICE can be regarded as an important indirect indicator of the coating coverage. Clearly,\nbased on this criterion, the FJ300 sample with an ICE of 91.2% still exhibits incomplete\nsurface coverage. Conversely, the lower ICE of FJ400 compared with FJ300 arises from the\nexcessively thick spinel shell, which hinders Li+ transport and induces polarization. In this case,\nthe decrease in ICE is no longer attributable to insufficient surface coverage. The spinel shell\non the surface of the FJ350 sample not only exhibits excellent completeness but also\npossesses an appropriate thickness, thereby delivering the best electrochemical performance.\nIn addition, the importance of an optimized shell layer for NCM811, both in the first cycle\nand during long-term cycling, has been thoroughly discussed in the manuscript (Figure 3d and\nFigure 4b). However, the nanoscale spinel-phase shell cannot be directly observed by SEM,\nand TEM can only provide local structural information, making it impossible to obtain an\noverall assessment of the coating coverage.\nIn this case, we chose to employ in situ differential electrochemical mass spectrometry\n(DEMS) to assess the coverage of the spinel-phase shell by monitoring oxygen evolution\nat high voltages (e.g., 4.7 V) for different samples. Although this method is also indirect,\nthe amount of released oxygen provides a relatively objective measure of the completeness\nand confining effect of the shell layer. It should be noted that the upper cutoff voltage used\nhere represents an extremely demanding condition; this is because under more conventional\nupper voltages, such as 4.3 V, the oxygen evolution from all samples is very low, making\nit difficult to distinguish the advantages of a highly covered shell layer.\nAs shown in Supplementary Fig. 13, the FJ300 sample with a thin and incomplete spinel-\nphase shell already exhibits a certain degree of suppression of oxygen evolution during the first\ncharge compared with the pristine sample, but the reduction is limited, decreasing only from\n4.6 to 3.1 nmol min-1. In sharp contrast, the oxygen-release signals of the FJ350 and FJ400\nelectrodes are significantly lower, with FJ350 in particular reduced to as low as 1.2 nmol\nmin-1. This result indicates that the spinel-phase shell can effectively inhibit oxygen evolution\neven under the extremely high cutoff voltage of 4.7 V, directly reflecting both the completeness\nand the strong spatial confinement effect of the FJ350 coating layer. The suppression of oxygen\nrelease is also consistent with the increased Coulombic efficiency and reduced irreversible\nlithium consumption observed in the electrochemical tests, further demonstrating the critical\nrole of an optimized coating thickness in achieving stable cycling performance.\nGiven the above, the FJ350 coating not only delivers superior first-cycle and long-term\ncycling performance, but is also substantiated by theoretical calculations and DEMS analyses\nto form a fully covered, optimally thick shell. Such an engineered shell effectively suppresses\noxygen release and interfacial side reactions, while simultaneously enhancing lithium-ion\ndiffusion efficiency.\nNew Supplementary Figure 13\nSupplementary Figure 13 | In situ DEMS data of (a) Pristine, (b) FJ300, (c) FJ350, (d) FJ400\nduring the first charge within 2.7-4.7 V.\nAccording to the reviewer’s constructive suggestions, relevant information have been\nprovided on page 9 of the revised manuscript and page 16 the supplementary information. We\ncopy the relevant text here for your check:\n“In situ differential electrochemical mass spectrometry (DEMS) was employed to monitor the\noxygen evolution of different samples under high voltages (e.g., 4.7 V) in order to evaluate the\ncoverage of the spinel-phase shell. As Supplementary Fig. 13 demonstrates, the FJ300 sample\nwith a thin and incomplete spinel-phase shell already exhibits a certain degree of suppression\nof oxygen evolution during the first charge compared with the pristine sample, but the reduction\nis limited, decreasing only from 4.6 to 3.1 nmol min-1. In sharp contrast, the oxygen-release\nsignals of the FJ350 and FJ400 electrodes are significantly lower, with FJ350 in particular\nreduced to as low as 1.2 nmol min-1. This result indicates that the spinel-phase shell can\neffectively inhibit oxygen evolution even under the extremely high cutoff voltage of 4.7 V,\ndirectly reflecting both the completeness and the strong spatial confinement effect of the FJ350\ncoating layer. The suppression of oxygen release is also consistent with the increased\nCoulombic efficiency and reduced irreversible lithium consumption observed in the\nelectrochemical tests, further demonstrating the critical role of an optimized coating thickness\nin achieving stable cycling performance.”