Intraplanar percolation and interplanar bridge enables layered matrix for high-performance negative electrode

Domain: neuroscience
Published: 2026-02-10
Editorial outcome: minor_revision

Abstract
The rapid advancement of modern technologies has increasingly heightened the demand for rechargeable batteries that exhibit higher energy density, faster charging capabilities, and longer cycle life. However, current battery materials are unable to meet these requirements due to their inherent structural limitations. Therefore, it is essential to design a multifunctional battery material structure, including ample Li+ storage, rapid Li+ diffusion, and enhanced structural stability. Here we show a layered structure featuring interplanar bridges and intraplanar percolation that integrates the advantages of currently commercialized mainstream battery material structures. It enables efficient ion transport and structural stability, delivering high capacity/rate capability, long cycle life, and nearly zero-strain structural evolution over a wide temperature range. The versatility of such structural design principle is further confirmed by other materials, and this general principle of discovering and designing advanced battery materials should accelerate the development of next-generation rechargeable batteries.

Introduction
Lithium-ion batteries (LIBs) have shown significant success in powering portable electronic devices, renewable energy storage, and electric vehicles. However, numerous technical demands for the LIBs, which current batteries cannot meet1–4. The ultimate pursuit of future rechargeable batteries is high energy/power density, long cycle life, all-climate durability, and high safety. Currently, conversion-type and alloy-type battery materials with huge volume change and poor cycle life are difficult to meet these requirements, while the intercalation-type materials maintain relative stability with comparable volume change during lithiation/delithiation, showing great potential to develop versatile LIBs for future demands5–7.

Generally, two-dimensional (2D) layered materials in Fig. 1a, have garnered significant commercial interest owing to their notable advantages, including large reversible theoretical capacity and 2D transport pathways, such as graphite negative electrodes, layered LiNi MnCoO positive electrodes, and LiCoO . However, the 2D transport pathways in layered matrix are still hard to afford rapid Li+ diffusion, especially for low-temperature output, and their anisotropic volume changes would also lead to stress accumulation and poor cycling performance8–11. On the other hand, the spinel structure in Fig. 1b, with three-dimensional (3D) Li+ diffusion channels, significantly improves the Li+ diffusion coefficient, resulting in enhanced rate performance. However, excess interplane bridges between adjacent layers occupy too many sites for Li+ intercalation, resulting in the low theoretical capacity for high-energy applications. Among various reported electrode materials, Li [Li Ti ]O (0≤y≤1) stands out as a representative of spinel structure with a lattice expansion of less than 0.2% for long-cycle-life negative electrode materials, while the spinel structure also endows the material with low theoretical capacity (175 mAh g−1) and rather high operating potential of approximately 1.55 V vs. Li+/Li, both of these factors significantly constrain its energy density12,13.

Recently, plenty of efforts have been devoted to developing high-capacity and fast-charging negative electrode materials for LIBs. To realize this goal, researchers focus on searching for materials that maintain the layered structure with high capacity while constructing tunnels along the layers to realize fast ionic transportation by intraplanar percolation. In recent years, the shear ReO structures (e.g., Ti Nb O , Nb W O , Nb WO , and Ni Nb O ) and tungsten bronze structures (e.g., Nb W O and T–Nb O) as shown in Fig. 1c, have garnered considerable attention in the field of fast-charging negative electrode materials, and such structures not only facilitate rapid Li+ diffusivity, but also show enhanced intercalation-pseudocapacitive behavior14–18. Despite their advantages, neither the shear ReO structure nor the tungsten bronze structure demonstrate sufficient stability after deep charge/discharge cycling, and the huge cell volume expansion and phase transitions during Li+ insertion would result in a gradual accumulation of lattice strain, ultimately leading to attenuation during prolonged cycling19–21. Consequently, to develop future negative electrode materials for high-energy, fast-charging and long-cycling LIBs, it is essential to design material structures with versatile functions: vast Li+ storage sites, fast Li+ diffusivity and stable framework under deep and long cycling.

In this work, to develop high-capacity, fast-charging and long-cycling battery materials, we propose a strategy utilizing a layered structure with interplanar bridge and intraplanar percolation by combining the merits of the above-mentioned material structures as shown in Fig. 1d. In such structure, the active transition metal (TM) layer contains numerous tunnels that promote intraplanar percolation. Additionally, the proper amounts of inactive ions with large radius between the layers not only effectively expand the interlayer spacing without consuming many Li+ intercalating sites, but also firmly suppress the expansion of the active TM layers, even endowing the material with zero-strain characteristic22–24. After material structural searching and evaluation, K VO (KVO), a low-cost vanadium-based compound, was selected to verify our design strategy. In KVO, the edge-shared VO tetrahedrons and VO pentahedrons consist of the active TM layers with multiple open tunnels for Li+ percolation and dual redox couples (V5+/V3+), while the KO polyhedron serve as pins to rivet these active TM layers, which not only expand interlayer spacing, but also attribute to this wide 3D network with robust mechanical stability. Meanwhile, KVO with the low cost of potassium (K) and vanadium (V) elements renders it a cost-effective option for large-scale applications. The as-prepared KVO compound displays a high capacity of 377 mAh g−1, enhanced rate performance (approximately 146 mAh g−1 at 10C) and cyclability retaining 88.3% capacity after 19,000 cycles. In-situ X-ray diffraction (XRD), in-situ/ex-situ transmission electron microscopy (TEM), synchrotron radiation based X-ray absorption spectroscopy (XAS), neutron diffraction (ND), time of flight secondary ion mass spectrometry (TOF-SIMS) and density functional theory (DFT) calculations unveil that the prolonged cycling stability is attributed to the zero-strain structural evolution of KVO, and the high rate performance can be explained by the 3D conduction networks for Li+ and robust LiF-enriched solid electrolyte interphase (SEI) of KVO. This strategy to develop negative electrode materials is also further verified by other materials and sheds light on the development of fast-charging, long-life negative electrode materials for future LIBs.

Results
Crystal structure and characteristics
The KVO material is prepared by the simple ball milling process, and the corresponding XRD and ND patterns of the as-prepared KVO are depicted in Fig. 1e, f and Supplementary Tables 1 and 2. Rietveld refinement shows that all Bragg diffraction peaks can be well indexed to an orthorhombic system with a space group of P31m (JCPDS card No. 14-555). The field emission scanning electron microscopy (FESEM) images in the inset of Fig. 1f reveal that KVO presents as micron-sized particles ranging from 1 to 20 μm. HRTEM image in Fig. 1g confirms well-defined (111) planes with an interspacing of 0.328 nm, and TEM with energy-dispersive X-ray spectroscopy (EDS) (Supplementary Fig. 1) shows the uniform distribution of K, V, and O in the KVO sample. The crystal structure of KVO is further depicted in Fig. 1h, i. The VO tetrahedrons and VO square pyramids collectively form (VO )3− TM slabs, and the top view shows these TM slabs contain abundant pentagonal tunnels, which would be favorable for Li+ percolation along the slabs. Meanwhile, all these TM layers are stacked along the c axis direction, and the K+ with neighboring O to form KO polyhedrons, serve as bridges to rivet the adjacent TM slabs with A-B-A stacking. Thus, we successfully synthesized the KVO layered material, and the TM slabs with intraplanar percolating ability and KO with robust interplanar bridge in such layered material generate abundant open tunnels for 3D Li+ transportation and vast sites for Li+ intercalation and volume accommodation. Furthermore, KVO exhibits a remarkably small band gap of 2.17 eV by UV–vis absorption test (Supplementary Figs. 2 and 3), demonstrating much higher electronic conductivity of 6.34×10−8 S cm−1 than Li TiO materials25.

Electrochemical advantages of the as-prepared KVO
To explore the electrochemical advantages of the as-prepared KVO material, it was fabricated as the negative electrode to be charged/discharged in the half cell and full cell as shown in Fig. 2. Supplementary Fig. 4 displays the initial four discharge/charge curves of KVO in half cell at 0.1C (1C = 450 mA g−1), ranging from 0.2 to 3.0 V at 30°C. In the first discharge cycles, the capacity of KVO exhibited more than 700 mAh g−126–28. The origin for the extra capacity can be attributed to 1) the formation of the SEI, resulting from electrolyte decomposition above 0.4 V vs. Li+/Li, 2) the obvious Li+-storage by conductive carbon (Super P) at low potentials, and 3) intercalation-type negative electrode materials that the inserted Li+ ions during the first lithiation cannot be totally extracted. (Supplementary Figs. 5 and 7) When we use the active mass up to 90 %, the initial coulomb efficiency and charge capacity becomes more reasonable. (Fig. 2a) After initial lithiation, the KVO negative electrode presents similar discharge/charge curves in subsequent cycling, suggesting a highly reversible delithiation pr/delithiation process. Figure 2b shows that the KVO sample displays a reversible capacity of 377 mAh g−1 (84% of its theoretical capacity, which is based on the results of theoretical calculations (Supplementary Figs. 8 and 9 and Supplementary Data 1)) at 0.1C, which is 2.5 times higher than that of commercial Li TiO , and higher than that of commercial graphite. Meanwhile, it can also be clearly seen that KVO operates at a quite suitable working potential of ~1.45 V vs. Li+/Li, which is notably higher than the threshold for lithium dendrite formation (Supplementary Figs. 10 and 11 and Supplementary Table 3) and lower than 1.55 V vs. Li+/Li of the renowned Li TiO 29,30. The low-voltage operation of KVO primarily originates from diminished crystal field stabilization in tetrahedral VO units, augmented by isotropic 3D ion diffusion and zero-strain lattice accommodation31–34. Supplementary Fig. 12 demonstrates high rate performance, delivering reversible capacities of 298, 263, 233, 192, and 147 mAh g−1 at 0.5, 1, 2, 5, and 10C, respectively. At 10C, the capacity of the KVO material maintains almost half of that at 0.5C. Under the identical 10C condition, the graphite negative electrode exhibits a lower capacity retention rate of only 6% relative to its own 0.5C capacity. (Supplementary Figs. 13 and 14).

The cycling performance is further characterized at different C rates. Notably, the KVO negative electrode demonstrates a high-capacity retention of 90.2% over 500 cycles at 1C in Fig. 2c. At 5C, it remains a capacity retention of nearly 100.0 % without capacity attenuation over 2000 cycles in Fig. 2d. At a high rate of 10C, it maintains 95.0 % capacity retention over 10,000 cycles. Even after 19,000 cycles, the KVO negative electrode still exhibits capacity retention of 88.3% in Fig. 2e. The cycling stability of KVO in this study is at a relatively high level in the field of electrochemical energy storage materials. (Fig. 2f and Supplementary Table 4). To further validate the advantages of KVO as a negative electrode, we narrowed the voltage range to 0.1 V − 2 V. The average operating voltage within this range was ~0.78 V vs. Li+/Li, the material demonstrates promising electrochemical performance. (Supplementary Fig. 15)

To further investigate the electrochemical merits of the KVO negative electrode material, it was cycled at both high temperature (60°C) and low temperature (−10°C) in Fig. 2g, h. The capacity and capacity retention of KVO at different C rates shows slight enhancement at 60°C. It shows a high reversible capacity of 396 mAh g−1 at 0.1C (Supplementary Fig. 16), slightly higher than that at 30°C. And it also maintains capacities of 348, 317, 282, 224, and 175 mAh g−1 at 0.5C, 1C, 2C, 5C, and 10C, respectively (inset of Fig. 2g). More importantly, the KVO sample yields a capacity retention of 100.2% after 8000 cycles at 10C in Fig. 2g. In stark contrast, Li TiO exhibits a retention of only 20.9 % after 500 cycles at 5C, and commercial graphite particles exhibit a retention of only 36.7% after 1000 cycles at 10C16,25. Meanwhile, when cycled at −10°C, KVO demonstrates a capacity of 184 mAh g−1 at 0.1C (Supplementary Fig. 17) and a rate performance of 84 mAh g−1 at 2C (inset of Fig. 2h). Additionally, it shows capacity retention of 97.9 % at 2C over 2000 cycles (Fig. 2h). To assess the commercialization potential of KVO, KVO||LiFePO full cell was assembled. It is worth noting that the KVO||LiFePO full cell exhibits a large reversible capacity (326 mAh g−1) at a current density of 0.1C and average operating voltage of 2.65 V in inset of Fig. 2i. And its specific energy (on an active material mass basis) is 3.5 times that of Li TiO ||LiFePO . Figure 2i shows that KVO maintains a capacity retention of 96.1% after 1000 cycles at 2C. Meanwhile, the KVO||LiFePO interdigital battery in Supplementary Fig. 18 also exhibits high capacity and prolonged cycling stability. The negligible voltage polarization changes observed in the long-term cycling curves of half cell and full cell at different rates (Supplementary Fig. 19) indicate that KVO has high structural integrity. It should be mentioned that the slight capacity increases of the Li||KVO half cells during the electrochemical reaction can be explained by the continuous KVO activation. (Supplementary Fig. 20 and Supplementary Table 5)35,36 Therefore, the electrochemical characterizations demonstrate that the KVO material shows a great potential to stand out as a practical negative electrode material with high capacity, long lifespan, fast-charging capability and all-climate operation for wide-range applications such as electric vehicles, large-scale energy storage and so on.

Reversible bulk chemical evolution and stable surface chemistry
To better understand the electrochemical performance of KVO, we probe the bulk chemical evolution in KVO during cycling by hard XAS. The samples for detailed tests were selected according to the Cyclic Voltammetry (CV) curves in Supplementary Fig. 21. There are two characteristic regions in the spectra: the pre-edge peaks in the range 5465–5470 eV, and main peaks at about 5484 eV. The oxidation state of the V is known to be sensitive to the intensity ratio between the pre-edge and main peaks (referred to as P/M hereafter)1,37,38. It can be seen that the valence state of V is +5 at pristine state in Fig. 3a, and then the P/M value decreases from 0.88 to 0.77 when KVO is discharged to 0.6 V, and the P/M value further decreases to 0.46 when discharged to 0.2 V, which is significantly lower than the reference value of 0.70 for V4+. This observation indicates the a two-electron transfer process from V5+ to V3+ in the KVO sample. Conversely, during charging, the P/M value gradually increases and nearly approaches to the original state at the end of charge (Fig. 3b). To investigate the chemical evolution of the near surface, soft XAS were collected at total fluorescence yield (TFY) mode with a detection depth of 50~100 nm. Figure 3c shows the V L-edge spectra of the pristine KVO sample exhibit two main peaks at 517.5 eV for V-2p 3/2 and 524.9 eV for V-2p 1/2, respectively, which can be used as V5+ reference. When the electrode is gradually discharged to 0.2 V, the two main peaks gradually drift to lower energy levels, and interestingly, the intensities of their shoulder peaks gradually increase. The comparison between the sample at 0.2 V and standard reference samples indicates the coexistence of V4+ and V3+ at the end of discharge. Conversely, during charging, the V L-edge spectra progressively recover and return to their pristine state upon charging to 3 V. The ex-situ X-ray Photoelectron Spectroscopy (XPS) characterization also confirms this surface evolution in Supplementary Fig. 22. The valence state changes of V in the surface (50~100 nm) and the bulk are further illustrated clearly in Fig. 3d, e. It is worth noting that V can be almost recovered to +5 during the charge process in both bulk and surface, indicating the highly reversible V redox, the slight discrepancy from the original state is attributed to incomplete extraction of trace Li⁺ ions during the initial cycle39,40. At the same time, V undergoes a two-electron transfer from V5+ to V3+ in both bulk and surface, while there is slight charge heterogeneity from bulk to surface as well. Ex-situ Raman spectroscopy (Supplementary Fig. 23) demonstrates the highly reversible structural behavior of KVO, with characteristic V–O bending vibrations at 930 and 962 cm−1 showing no peak shifts or new phases, consistent with its zero-strain characteristic41–45. However, a slight attenuation in restored Raman intensity after delithiation suggests minor irreversibility. Electron Energy Loss Spectroscopy (EELS) analysis (Supplementary Fig. 24) reveals slight surface oxygen defects, where Li⁺ ions become trapped during initial lithiation through stabilization by defect sites and local reduction of V5+ to V4+ (Supplementary Table 6, Supplementary Fig. 25 and Supplementary Data 2)41,46,47. The synergistic effect of shallow Li trapping and zero-strain lattice confinement maintains bulk V5⁺ stability while accommodating surface reduction, preventing bulk degradation and enabling favorable structural integrity and cycling stability48–51.