\nQ2: “The advantage of FJ350 over the pristine sample in reaction kinetics can be elucidated\nby measuring the Li-ion diffusion coefficient (D +)”. Here the authors are comparing their\nLi\nnovel processed cathode with the pristine one. However, the novelty of this point should be a\nbetter coating method. Thus, the authors should compare properties between coated NCM and\nself-induced coated NCM, not just with the pristine one.\nOur response: We are very grateful for the highly professional and important suggestions from\nthe reviewers. To further highlight the electrochemical advantages of the self-induced coated\nFJ350, we conducted a comparative study on the electrochemical performance and DEMS\nresults of FJ350, commercial LiNbO -coated LiNi Co Mn O (LiNi Co Mn O\n3 0.8 0.1 0.1 2 0.8 0.1 0.1 2\n@LiNbO ), and commercial Li BO -coated LiNi Co Mn O (LiNi Co Mn O @B).\n3 3 3 0.8 0.1 0.1 2 0.8 0.1 0.1 2\nAs manifested in Supplementary Fig. 23, compared with FJ350 (which delivers an initial\ncoulombic efficiency of 95.3% and a discharge capacity of 210 mAh g-1), the\nLiNi Co Mn O @LiNbO cathode exhibits a significantly lower initial coulombic\n0.8 0.1 0.1 2 3\nefficiency of only 86.1%, along with a discharge capacity of 199.8 mAh g-1. Similarly,\nSupplementary Fig. 24 demonstrates that the LiNi Co Mn O @B material also\n0.8 0.1 0.1 2\nunderperforms relative to FJ350, showing a first discharge capacity of only 198.2 mAh g-1,\ncorresponding to a low Coulombic efficiency of 85.2%. These results further highlight the\nsuperior electrochemical performance of FJ350. In terms of cycling stability, after 200 cycles\nat 1 C, the FJ350 electrode still maintains a high capacity retention of 99.2%, whereas\nLiNi Co Mn O @LiNbO and LiNi Co Mn O @B exhibit much lower retentions of\n0.8 0.1 0.1 2 3 0.8 0.1 0.1 2\n82.7% and 82.2%, respectively.\nMoreover, the DEMS results show that the oxygen release from FJ350 (1.2 nmol min-1) is\nmarkedly lower than that from LiNi Co Mn O @LiNbO and LiNi Co Mn O @B,\n0.8 0.1 0.1 2 3 0.8 0.1 0.1 2\nwith the latter two reaching oxygen evolution values of 2.67 and 3.46 nmol min-1 at 4.7 V,\nrespectively. This once again confirms the higher structural stability of FJ350, strongly\nsupporting the effectiveness of the proposed subtraction-coating strategy in enhancing battery\nsafety and performance.\nTo further highlight the advantages of the subtraction-coating strategy, we also compared\nFJ350 with NCM811 cathodes coated using traditional wet-chemical methods as well as\nother typical methods reported in the literature. As shown in Table S4, FJ350 exhibits\nremarkable improvements in both cycling performance and rate capability, underscoring that\nthe subtraction-coating strategy not only enhances electrochemical performance but also\nimproves battery stability and extends service life.\nNew Table S4\nSupplementary Table S4 Comparison of the initial Coulombic efficiency and cycling stability\nof FJ350 with NCM811 cathodes prepared using various coating methods reported in the\nliterature.\nCathodes ICEs Retention Ref.\nFJ350 (this work) 95.3% 95.3%, 500th (1 C) This work\nNCM811@In O &LiInO 90.9% 92.7%, 100th (1 C) [19]\n2 3 2\nNCM811@SPDX 91.4% 82.5%, 200th (1 C) [20]\nNCM811@Co B 91.5% 82.2%, 200th (1 C) [21]\nx\nNCM811@LiTMPO 81.3% 92.0%, 100th (1 C) [22]\nN-FCG-NCM811@LiNbO 89.4% 77.2%, 300th (1 C) [25]\n2 2 9\nNew Supplementary Figure 23\nSupplementary Fig. 23 | The electrochemical performance of the commercial\nLiNi Co Mn O @LiNbO cathode. (a) The Initial charge-discharge profiles at 0.1 C. (b)\n0.8 0.1 0.1 2 3\nCycling performance over 200 cycles within 2.7-4.3 V. (c) In situ DEMS data during the 1st\ncharge cycle within 2.7-4.7 V. (d) GITT profile of the charge/discharge processes of and the\ncorresponding D + vs. potential curve.