The SEI is crucially important for the cycling performance of negative electrode materials, and to elucidate the chemical evolution of SEI in the KVO material, the O and F K-edge soft XAS spectra at both total electron yield (TEY) and TFY modes were collected at different charge and discharge states52. Figure 3f, g show the TEY and TFY data for O K-edge XAS spectra, and at pristine state, only the lattice oxygen signal can be seen in both modes. When the electrode is discharged to 1.1 V, a slight oxygen signal of SEI emerges at 534 eV at both TEY and TFY modes, and these signals are resembled those of organic carbonates such as LEMC (lithium ethylene mono-carbonate), LEDC (lithium ethylene di-carbonate) (Supplementary Figs. 26 and 27). The peak intensities for SEI at both TEY and TFY modes become more obvious when the electrode is further discharged to 0.2 V, indicating a significant oxygen compound formation in SEI. The intensity for SEI remains stable even after 100 cycles, further suggesting the stable oxygen compounds form in SEI after the initial cycle. Meanwhile, the peak intensity for oxygen compounds of SEI at TFY mode is much lower than that at TEY mode, possibly due to the enrichment of combinates in the surface layer of the electrode. Additionally, F K-edge soft XAS spectra at TEY and TFY modes are depicted in Fig. 3h, i and Supplementary Fig. 28, respectively. There is only a signal for PVDF binder at pristine state at both modes, and after discharging to 1.1 V, LiF starts to emerge. The LiF becomes a dominant compound at 0.2 V discharge and remains quite stable even after 100 cycles. There are no obvious spectral differences between the TFY and TEY data, inferring the LiF-rich SEI formation in the KVO sample.

To further investigate the structure and chemical composition of the SEI, TOF-SIMS analysis was conducted on the KVO sample after 100 cycles. Supplementary Fig. 29 displays the depth analysis curve of the KVO electrode, revealing characteristic ion fragments of LiF- and C HO-, which confirm the presence of LiF and organic carbonates in SEI. As the sputtering time increases, the LiF content gradually increases while the organic composition sharply decreases, as depicted in Fig. 3j. Hence, combined with the soft XAS data, it can be found that SEI on the KVO electrode exhibits a clear bilayer structure, comprising an external organic layer and an internal inorganic layer (markedly enriched in LiF). This result is well consistent with the TEM results in Supplementary Fig. 30. Therefore, the stable LiF-enriched SEI formed in the KVO material is believed to play a crucial role in enhancing the Li+ diffusivity and long cycling performance. To determine whether LiF enrichment was specific to KVO or inherent to the vanadate framework, we performed comparative TOF-SIMS studies on isostructural K MgV O (Supplementary Fig. 31). Crucially, this K MgV O exhibited the same pronounced LiF dominance in its SEI. DFT calculations uncovered a surface-selective reaction pathway driving LiPF decomposition: Fluorine exhibits significantly higher electronegativity than oxygen, enabling stronger ionic-covalent hybrid bonds with electropositive vanadium sites. From a thermodynamic perspective, this is conducive to the selective adsorption of LiPF on the electrode surface. Compared with common electrolyte decomposition products such as EC (−0.59 eV) or DEC (−0.056 eV), the calculated adsorption energy (−0.708 eV) is significantly higher. When LiPF approaches the KVO surface, PF ⁻ anions undergo site-specific chemisorption (Fig. 3k, Supplementary Fig. 32, Supplementary Table 7 and Supplementary Data 3). The LiF-rich SEI originates from the inherent fluorophilicity of exposed vanadium centers in layered vanadates, where strong V–F bonding drives preferential PF ⁻ decomposition and localized LiF precipitation. (Supplementary Figs. 33 and 34 and Supplementary Data 4) This mechanism, which is generalizable across vanadate frameworks, represents an intrinsic interfacial property that complements the structural advantages of KVO.

Zero-Strain structural evolution at wide-temperature cycling
The crystal-structure evolution of KVO during charge and discharge was explored through in-situ XRD characterization15. After the initial activation in Supplementary Fig. 35, the in-situ XRD patterns of the subsequent two cycles at 0.2 °C and 30 °C are illustrated in Fig. 4a. It can be seen that all patterns nearly show no change upon cycling. The enlarged view of patterns in Supplementary Fig. 36 shows that the (001) and (111) peaks slightly shift to higher 2θ angles during the first charge, and then drift back to their original positions during the second charge, and reversibly evolve after cycling as shown in Fig. 4a. The (001) and (111) peaks change in P31m structure indicate three possible Li+ occupation sites, 1a, 3c(1), and 3c(2) in KVO (Fig. 5d and Supplementary Fig. 51), which will be discussed later. Meanwhile, the intensities of all diffraction peaks slightly decrease and then almost fully recover during cycling. Such highly reversible single-phase evolution serves as strong evidence to rather favorable structural reversibility and stability of KVO during the repeated charging and discharging process. Furthermore, all the precise lattice parameters of KVO during cycling are calculated by Rietveld refinement as shown in Fig. 4b and Supplementary Fig. 37. All the lattice parameters a, c and V exhibit high reversibility and zero-strain characteristic with maximal variations of approximately 0.11%.

To further explore the structural merits of the as-prepared materials, an additional in-situ XRD test was conducted at 60°C in Fig. 4c and Supplementary Fig. 38. The directions of peak shifts remain consistent with those observed at 30°C, while the peak drift degree is slightly larger than that at 30°C. Rietveld refinement of the in-situ XRD patterns at 60°C shows the maximum change in V-value slightly increases to 0.36% in Fig. 4d and Supplementary Fig. 39. The increment is evidently attributed to the larger reversible Li+-storage capacity at 60°C (Supplementary Fig. 16). Notably, the zero-strain characteristic persists at 60°C for the KVO material. This desirable trait of zero-strain at various temperatures undoubtedly contributes to its prolonged cycling stability across a wide-temperature range, particularly at 60°C. This is one of the lower percentages observed in Li⁺ storage materials (Fig. 4e and Supplementary Table 8).

In-situ TEM was employed to further investigate the microstructural and morphological evolution of KVO during lithiation53,54. As highlighted in Fig. 4f, Supplementary Fig. 40 and Supplementary Movie 1, significant movement of strain fringes is observed, originating from Li+ intercalation in the KVO primary particles, while the microstructure, morphology, and volume of the as-observed KVO particle exhibit negligible change. Ex-situ HRTEM characterization in Supplementary Fig. 41 aligns well with the findings of in-situ TEM, demonstrating unchanged interplane spacing of (201) planes from the pristine to the lithiated and delithiated states. In-situ optical microscopy (Supplementary Fig. 42, Supplementary Movie 2, and Supplementary Movie 3) reveals KVO particles maintain structural integrity with minimal volume change after 100 cycles—in sharp contrast to the severe expansion of graphite—while in-situ XRD (Supplementary Fig. 43) confirms that the zero-strain character of KVO persists even at 0.01 V, demonstrating macroscopic stability. Furthermore, the stress simulation diagrams in KVO multi-particle simulation by COMSOL in Fig. 4g, Supplementary Fig. 44 and Supplementary Movie 4 show the minimal stress and strain reversibility during the initial cycle, and the XRD, FESEM, and TEM characterization in Supplementary Fig. 45 display favorable stability of the KVO electrode after 500 cycles at 5C. Based on our analysis, we believe that the slight changes in the XRD after the cycle are mainly due to the difference inherent to measuring electrode sheets versus pure powders, as well as the preferred orientation generated during the rolling process, rather than indicating the presence of phase transformation or structural degradation. (Supplementary Figs. 46 and 48 and Supplementary Table 9) Thus, it is believed that the zero-strain single-phase change at a wide-temperature range, moderate stress, and strain reversibility during cycling enable KVO to sustain over tens of thousands of cycles.