\nLi\nNew Supplementary Figure 24\nSupplementary Fig. 24 | The electrochemical performance of the commercial\nLiNi Co Mn O @B cathode. (a) The Initial charge-discharge profiles at 0.1 C. (b) Cycling\n0.8 0.1 0.1 2\nperformance over 200 cycles within 2.7-4.3 V. (c) In situ DEMS data during the 1st charge cycle\nwithin 2.7-4.7 V. (d) GITT profile of the charge/discharge processes of and the corresponding\nD + vs. potential curve.\nLi\nAccording to the reviewer’s constructive suggestions, relevant information have been\nprovided on pages 12 and 13 of the revised manuscript and pages 26 and 27 the supplementary\ninformation. We copy the relevant text here for your check:\n“To further highlight the electrochemical advantages of the self-induced coated FJ350, we\nconducted comparative studies on the electrochemical performance and DEMS results of FJ350,\ncommercial LiNbO -coated LiNi Co Mn O (LiNi Co Mn O @LiNbO ), and\n3 0.8 0.1 0.1 2 0.8 0.1 0.1 2 3\ncommercial Li BO -coated LiNi Co Mn O (LiNi Co Mn O @B). As shown in\n3 3 0.8 0.1 0.1 2 0.8 0.1 0.1 2\nSupplementary Figs. 23 and 24, the commercial NCM811 cathodes coated with LiNbO or\n3 3\nalso inferior suppression of oxygen release under high voltage (4.7 V). In addition, we further\ncompared FJ350 with NCM811 cathodes coated using traditional wet-chemical methods and\nother representative coating approaches reported in the literature. As presented in\nSupplementary Table S4, the FJ350 electrode demonstrates remarkable improvements in both\ncycling performance and rate capability. These results collectively highlight that our\nsubtraction-coating strategy not only enhances electrochemical performance but also improves\nbattery stability and extends service life, making it the most effective approach among currently\navailable coating methods.”\nQ3: “This phenomenon indicates a substantial improvement in the electrical conductivity of\nmaterial, suggesting a superior rate capability compared to the pristine sample. In contrast, the\npristine sample demonstrates a sharp decrease in electronic density of states near the Fermi\nlevel at high delithiation states, indicating suppressed redox activity of oxygen and nickel.” I\ndo not agree with this statement. A sharp decrease in electronic density of states near the Fermi\nlevel cannot alone indicate suppressed redox activity of oxygen and nickel. The oxygen redox\nshould be considered not only from the electrons right at the Fermi level but also from the\nregion close to it. In this sense, I cannot see much difference between Fig. 5 (a) top and bottom.\nThe authors may use LOBSTER to accurately calculate this.\nOur response: Thank you for your insightful comment regarding the relationship between\nelectronic density of states (DOS) near the Fermi level and the redox activity of oxygen and\nnickel. We agree that a sharp decrease in DOS near the Fermi level alone cannot fully\ndemonstrate suppressed redox activity, and we apologize for the oversimplified statement in\nour original manuscript. As suggested, we have performed additional LOBSTER calculations\nto accurately analyze the crystal orbital Hamilton population (COHP) and integrated COHP\n(ICOHP) for the Ni-O bonds. As showed in Supplementary Fig. 43, the results reveal that in\nthe fully lithiated state, the Ni-O bonds in the LiTMO @LiTM O heterostructure exhibit\n2 2 4\nstronger bonding interactions (larger –ICOHP values) compared to the uncoated LiTMO .\n2 4\ncomplete delithiation, the spinel-phase skin enhances the mechanical and chemical stability of\nthe framework. This strengthening of the Ni–O bonds effectively suppresses oxygen release\nand irreversible phase transitions during high-voltage operation, thereby improving the\nstructural integrity and redox reversibility.\nRegarding the electronic structure, the reduced DOS near the Fermi level at high delithiation\nstates in the pristine sample indeed indicates decreased electronic conductivity. However, as\nthe reviewer rightly pointed out, redox activity, especially anionic redox, is influenced not only\nby states exactly at the Fermi level but also by the overall bonding/anti-bonding character of\noxygen-related bands in a broader energy range. Our COHP analysis confirms that in the\npristine sample under high delithiation, the Ni-O anti-bonding states near the Fermi level are\nmore pronounced, which correlates with higher oxygen instability. In contrast, the\nLiTMO @LiTM O heterostructure maintains a more stable electronic structure with\n2 2 4\nmoderated anti-bonding interactions, contributing to the inhibition of oxygen oxidation and\nloss.