Local structure for zero-strain evolution and electrochemical kinetics
To investigate the origin of the zero-strain characteristic of KVO, we carefully analyze the changes of local structure and chemical bond length through EXAFS and XRD refinement at the lithiated and delithiated states55–57. Fig. 5a shows the Fourier transform of V K-edge EXAFS with a selected k-space range of 0~14 Å−¹. The R-space curves of V K-edge EXAFS reveal two main peaks corresponding to the first shell for the V–O coordination and the second shell for the V–V coordination, and the horizontal position of the shell indicates the uncorrected bond length, while the intensity the changes in the coordination number (CN) and the Debye-Waller (DW) factor58. (Supplementary Table 10)

From the pristine state to 0.2 V discharge, the first shell for the V–O coordination slightly shifts rightward due to the reduction of V5+ ions to larger V4+/V3+ ions, and then the position of the second shell for interplanar V–V coordination also simultaneously shifts rightward upon 0.2 V discharge. This phenomenon is well consistent with the XRD refinement result, as shown in Fig. 5b and Supplementary Tables 11 and 13, and the lithium intercalation increases the volume of the VO tetrahedron by 1.05% and the volume of the VO pentahedron by 0.19 %. More interestingly, the KO polyhedron contracts its volume by 0.42%, which is possibly attributed to the pillar effect of K buffering the volume expansion of V in redox reactions. During charging, these two shells almost horizontally drift back to their original positions. We performed quantitative EXAFS fitting for both pristine, lithiated and delithiated states of KVO. A synergetic variation of parameters is observed: upon discharge, the CN of V–O and V–V decrease slightly from 4.6 to 4.5 and from 2.5 to 2.3, respectively, accompanied by an increase in the DW factor. All parameters recover to values close to the pristine state after charging. This phenomenon depicts a dynamic but perfectly resilient coordination environment undergoing elastic distortion, which is fully consistent with the long-range zero-strain behavior confirmed by our in-situ XRD. (Supplementary Table 10 and Supplementary Figs. 49 and 50) Additionally, the Wavelet-Transformed EXAFS (WT-EXAFS) in Fig. 5c further clearly illustrates that the maximum WT values of the V–O and V–V coordination do not significantly deviate during charging and discharging. Thus, the inactive KO slabs and active TM slabs in the as-proposed KVO material can mutually compensate each other and effectively buffer the volume change in lattice, demonstrating zero-strain structural change at a wide-temperature range59.

To gain the origin of the fast charge capability of KVO, the Li+-transport paths in KVO were further simulated by using DFT (Supplementary Data 5). In-situ XRD in Fig. 4 indicates three possible Li+ occupation sites, 1a, 3c(1), and 3c(2) in the P31m structured KVO as shown in Fig. 5d and Supplementary Fig. 51. The formation energy calculation shows Li+ favorably occupy 1a sites and then the 3c(1) sites (Supplementary Fig. 52). Interestingly, in the optimized crystal model, since the V–O bonds are in a negatively charged state, they are more likely to adsorb Li+, and the Li+ are more likely to shift towards the V–O position. For KVO, the maximum inserted Li+ number is as high as 10 (Supplementary Fig. 9). Within the range of 1≤x≤10, the binding energies of xLi and KVO are both negative, indicating that KVO is suitable for utilization in LIBs. Figure 5e–g illustrate three potential paths for Li+ transmission: 1a→3c(1)→3c(2) (path I), 1a→3c(1)→3c(2) (path II), and 1a→3c(1)→1a (path III). Path I illustrates the linear migration of Li+ along the pathway 1a→3c(1)→3c(2). The results obtained from the Climb Image-Nudged Elastic Band (CI-NEB) method demonstrate that Path I encounters a relatively high energy barrier of 1.85 eV. The transition state, situated near the 3c(1) sites in Fig. 5e, corresponds to the high energy barrier due to the competitive affinity between V–O and Li–O. Through the transition state, the energy of the V–O bond is reduced to form the Li–O bond. Path II in Fig. 5f represents a special pathway through Li+ percolation perpendicular to the TM slabs. Li+ at 1a sites initially migrates to a location adjacent to 3c(1) sites before shifting to 3c(2) sites, featuring a relatively low energy barrier of 0.98 eV. Path III in Fig. 5g shows a bent transport pathway in the inactive KO slab, which involves Li+ migration between sites 1a and 3c(1) with a moderate energy barrier of 1.01 eV. It is evident that the barriers of Path II and Path III are notably smaller than that of Path I, and these favorable Li+ diffusive paths can establish an interconnected 3D large channel for Li+ transport in KVO, which surely facilitates a rapid Li+ diffusion. Inductively coupled plasma (ICP) analysis confirming negligible K⁺ loss (Supplementary Table 14) combined with DFT calculations revealing a prohibitively high K+ deintercalation energy (+5.28 eV) versus favorable Li+ removal (−1.58 eV) (Supplementary Fig. 53, Supplementary Table 15 and Supplementary Data 6) demonstrate that Li⁺/K⁺ ion exchange is thermodynamically prohibited in KVO. The Galvanostatic intermittent titration technique (GITT) experiments in Supplementary Figs. 54 and 55 and Supplementary Table 16 also indicate a high Li+ diffusion coefficient (D ) of KVO in comparison with other materials. Additionally, COMSOL simulations in Fig. 5h, Supplementary Fig. 56 and Supplementary Movie 5 visualize the Li+ concentration distribution diagram in KVO multi-particles, and the 3D channels enable that the Li+ concentration distribution is relatively uniform with small gradient and almost non-disparity from bulk to surface in particles during discharge and charge. Therefore, the theoretical calculations combined with the experiment results elucidate the 3D pathway for Li+ diffusion, which is attributed to the rapid Li+ diffusion and fast-charging capability of KVO.

Universality of the strategy to discover advanced zero-strain battery materials
The proposed strategy to search a high capacity, high-rate and long-cycling zero-strain negative electrode material is successfully demonstrated by the KVO sample. The KVO sample displays a layered structure with intraplanar percolation and interplanar bridge. To explore the universality of this proposed strategy, we further search and investigate some other compounds with similar layered structures in Fig. 6. Figure 6a shows the XRD patterns of the K VP O material prepared via a simple ball milling method. Rietveld refinement reveals that all Bragg diffraction peaks can be successfully indexed to a tetragonal system with the space group P-421m (JCPDS card No. 40-283). (Supplementary Table 17) The schematic diagram of the crystal structure for K VP O is illustrated in Fig. 6b. Similar to KVO, the VO pentahedrons and PO tetrahedrons interconnect to form tunnels in the TM slab, and KO serves as an inactive layer to pin-bridge the adjacent TM slabs. As anticipated, K VP O exhibits reversible capacity (306 mAh g−1 at 0.1C), rate capability (133 mAh g−1 at 10C) and cyclability, with a capacity retention of 93.8% at 10C over 1200 cycles (Fig. 6c and Supplementary Fig. 57). More importantly, Fig. 6d illustrates in-situ XRD patterns at the initial two cycles at 0.3C at 0.01–3.0 V. It is evident that nearly all peaks display negligible shift during the repeated Li+ insertion/extraction process. Additionally, Fig. 6e shows the XRD spectra of another compound, K MgV O with a crystalline space group of P42/mnm. (Supplementary Table 18) The schematic diagram in Fig. 6f indicates that the MgO tetrahedrons and VO tetrahedrons interconnect to form tunnels in the TM slab, and KO serves as an inactive layer to pin-bridge the adjacent TM slabs. Similarly, K MgV O delivers a capacity of 339 mAh g−1 at 0.1C, rate capability of 126 mAh g−1 at 10C, and a capacity retention of 95.2% after 2000 cycles at 10C. (Fig. 6g and Supplementary Fig. 58). As anticipated, the K MgV O sample also exhibits zero-strain structural evolution during cycling as shown in Fig. 6h. Similarly, we have synthesized phase-pure Rb VO via solid-state reaction, confirmed by XRD, SEM, and TEM without impurities (Fig. 6i, Supplementary Fig. 59a–c and Supplementary Table 19), and demonstrated electrochemical performance (Supplementary Fig. 59d–f): rate capability of 107 mAh g−1 at 5C, long-term cycling stability with 87.1% capacity retention after 1000 cycles at 5C, and near zero-strain behavior (Fig. 6j). Thus, our study demonstrates a general principle for searching and designing materials with zero-strain, high capacity, long-cycling and fast charging capabilities. This structural innovation differs significantly from conventional doping. (Supplementary Table 20) These layered structure materials are essential to maintain active TM layers with abundant open tunnels that facilitate intraplanar percolation, and also rivet appropriate amount of inactive ions between TM slabs to form inactive layers with vast interlayer space for Li+ storage and diffusion, which are also mainly used to buffer the expansion and contraction of TM slabs during cycling. Meanwhile, the successful realization of gradual decreasing in operating voltage strongly validates the effectiveness and tunability of our structural design. (Supplementary Fig. 60)

Overall, this study proposes a structural design strategy for negative electrode materials by integrating intraplanar percolation-interplanar bridging in the layered matrix, which combines the advantages of mainstream battery material structures to develop negative electrode materials for batteries with high capacity/rate capability, long cycling, and zero-strain behavior over a wide-temperature range. Taking the low-cost vanadium-based compound KVO as a representative, the abundant pentagonal tunnels in its active TM layers facilitate rapid intraplanar Li+ percolation, while the KO polyhedrons formed by large-radius K+ between layers act as bridging units, which cannot only widen the interlayer spacing to retain sufficient Li+ intercalation sites, but also effectively constrain the expansion of TM layers, endowing the material with enhanced structural stability. At 30//60//−10°C, the as-prepared KVO compound displays high capacity of 377//385//184 mAh g−1 at 0.1C, and exhibits rate performance and long-term cyclability with capacity retention of 88.3% at 10C rate after 19000 cycles at 30°C. The in-situ XRD, ex-situ XAS, and other characterizations confirm that V in KVO undergoes a two-electron transfer process (V5+ ↔ V3+) during Li+ (de)lithiation. DFT calculations combined with GITT reveal that the 3D Li+ diffusion pathways in KVO, ensure fast charging/discharging capability. Most importantly, the complementary volume changes between KO polyhedrons and the VO layers in KVO result in a maximum lattice parameter variation of only 0.11% during cycling, realizing zero-strain characteristics. The universality of this structural design strategy has been also verified by other materials, such as K VPO , K MgV O, and Rb VO , all of which exhibit high-capacity/rate capability, long cycling, and zero-strain behavior. In summary, this study has successfully demonstrated a general structural design strategy of high-performance electrode materials for next-generation batteries, while plenty of work is still needed for the commercial application of the as-designed materials in the future.