\nNew Supplementary Figure 43\nSupplementary Fig. 43 | The crystal orbital Hamilton population (COHP) and integrated\nCOHP (ICOHP) for TMO , LiTMO , TMO @TM O , and LiTMO @LiTM O .\n2 2 2 2 4 2 2 4\nAccording to the reviewer’s constructive suggestions, relevant information have been\nprovided on page 21 of the revised manuscript. We copy the relevant text here for your check:\n“The PDOS results reveal a notable increase in the electronic density near the Fermi level for\nthe coated material, along with a narrowed band gap and a reduced electronic transition barrier\n(Figure 5a). These electronic structure modifications suggest a marked enhancement in\nelectrical conductivity, which is expected to improve the rate performance relative to the\nuncoated sample.”\n“To further understand the redox behavior and structural stability, we analyzed the crystal\norbital Hamilton population (COHP) of Ni-O bonds (Supplementary Fig. 43). The integrated\nCOHP (−ICOHP) values indicate that in the fully lithiated state, the Ni-O bonds in\nLiTMO @LiTM O possess a stronger bonding strength (−ICOHP = 0.803 eV) than those in\n2 2 4\npristine LiTMO (−ICOHP = 0.761 eV). In the delithiated state, both TMO and\n2 2\nTMO @TM O show comparable Ni-O bond strengths. Given that LiTM O generally retains\n2 2 4 2 4\nsome Li+ during cycling, the LiTMO @LiTM O heterostructure helps maintain robust bond\n2 2 4\nstability throughout charge and discharge, thereby reinforcing the host framework and\nsuppressing oxygen loss as well as phase degradation. This effect contributes to the inhibition\nof irreversible oxygen and nickel redox activity. In contrast, the pristine sample exhibits a\npronounced decrease in the electronic density of states near the Fermi level under high degrees\nof delithiation, along with lower −ICOHP values in the lithiated state. These features imply\ninferior electronic transport and weaker Ni-O bonding during cycling, which can trigger\noxygen release and structural collapse, ultimately resulting in rapid capacity fading. These\ntheoretical insights are consistent with experimental observations, i.e., the spinel-coated\nLiTMO @LiTM O maintains high structural stability even under high-voltage conditions,\n2 2 4\neffectively suppressing oxygen release and irreversible phase transitions.”\nQ4: I cannot agree with the NEB calculation. Only the ODH migration path is considered. In\nlayered oxide cathode materials, the TSH is regarded as the common diffusion pathway as it\nhas ~1/3 Ea of ODH. This is well studied and reported by Ceder et al. Please refer to:\nhttps://doi.org/10.1103/PhysRevB.79.014305\nhttps://doi.org/10.1021/acsami.3c14444 “Theoretical methods: All spin-polarized xxxx”.\nOur response: We sincerely thank the reviewer for raising this important point regarding the\nLi+ diffusion pathways in layered oxide cathode materials. We agree that the tetrahedral site\nhop (TSH) mechanism is indeed the dominant diffusion pathway in layered oxides like LiCoO ,\n2\nas established in the previous works as you mentioned. In our original manuscript, we did\ncalculate the Li+ diffusion barriers along the TSH pathway, but we acknowledge that this\nwas not clearly presented. We have now revised the manuscript to explicitly state that our\nNEB calculations focused on the TSH mechanism. Our calculated diffusion barrier for TMO\n2\nis approximately 0.3 eV, which aligns well with values reported in the literature for similar\nsystems. To prevent any potential misunderstanding, in revised Figure 5e, we have added\nfigures showing both side and top views of the Li+ diffusion paths along the TSH\nmechanism. These illustrations clearly demonstrate the curved migration path through\ntetrahedral sites, which characteristic of the TSH mechanism.\nRegarding our model selection for the interface structure, we would like to clarify our\nrationale: 1) Unlike conventional coating methods where the coating layer is deposited\nexternally, the spinel phase in our work forms through an in-situ phase transition from the\nlayered structure. This would result in excellent orientation consistency between the two phases,\nparticularly along the layered direction, creating a well coherent interface. 2) We selected the\n4-coordinated facet as our model system because we believe that phase transition would\npreferentially occur on the less stable facets (J. Phys. Chem. C 2017, 121, 8290-8299). By\nimproving the stability of these vulnerable surfaces, we can significantly enhance the overall\nperformance, similar to the barrel effect where the shortest stave determines capacity. While\nother interface types exist in practice, our chosen model is representative and provides\nqualitatively sound conclusions for this study. We have incorporated these clarifications into\nthe revised manuscript, along with appropriate references to the works by Ceder et al. These\nrevisions ensure that our methodology is clearly presented and aligned with the established\nunderstanding of Li+ diffusion in layered oxide materials.\nAccording to the reviewer’s constructive suggestions, relevant information have been\nprovided on pages 23, 24 and 26 of the revised manuscript (Refs [47] and [48]). We copy the\nrelevant text here for your check:\n“Based on the analysis of average lithium ion adsorption energy in LiTMO and LiTM O ,\n2 2 4\nthe spinel phase exhibits a significantly stronger adsorption capability than LiTMO throughout\n2\nthe charge/discharge process, especially under lithium-rich conditions (Figure 5d). This\nenhanced adsorption behavior is closely related to Li⁺ diffusion kinetics. In line with\nestablished studies on layered oxide cathodes[47,48], we employed the CI-NEB method to\nevaluate the tetrahedral site hop (TSH) mechanism, which is the dominant Li+ diffusion\npathway characterized by a relatively low activation energy. The computed diffusion barrier\nfor TMO is approximately 0.3 eV, consistent with reported values. For clarity, schematic\n2\nillustrations of the TSH pathway from both side and top views are provided. As shown in\nFigure 5e, in the high-delithiation state, the LiTMO region in the coated sample shows\n2\nimproved lithium adsorption compared to the uncoated material. However, when the spinel\ncoating is in a lithium-rich state, this enhancement diminishes with increasing lithium content.\nThis implies that an excessively thick LiTMO outer layer could impede Li+ transport into the\n2\nbulk and weaken the adsorption contribution from LiTMO . Therefore, an optimal thickness of\n2\nthe spinel coating is essential to enhance lithium adsorption (Figure 5d) while establishing an\nefficient Li+ diffusion channel (Figure 5e). Moreover, Li+ diffusion across the coating interface\nexhibits strongly asymmetric energy barriers (Figure 5f and 5g), suggesting that extraction of\nLi⁺ from LiTMO is energetically more favorable than insertion. Hence, rational design of the\n2\nthickness and microstructure of the LiTM O spinel layer is critical to improve structural\n2 4\nstability while minimizing adverse effects on lithium-ion transport kinetics. These findings are\nconsistent with the superior initial Coulombic efficiency, rate capability, and cycling\nperformance observed in sample FJ350, which possesses an optimally tuned spinel-phase\nthickness.”\n“Regarding our interface model, the spinel phase formed through in-situ transformation\nmaintains crystallographic continuity with the parent layered structure, particularly along the\nbasal plane direction. We focused on the 4-coordinated facets as model surfaces since they\nrepresent vulnerable sites where phase transition likely initiates. Stabilizing these critical\ninterfaces through our approach provides substantial improvement in overall performance.”\nQ5: For the DFT calculation, the authors should provide a comprehensive picture and data of\ntheir structure. If the structure is as short as given in Fig. 5, the bottom or top vacuum layer\nwould impact the calculation at the interface. The authors are recommended to refer to DFT\ninterface modeling papers:\nhttps://doi.org/10.1021/acsami.2c07560\nhttps://doi.org/10.1021/acsami.4c05737\nMoreover, how did the authors sample such Li-deficient structures? If they simply picked one\nstructure without careful sampling, and if that structure is not the ground state energy, then\nsuch DFT calculation data would be totally meaningless.