Methods
Material synthesis
The synthesis of KVO sample was conducted via a straightforward and cost-effective solid-state reaction method. Specifically, NH VO (99%, Macklin) and K CO (99%, Macklin) powders were thoroughly mixed in a molar ratio of 10:3, with the weighed masses of 5.849 g for NH VO and 2.073 g for K CO . Subsequently, the mixture underwent energetic milling using a high-energy ball-milling machine (SPEX 8000M, SPEX SamplePrep, USA) equipped with an 8007 stainless steel vial (volume: 65 mL). The milling media consisted of two 1/2 in. (12.7 mm) and four 1/4 in. (6.35 mm) stainless steel balls. The milling program was set as follows: continuous operation for 60 min, followed by an intermittent rest time of 30 min, and this cycle was repeated 10 times (total milling duration: 15 h) at a rotational speed of 1725 r min−1. After ball-milling, the resultant powder was collected and subjected to calcination at 400°C for 8 h in air using a muffle furnace (KSL-1200X-J, Hefei Kejing Material Technology Co., Ltd) with a ramp rate of 5°C min−1. Similarly, K VP O samples were fabricated through a solid-state reaction route. Initially, NH VO (99 %, Macklin), K CO (99 %, Macklin), and NH H PO (99.9 %, Macklin) powders were uniformly blended at a molar ratio of 1:1:2, with the weighed amounts of 1.170 g, 1.382 g, and 2.301 g respectively. Subsequently, the mixture underwent high-energy ball milling using the same SPEX 8000M equipment equipped with an 8007 stainless steel vial, adopting the same milling protocol: 60 min of continuous grinding, 30 min of intermittent rest, and 10 cycles of repetition at 1725 r min−1. Following milling, the obtained powder was gathered and calcined in air at 550°C for 8 h via a KSL-1200X-J muffle furnace with a heating rate of 5°C min−1. Micrometer-sized K MgV O particles were prepared via a solid-state reaction protocol. NH VO (99%, Macklin), K CO (99%, Macklin), and MgO (99.9 %, Macklin) were first weighed out as the starting materials (2.340 g, 1.382 g, and 0.403 g respectively) and then mixed homogeneously at a molar ratio of 2:1:1. The mixture was subsequently processed by high-energy ball milling using a SPEX 8000M apparatus paired with an 8007 stainless steel via. Two 1/2 in. (12.7 mm) and four 1/4 in. (6.35 mm) stainless steel balls served as the milling media, and the milling was carried out following the same procedure as above: 60 min of continuous operation alternating with 30 min of intermittent rest, repeated 10 times (total milling duration: 15 h) at 1725 r min−1. After milling, the powder was collected and calcined in air at 850°C for 8 h in a KSL-1200X-J muffle furnace with a heating rate of 5°C min−1. Rb VO samples were fabricated via a solid-state reaction method. Firstly, Rb CO (99%, Macklin) and NH VO (99%, Macklin) were weighed as 3.464 g and 5.849 g, respectively, and mixed homogeneously at a molar ratio of 3:10. The mixture was then processed by high-energy ball milling following the same procedure as above. After milling, the powder was collected and calcined in air at 500°C for 6 h using a KSL-1200X-J muffle furnace with a heating rate of 5°C min−1. Both commercial graphite (SFC-E) and commercial LiFePO (P198) were purchased at Shenzhen BTR New Energy Materials Co., LTD. Li TiO (MA-EN-AN-0020) was purchased at Guangdong Canrd New Energy Technology Co., LTD.

Material characterizations
The KVO crystal structure was characterized via XRD measurements conducted on a Bruker D8 Discover diffractometer equipped with Cu-Kα radiation. The instrument operated at 40 kV and 40 mA in reflection geometry. The XRD patterns underwent Rietveld refinement with the General Structure Analysis System (GSAS) software. Neutron powder diffraction (NPD) data were collected using the Echidna instrument at the Australian Nuclear Science and Technology Organisation (ANSTO). Structural refinements were executed via the Rietveld method, implemented in the GSAS (General Structure Analysis System) software suite. Morphological characteristics were examined using FESEM on an FEI Tecnai G2 F20 STWIN instrument. Additionally, TEM combined with EDS was performed on a Tecnai G2 F20 system. Additionally, ex-situ XPS measurements were performed with a PHI 5000 VersaProbe II system, using monochromatic Al-Kα irradiation at 1486.6 eV.

Electrochemical tests
The four electrodes were fabricated by mixing 80 wt.% active material (KVO, K VP O , K MgV O or RbVO ), 10 wt.% Super-P conductive carbon (99.5%, Canrd Technology Co., Ltd.), and 10 wt.% polyvinylidenefluoride (PVDF, MA-EN-BI-0009, Arkema, molecular weight: ~1,000,000) in N-Methyl-2-pyrrolidinone (NMP, 99.5%, Sigma-Aldrich). The resulting slurry was evenly spread onto copper foils (≥99.8%, 9 μm) using a coating instrument (MSK-AFA-ES200, Shenzhen Kejing Star Technology Co., Ltd.). Active materials (Li TiO , graphite or LiFePO ), Super-P conductive carbon and PVDF in a mass ratio of 8:1:1 were mixed in NMP to spread on current collectors using the same coating instrument—copper foils (≥99.8%, 9 μm) for Li TiO or graphite, and aluminum foils (≥99.6%, 16 μm) for LiFePO . After vacuum drying at 60°C for 12 hours, the electrodes were further heated to 120°C for 2 h, and then the resulting electrodes were trimmed into 12 mm diameter circular sheets using an MSK-T10 instrument (Shenzhen Kejing Star Technology Co., Ltd.).

In the argon-filled glove box, half cells were assembled using KVO, K VP O , K MgV O, RbVO , Li TiO or graphite as the working electrode and Li metal (diameter: ~16.0 mm, thickness: ~0.5 mm) as the counter and reference electrode. The average loading of negative electrodes was ~5.0 mg/cm2. To fabricate pre-lithiated KVO negative electrodes used in KVO||LiFePO full cells, Li||KVO half cells were subjected to one cycle at 0.1C within a voltage range of 0.2–3 V prior to disassembly in an argon-atmosphere glove box. The positive electrodes in full cells utilized commercial LiFePO with an N/P ratio of ~1.05 (the average loading of LiFePO is ~12.5 mg/cm2). The current rate of 1C for both half and full cells was equal to 450 mA g⁻1, corresponding to the theoretical capacity of KVO. Their specific capacities were calculated based on the mass of KVO on the working electrodes. Similarly, Li TiO ||LiFePO full cells were assembled using commercial LiFePO for the positive electrodes and commercial Li TiO for the negative electrodes (an N/P ratio of ~1.05, 1C = 175 mA g−1 for Li TiO ). All half and full cells was fabricated in the 2032-type coin cell. The electrolyte consisted of 1 M LiPF dissolved in an equal-volume mixture of ethylene carbonate, diethyl carbonate, and dimethyl carbonate (LB315, Zhangjiagang Guotai Huarong New Chemical Materials Co., Ltd.). The electrolyte for all coin cells was controlled at 80 μL. As separators, microporous polypropylene films (Celgard® 2325, thickness: 25 μm, Gurley value: 620 s, porosity: 39%, one piece per cell, Canrd Technology Co., Ltd.) were utilized. Galvanostatic charge/discharge tests and GITT studies were conducted using a multichannel battery tester from Wuhan Land Electronic Co., Ltd., China. Electrochemical impedance spectroscopy (EIS) tests and CV tests were performed using a CHI660E electrochemical workstation. The EIS measurements were conducted under a potentiostatic mode with an alternating current (AC) signal amplitude of 5 mV. The frequency range was set from 105 Hz to 10⁻1 Hz, with 73 total data points for sufficient spectral resolution. after target cycles, all cells stood for 30 min to eliminate galvanostatic-induced concentration polarization and bring electrode-electrolyte interface to quasi-stationary state, ensuring reliable EIS data. All ex-situ measurements were performed following the procedure below: all cycled electrodes were promptly collected from the cells at designated voltages to mitigate side reactions between the cycled electrodes and electrolyte. Subsequently, the electrodes were thoroughly rinsed with dimethyl carbonate to eliminate soluble surface species. After drying, all electrodes were transferred into the experimental vacuum chamber via a custom-designed sample transfer module in an argon-filled glove box, ensuring no exposure to ambient air. All electrochemical tests were conducted in climatic chambers at different temperatures, including: 30°C and 60°C climatic chambers (SPX-150BIII, temperature fluctuation: ±0.5°C, Tianjin Taiste Instrument Co., Ltd.) and −10°C low-temperature climatic chamber (ECT-150-70CPSD, temperature fluctuation: ±0.2°C, Giant Force Instrument Enterprise Co., Ltd.).