\nOur response: We thank the reviewer for the insightful comments regarding our DFT\ncalculations. In our DFT calculations, the core interface model between LiTMO and LiTM O\n2 2 4\nwas constructed as a symmetric, periodic slab without vacuum layer. This approach ensures\nthat the interface is embedded in a bulk-like environment, eliminating artificial effects from\nvacuum boundaries. The atomic layers for both phases were made sufficiently thick to mimic\nbulk properties and avoid surface-induced distortions. The entire model was fully optimized\nwithout any constraints, allowing the interface to relax naturally. This unconstrained\noptimization resulted in excellent lattice matching between the two phases at the interface, with\nminimal strain and no significant atomic reconstruction. The vacuum layer was only introduced\nin separate molecular dynamics (MD) simulations for surface stability assessments, e.g.,\nevaluating oxygen evolution barriers, where a vacuum spacer is necessary to model isolated\nsurfaces. This methodology aligns with best practices in DFT interface modeling, as\nemphasized in references such as Zhou et al. (ACS Appl. Mater. Interfaces 2022,\n10.1021/acsami.2c07560) and Gao et al. (ACS Appl. Mater. Interfaces 2021,\n10.1021/acsami.0c19091), which highlight the importance of avoiding vacuum-induced\nartifacts in interfacial studies.\nNew Supplementary Fig. 42\nSupplementary Fig. 42 | DFT model on the role of the spinel-phase “skin”. (a) LiTMO\n2\nstructure. (b) LiTM O structure. (c) LiTMO @LiTM O structure. (d) surface of TMO\n2 4 2 2 4 2\nstructure. (e) surface of TM O structure.\n2 4\nTo address the Li-deficient configurations, we employed the Special Quasirandom Structure\n(SQS) method to generate statistically representative models of random Li-vacancy\ndistributions. The SQS approach ensures that the chosen structures exhibit energy minima\nwithin the random solid-solution regime, effectively approximating the ground state for a given\nLi concentration. For each Li content, the same SQS distribution was applied across different\nsystems to enable direct comparison of energies and diffusion barriers. This rigorous sampling\navoids the pitfall of selecting arbitrary, non-representative configurations and ensures that our\nDFT results reflect realistic, energetically favorable states.\nAccording to the reviewer’s constructive suggestions, relevant information have been\nprovided on pages 21 and 26 of the revised manuscript (Refs [56] and [57]). We copy the\nrelevant text here for your check:\n“Based on the structural model presented in Supplementary Fig. 42, we performed a\ncomparative analysis of the partial density of states (PDOS) for both pristine LiTMO (TM =\n0.8 0.1 0.1 2\nLiTM O under charged and discharged states.”\n2 4\n“The interface model was constructed as a symmetric slab without vacuum layers[56,57], with\nthick atomic regions to bulk-like conditions. Full geometry optimization was performed\nwithout constraints to minimize artificial strain. Li-deficient structures were generated using\nthe SQS method[58] to approximate random vacancy distributions, with consistent sampling\napplied for all comparative studies.”\nPoint-by-point Response to Review Comments\nReviewers' comments and our Response:\nReviewer #1 (Remarks to the Author): ================================\nThe authors did plenty of work and analysis to address my initial concerns, and the paper is\nwell organized and it is worthy to be published.\nOur response: We greatly appreciate the reviewer’s careful evaluation and encouraging\ncomments. According to the constructive feedback, our research has become more rigorous,\nrefined, and scientifically sound.\nReviewer #2 (Remarks to the Author): ================================\nAmong the questions raised, the theory of capacity loss driven by lithium loss cannot\nrealistically represent the behavior of actual particles, and the mismatch between the cutoff\nvoltage used in the in-situ XRD experiment and the TOF-SIMS results remains a limitation.