In-situ characterizations
To investigate the crystal-structural evolution, in-situ half-cells were assembled for XRD analysis using a LIB-LHTXRD-LN system (Beijing Scistar Technology Co., Ltd.). Each cell was configured with a fiberglass separator and a beryllium window, enabling in-situ XRD measurements. The Rietveld refinement of the XRD pattern was performed by using the GSAS software with the EXPGUI interface. For real-time observation of morphological and structural evolution, in-situ TEM was employed. This was facilitated by an electrochemical holder developed by the X-mech Center at Zhejiang University. In the in-situ TEM half cell, the two electrodes consisted of KVO particles on a silver rod and lithium metal on a tungsten rod, respectively. A Li O coating, acting as a solid electrolyte, formed on the surface of the lithium metal. Lithiation processes, transitioning from lithium metal to Li O and further to KVO, were induced by applying a negative voltage of –3.0 V when the two electrodes were in contact.

Related characterization of synchrotron radiation based
Soft XAS experiments were performed at the V L-edge, O K-edge, and F K-edge using beamline 02B02 at the Shanghai Synchrotron Radiation Facility (SSRF). Photons ranging from 50 to 2000 eV were provided by the bending magnetic beamline, with a sample beam size of 150 μm × 50 μm. Spectra for V, O, and F were simultaneously collected using surface-sensitive TEY and TFY modes. Beam energy calibration was carried out using the spectra of SiTiO and LiF samples. The bulk chemical evolution was investigated by hard XAS at the SSRF Beamline 1W1B. Additionally, V O, VO and C H O V were utilized as standard samples for V5+, V4+ and V3+, respectively.

Theoretical calculations
To elucidate the Li+-storage mechanisms and determine structural energetics, DFT calculations60 were implemented with the Vienna Ab-initio Simulation Package (VASP). Structural optimization employed the projector augmented wave (PAW) approach61 together with the Perdew–Burke–Ernzerhof (PBE) functional under the generalized gradient approximation62. A simplified grid model (1×1×1) of KVO served as the initial framework. The plane-wave basis set was defined by a cutoff energy of 400 eV, while convergence thresholds were set at 10−6 eV for total energy and 0.005 eV Å−1 for atomic forces.

Peer review rounds
[{"round": 1, "reviewer_comments": {"reviewer_1": "In this study, the authors developed anode materials with three-dimensional transport pathways and zero-strain behavior for fast-charging Li-ion batteries. In particular, they highlighted the structural advantages of K3V5O14 (KVO), which features a unique layered structure with both intraplanar percolation and interplanar bridging. The authors investigated its electrochemical and structural properties through various methods, including rate capability test, long-term cycling, X-ray analyses, microscopic observations, and computational research. Furthermore, based on KVO, they extended their design strategy to similar compounds, K2VP2O8 and K2MgV2O7, to highlight the versatility of their approach. The overall scientific question addressed by this work is timely and relevant, and the data presented, particularly operando measurements, are sufficient to support the key claims. However, the manuscript currently lacks clarity in the presentation of electrochemical and structural analyses. The scientific logic should be more clearly articulated, and key details must be included to avoid confusion and strengthen the conclusions. Accordingly, the manuscript requires revision before it can be considered for publication in Nature Communications.\n1. Please revise the definition of XAS from “synchrotron radiation (XAS)” to “synchrotron radiation based X-ray absorption spectroscopy (XAS)”.\n2. How did the authors define C-rates or theoretical capacity for KVO, 1 C = 450 mA g-1? Is it based on the theoretical calculation? In the first discharge cycles, the KVO exhibited more than 700 mAh g-1 at 30 °C (Fig. S4), which appears to surpass the theoretical capacity. The authors should clarify this by presenting the electrochemical reaction equation and detailing the basis for the theoretical capacity used.\n3. Please provide corresponding voltage profiles for long-term cycling data (Fig. 2c, 2d, 2e, 2g, 2h, and 2i). These are essential for understanding capacity retention and degradation mechanisms.\n4. In Fig. 2g and 2h, the x-axis in voltage profiles should be labeled as specific capacity (mAh g-1) instead of cycle number.\n5. Following Question 2, the formation cycles likely induce substantial structural changes in KVO due to the potential trapping of large amounts of lithium ions to stabilize the structure. Although XPS (Fig. S11) and in-situ XRD (Fig. S16-S19) are provided, they are insufficient to reveal the full extent of these transformations. Please include complementary structural characterizations, such as Raman spectroscopy, nuclear magnetic resonance (NMR), or infrared (IR) spectroscopy. Ex-situ analysis would be suitable for the supporting information.\n6. XPS results (Fig. S11) suggest a shift in V oxidation state from 5+ to a 4+-dominated state after formation cycles. However, XAS results (Fig. 3d, 3e) show only slight changes, indicating a 5+-dominated bulk even after cycling. This discrepancy suggests that oxidation state changes may be surface-confined. Still, if lithium trapping occurs, this should significantly affect the bulk chemistry. The authors should provide a systematic explanation reconciling these observations.\n7. Even in the absence of significant peak shifts during cycling, notable changes in peak intensities are observed after 500 cycles (Fig. S23). The authors should discuss the origin of these changes-whether due to preferred orientation, phase transformation, or structural degradation.\n8. Please include the voltage profiles corresponding to the GITT measurements (Fig. S26).\n9. It is strongly recommended to discuss the formation mechanisms of KVO using computational results as demonstrated in Fig. 5.\n10. Please provide more details of electrochemical data for K2VP2O8 and K2MgV2O7, such as voltage profiles and rate capability.", "reviewer_2": "This manuscript introduces the design of “zero-strain” K3V5O14 (KVO) anode materials with high capacity, rapid Li ion conductivity, and enhanced structural stability. The authors characterize the structure of the anode material in detail and elucidate the relationship between the structural characteristics and the high energy storage performance. Combined with advanced characterization technology and computational simulation, the lithium storage mechanism of the KVO anode was studied from the perspectives of the chemical state, bulk structure, and microscopic scale. The authors demonstrate a general principle for searching and designing materials with high energy storage performance and implement it in K2VP2O8 and K2MgV2O7 anodes. This manuscript is clearly written, and the data are nicely presented and analyzed. I recommend publication after some minor corrections are made, as suggested below.\n1. Why does the discharge specific capacity increase during the charge-discharge cycling in Figure 2?\n2. Please explain the reason why the first circle discharge capacity is significantly higher than the charging capacity in Figures S7 and S8, and analyze the reasons for the inconsistency with the illustrations in Figures 2g and 2h.\n3. Figure 3 demonstrates the formation of a LiF-rich SEI layer, which is beneficial for cycling stability. But what is the formation mechanism of such LiF-Rich layer?\n4. On page 4, line 89, the manuscript says that there are dual redox couples (V5+/V3+), but this appears to be only one redox couple.\n5. The insert in Figure 2i is not clear enough.\n6. Please explain the phenomenon in Figures 3d and 3e where the valence state changes in the bulk and near-surface are not completely consistent.\n7. Please provide an enlarged display of the strain region in Figure 4f.", "reviewer_3": "Ma et al. report the use of K3V5O14 (KVO) as a fast-charging anode material for Li-ion batteries, presenting extensive data on its phase evolution and surface chemistry during cycling. However, in my view, the V2O5-based anode system has already been widely explored, and this study does not deliver sufficient conceptual novelty or electrochemical advancement to meaningfully progress the field. Furthermore, the characterization data contain several errors, lack standardization, and are not able to support the authors’ conclusions. I do not recommend publication in Nature Communications.\nMajor Comments:\n1. Why is the V5+/V3+ redox couple in KVO located within the anode voltage window, despite many reported cathode materials with similar compositions (e.g., K0.25V2O5, KV3O8) operating at higher potential?\n2. Have the authors investigated Rb- or Cs-substituted analogues of KVO? If the proposed mechanism is valid, such substitutions (which also crystallize in the M3V5O14 phase) might further enhance rate performance despite with lower capacity.\n3. Figure S1 appears to show non-uniform oxygen distribution, contradicting the claim of uniformity in the main text. Also, abbreviations like “EDS” and “EDX” should be used consistently throughout the manuscript.\n4. The voltage profile of KVO is notably sloped, likely due to pseudocapacitive contributions or the amorphization process revealed in ex-situ HRTEM images shown in Figure S21.\nA key factor driving the commercialization of LTO is its stable Li-ion storage behavior, which delivers a flat redox plateau and low polarization, features that extend beyond mere coin-cell performance.\n5. Have the authors considered potential Li-/K- ion exchange between KVO and the electrolyte during cycling, given the large chemical potential difference? It is plausible that K-ion re-intercalation to KVO partially contributes to the observed rate performance due to the widely accepted fact the K-ion exhibit lower desolvation energy at the electrode surface.\n6. Please report the average charge voltage of KVO. A charge plateau can be observed at ~2.5 V vs. Li⁺/Li, which is high for an anode and directly impacts cell-level energy density.\n7. The manuscript attributes superior SEI formation to the KVO electrode despite the use of a conventional electrolyte. How does KVO specifically contribute to the LiF-rich SEI?\n8. The text states that peaks in Figure S17 shift to lower 2θ values during lithiation and return upon delithiation. However, Figure S17 shows a shift to lower angles followed by a shift back to higher angles.\nCharacterization Concerns:\n1. Standard deviations or uncertainty estimates should be included for all XRD and NPD refinements.\n2. The XPS spectra in Figure S11 should be fitted to clearly distinguish V oxidation states.\n3. The authors claim the pre-edge peak in XAS is usually used to probe the valence state. In fact, the pre-edge is sensitive to local symmetry, while the edge position is typically used to infer oxidation state.\n4. In Figure 3c, the V L-edge should be clearly labeled (i.e., L1, L2 or L3).\n5. The caption of the XAS figure should indicate “photon energy”, not “binding energy” or “photo energy”. Please correct this across the manuscript.\n6. Figure S21 reveals a pronounced amorphous phase during discharge, suggesting that a crystalline-to-amorphous transition contributes to capacity. Consequently, claiming KVO is a “zero-strain” anode based solely on XRD refinement is misleading, as XRD cannot detect amorphous components.\n7. Using EXAFS R-space peak intensity alone to infer coordination number is inappropriate given the potential disorder revealed during lithiation. The peak intensity is also influenced by the Debye-Waller factor. Quantitative EXAFS fitting is necessary to support structural claims.\n8. Please cite the software used for structural and spectral analyses (e.g., GSAS, Athena, etc.)."}, "author_response": "This letter is in response to the reviewers’ comments regarding the manuscript entitled, “Intraplanar Percolation and Interplanar Bridge Enables Layered matrix for High-capacity/rate, Ultralong-cycling and “Zero-strain” Battery Operations”. The authors are grateful to the reviewers whose comments have helped the authors improve the scientific quality of the manuscript. Each comment has been addressed in the following and corresponding changes have been made in the manuscript.\nPoint-to-Point Response to Referees’ Comments and Suggestions\nReviewer # 1:\nIn this study, the authors developed anode materials with three-dimensional transport pathways and zero-strain behavior for fast-charging Li-ion batteries. In particular, they highlighted the structural advantages of K V O (KVO), which features a 3 5 14 layered structure with both intraplanar percolation and interplanar bridging. The authors investigated its electrochemical and structural properties through various methods, including rate capability test, long-term cycling, X-ray analyses, microscopic observations, and computational research. Furthermore, based on KVO, they extended their design strategy to similar compounds, K VP O and K MgV O , to highlight the 2 2 8 2 2 7 versatility of their approach. The overall scientific question addressed by this work is timely and relevant, and the data presented, particularly operando measurements, are sufficient to support the key claims. However, the manuscript currently lacks clarity in the presentation of electrochemical and structural analyses. The scientific logic should be more clearly articulated, and key details must be included to avoid confusion and strengthen the conclusions. Accordingly, the manuscript requires revision before it can be considered for publication in Nature Communications.\nResponse: We are very grateful to Reviewer 1 for the very positive comments.\nComment 1 of Reviewer 1:\nPlease revise the definition of XAS from “synchrotron radiation (XAS)” to “synchrotron radiation based X-ray absorption spectroscopy (XAS)”.