\nHowever, we acknowledge that there are clear technical constraints that prevent us from fully\naddressing this point. We believe that our responses to all of the reviewer’s questions are\nappropriate.\nOur response: We sincerely thank the reviewer for the positive evaluation and constructive\ncomments. Although the lithium loss during the flash Joule heating process is irreversible, its\nimpact on the initial Coulombic efficiency and capacity is negligible, primarily because the\nloss occurs only within a few nanometers at the surface and subsurface regions. This has been\neffectively captured through semi-quantitative calculations as proposed in Figure S14,\nconfirming that the effect is minimal. The following statement has been added to the SI: The\nabove model is used to semi-quantitatively estimate lithium loss during flash Joule processing.\nIn Figure 4, we tested two cutoff voltages, 4.3 V and 4.7 V, ensuring that the XRD and TOF-\nSIMS results are consistent with each other.\nReviewer #3 (Remarks to the Author): ================================\nThe authors have successfully addressed the concern. I would recommend it as publish.\nOur response: We sincerely thank the reviewer for the thorough review and positive\nassessment. According to the constructive feedback, our research has become more rigorous,\nrefined, and scientifically sound."}, {"round": 2, "reviewer_comments": {"reviewer_1": "The authors did plenty of work and analysis to address my initial concerns, and the paper is well organized and it is worthy to be published.", "reviewer_2": "Among the questions raised, the theory of capacity loss driven by lithium loss cannot realistically represent the behavior of actual particles, and the mismatch between the cutoff voltage used in the in-situ XRD experiment and the TOF-SIMS results remains a limitation.\nHowever, we acknowledge that there are clear technical constraints that prevent us from fully addressing this point.\nWe believe that our responses to all of the reviewer’s questions are appropriate.", "reviewer_3": "The authors have successfully addressed the concern. I would recommend it as publish."}, "author_response": "Point-by-point Response to Review Comments\nReviewers' comments and our Response:\nReviewer #1 (Remarks to the Author): ================================\nThe authors did plenty of work and analysis to address my initial concerns, and the paper is\nwell organized and it is worthy to be published.\nOur response: We greatly appreciate the reviewer’s careful evaluation and encouraging\ncomments. According to the constructive feedback, our research has become more rigorous,\nrefined, and scientifically sound.\nReviewer #2 (Remarks to the Author): ================================\nAmong the questions raised, the theory of capacity loss driven by lithium loss cannot\nrealistically represent the behavior of actual particles, and the mismatch between the cutoff\nvoltage used in the in-situ XRD experiment and the TOF-SIMS results remains a limitation.\nHowever, we acknowledge that there are clear technical constraints that prevent us from fully\naddressing this point. We believe that our responses to all of the reviewer’s questions are\nappropriate.\nOur response: We sincerely thank the reviewer for the positive evaluation and constructive\ncomments. Although the lithium loss during the flash Joule heating process is irreversible, its\nimpact on the initial Coulombic efficiency and capacity is negligible, primarily because the\nloss occurs only within a few nanometers at the surface and subsurface regions. This has been\neffectively captured through semi-quantitative calculations as proposed in Figure S14,\nconfirming that the effect is minimal. The following statement has been added to the SI: The\nabove model is used to semi-quantitatively estimate lithium loss during flash Joule processing.\nIn Figure 4, we tested two cutoff voltages, 4.3 V and 4.7 V, ensuring that the XRD and TOF-\nSIMS results are consistent with each other.\nReviewer #3 (Remarks to the Author): ================================\nThe authors have successfully addressed the concern. I would recommend it as publish.\nOur response: We sincerely thank the reviewer for the thorough review and positive\nassessment. According to the constructive feedback, our research has become more rigorous,\nrefined, and scientifically sound."}]

Decision letter