\nResponse: We sincerely thank the reviewer for their meticulous attention to terminology accuracy. The suggested revision to explicitly define the acronym XAS as \"synchrotron radiation based X-ray absorption spectroscopy (XAS)\" is indeed scientifically precise and necessary for clarity. We have implemented this change throughout the manuscript, ensuring full consistency with the recommended terminology. This refinement enhances the technical rigor of our methodology description, and we greatly appreciate the reviewer's valuable guidance on this point. The manuscript has been carefully revised. Similar language errors have been corrected. The corresponding corrections have been added in lines 4-9 on page 5 in the revised manuscript.\nComment 2 of Reviewer 1:\nHow did the authors define C-rates or theoretical capacity for KVO, 1 C = 450 mA g-1? Is it based on the theoretical calculation? In the first discharge cycles, the KVO exhibited more than 700 mA h g-1 at 30 °C (Fig. S4), which appears to surpass the theoretical capacity. The authors should clarify this by presenting the electrochemical reaction equation and detailing the basis for the theoretical capacity used.\nResponse: We greatly thank the reviewer for this comment.\nFig. S8. (a)The formation energy as a function of the number of inserted Li+ ions in a unit cell. (b)The voltage required for lithium within KVO crystal structure.\nThe formation energy of lithiated KVO as a function of the number of inserted Li+ ions per unit cell is shown in Fig. S8a. The result suggests that at most ten Li+ ions can be stored in one unit cell, corresponding to atheoretical capacity of 450 mAh g−1. For the ~10 Li⁺ ions, the intercalation voltage remains relatively low (~1.1 V), indicating facile insertion with minimal polarization. In contrast, when attempting to intercalate the 11th Li⁺, the voltage sharply rises to ~2.0 V, reflecting a significant structural or electronic reconstruction within the host framework. (Fig. S8b) Combined with formation energy calculations: the 11th Li⁺ exhibits a positive formation energy. Thermodynamically, this process is intrinsically non-spontaneous; kinetically, the high voltage corresponds to a large activation barrier, which further suppresses Li⁺ insertion. These two factors synergistically explain why the 11th Li⁺ can be hardly intercalated.\nThe practical reversible capacity of the KVO sample is 377 mAh g−1 (84% of its theoretical capacity), which indicates that methods such as nanoscale may further increase the capacity. Based on the results of theoretical calculations, we summarize the electrochemical reaction equation as follows:\nK V O +xLi++xe- ↔ Li K V O (0≤x≤10)\n3 5 14 x 3 5 14\nIn the first discharge cycles, the KVO exhibited more than 700 mA h g−1，The phenomenon that the initial discharge capacity is higher than the theoretical capacity is commonly observed in anode materials discharged to low potentials (Li, X. et al. Perovskite-type SrVO as high-performance anode materials for lithium-ion batteries. 3 Adv. Mater. 2107262 (2021)；Ette, P.M. et al. Ordered 1D and 3D mesoporous Co O 3 4 structures: effect of morphology on Li-ion storage and high rate performance. Electrochim. Acta 310, 184–194 (2019)；Su, L. et al. Ultrathin layered hydroxide cobalt acetate nanoplates face-to-face anchored to graphene nanosheets for high-efficiency lithium storage. Adv. Funct. Mater. 27, 1605544 (2017).). The origin for the extra capacity can be attributed to 1) the formation of the SEI, resulting from electrolyte decomposition above 0.4 V vs. Li+/Li (Yoon, S. et al. Surface modification of graphite by coke coating for reduction of initial irreversible capacity in lithium secondary batteries. J. Power Sources 94, 68–73 (2001); Yazami, R. Surface chemistry and lithium storage capability of the graphite–lithium electrode. Electrochim. Acta 45, 87–97 (1999).). To clarify the initial irreversibility, we compared the capacity differences observed below 0.4 V during the first and second discharge cycles. The results indicate that the irreversible capacity loss occurs predominantly above 0.4 V, mirroring the behavior observed in graphite. (Fig. S5) This similarity suggests that the origin of the initial irreversibility is analogous, primarily attributable to SEI formation. The formation of the SEI layer also correlates with the anodic peak centered at ~0.7 V, which is exclusively observed in the dQ/dV vs. voltage plot for the first cycle (Fig. S21). We carefully observed the KVO material after the first week of discharge by using the characterization method of Cryo-electron microscopy. It could be found that a uniform SEI film with a thickness of approximately 15nm was formed on the surface of the KVO electrode (Fig. S6).\n2) intercalation-type anode materials that the inserted Li+ ions during the first lithiation cannot be totally extracted. XANES and XPS demonstrated that the slight difference from the original state was due to the incomplete extraction of trace Li+ ions in the initial cycle, but this only accounted for a small portion of the irreversible capacity. (Fig. 3a, Fig. 3c, and Fig. S22)\n3) Contribution from conductive additives (Super P, 10 wt% of electrode).\nSuper P alone delivers ~260 mAh g⁻¹ of irreversible capacity within 0.2−3V. (Fig. S7)\nFig. S5. Voltage-capacity plots for the discharge processes of (a)K V O /Li, and 3 5 14 (b)graphite/Li at 0.1C.\nFig.S6. Cryo-Transmission Electron Microscopy images of K V O samples showing 3 5 14 SEI films after first lithiation.\nFig. S7. Discharge/charge profiles of Super P/Li half cell at 0.1C within 0.2−3.0 V.\nTherefore, we observed much higher charge capacity at initial cycle, and when we use the active mass up to 90 %, the initial couloubic efficiency and charge capacity becomes more reasonable. The corresponding discussions have been added in lines 10-19 on page 7 in the revised manuscript and in lines 6-15 on page 13 in the revised Supporting information.\nComment 3 of Reviewer 1:\nPlease provide corresponding voltage profiles for long-term cycling data (Fig. 2c, 2d, 2e, 2g, 2h, and 2i). These are essential for understanding capacity retention and degradation mechanisms.\nResponse: We really appreciate the reviewer's suggestion regarding voltage profiles, which are indeed crucial for capacity retention and degradation mechanisms. The corresponding long-term cycling voltage curves for all datasets (originally in Fig. 2c-i) are now provided in Fig.S19. Significantly, these profiles show negligible voltage polarization shifts. This exceptional stability demonstrates: (1) The integrity of KVO electrode particles: Zero-strain behavior (validated by in-situ XRD/TEM) prevents microcrack-induced phase separation and grain boundary migration. (2) Stable single-phase transition: Smooth solid-solution reaction kinetics maintain single-phase evolution throughout cycling. (3) Fast three-dimensional transportation network: Constant voltage hysteresis confirms sustained electronic/ionic conductivity. The invariant voltage profiles directly correlate with KVO’s structural integrity and explain its 88.3% capacity retention after 19000 cycles. This comprehensive dataset now strengthens our degradation mechanism analysis. The corresponding discussions have been added in lines 18-19 on page 10 in the revised manuscript.\nFig.S19. Discharge/charge profiles of KVO/Li half cells at (a)1C, (b)5C, (c)10C rates and at 30℃. (d) Discharge/charge profiles for selected cycles of the KVO/Li half cells under conditions of 10C and at 60℃. (e) Discharge/charge profiles for selected cycles of the KVO/Li half cells under conditions of 1C and at -10℃. (f) Discharge/charge profiles for selected cycles of the LiFePO /KVO full cell under conditions of 2C at 4 30℃.\nComment 4 of Reviewer 1:\nIn Fig. 2g and 2h, the x-axis in voltage profiles should be labeled as specific capacity (mAh g-1) instead of cycle number.\nResponse: We thank the reviewer for pointing out this problem. In Fig. 2g and 2h, “specific capacity (mAh g-1)” has been changed to “cycle number” in the revised manuscript.\nFig. 2. Electrochemical performance of KVO half and full cells. (1 C = 450 mA g−1 for KVO) (a) Discharge/charge curves of KVO/Li cell at 0.1C for initial three cycles within 0.2–3.0 V at 30 ℃ (90%KVO: 5%Super P: 5%PVDF). (b)Discharge/charge profiles of KVO/Li half cell from 0.1 C to 10C within 0.2−3.0 V. (80%KVO: 10%Super P: 10%PVDF) (c, d, e) Cyclability of KVO/Li half cell after rate test. (80%KVO: 10%Super P: 10%PVDF) (f) The cycling stability of KVO is compared with that of other well-known long-life electrochemical energy-storage materials documented in prior studies. (g) Cyclability at 10 C over 8000 cycles at 60 ℃ (inset: Discharge/charge profiles from 0.1C to 10 C at 60 ℃, 80%KVO: 10%Super P: 10%PVDF). (h) Cyclability at 1 C over 2000 cycles at −10 ℃ (inset: Discharge/charge profiles from 0.1 C to 2 C at −10℃, 80%KVO: 10%Super P: 10%PVDF). (i) Cyclability of LiFePO /KVO full cell at 2 C over 1000 cycles, N/P ratio is 1.05, and the capacity is 4 calculated based on mass of KVO, (inset: Comparison of LiFePO /KVO with 4 LiFePO /Li Ti O in terms of specific capacity, operating voltage, and specific energy 4 4 5 12 density).\nComment 5 of Reviewer 1:\nFollowing Question 2, the formation cycles likely induce substantial structural changes in KVO due to the potential trapping of large amounts of lithium ions to stabilize the structure. Although XPS (Fig. S11) and in-situ XRD (Fig. S16-S19) are provided, they are insufficient to reveal the full extent of these transformations. Please include complementary structural characterizations, such as Raman spectroscopy, nuclear magnetic resonance (NMR), or infrared (IR) spectroscopy. Ex-situ analysis would be suitable for the supporting information.\nResponse: We greatly thank the reviewer for this insightful comment. We acknowledge that XPS and in-situ XRD, have inherent limitations and fail to fully capture the complete scope of structural changes during the formation cycles. In response to this suggestion, we have selected and performed ex situ Raman spectroscopy on the KVO electrodes for further structural characterization. As requested, we have performed ex situ Raman spectroscopy (Fig.S23), which reveals highly reversible behavior. Peaks at 930 and 962 cm-1 are attributed to the bending vibration of V-O bonds. ( Zhang, F. et al. Boosting charge carrier transport by layer-stacked Mn V O /V C heterostructures for x 2 6 2 wide-temperature zinc-ion batteries. Adv. Funct. Mater. 34, 2402071 (2024); Wang, A. et al. Building stabilized Cu Mn V O −2.16H O cathode enables an outstanding 0.17 0.03 2 5 2 room-/low-temperature aqueous Zn-ion batteries. Carbon Energy 6, e512 (2024); Ding, J. et al. In situ lattice tunnel distortion of vanadium trioxide for enhancing zinc ion storage. Adv. Energy Mater. 11, 2100973 (2021); Zhao, D. et al. Unlocking the capacity o f v a n a d i u m o x i P A d G E e irreversibility mechanism. The corresponding discussions have been added in lines 3-9 on page 12 in the revised manuscript.\nFig.S23. Ex-situ Raman patterns of KVO (first cycle at 0.1C).\nComment 6 of Reviewer 1:\nXPS results (Fig. S11) suggest a shift in V oxidation state from 5+ to a 4+-dominated state after formation cycles. However, XAS results (Fig. 3d, 3e) show only slight changes, indicating a 5+-dominated bulk even after cycling. This discrepancy suggests that oxidation state changes may be surface-confined. Still, if lithium trapping occurs, this should significantly affect the bulk chemistry. The authors should provide a systematic explanation reconciling these observations.\nResponse: We appreciate the reviewer’s insightful observation regarding the apparent discrepancy between the surface-sensitive XPS data (indicating V4+ dominance) and the bulk-sensitive XAS data (indicating persistent V5+ dominance) after cycling. We agree this suggests a surface-confined reduction process. The potential role of lithium trapping in this phenomenon is well-noted. We provide the following systematic explanation reconciling these observations:\nInherent Initial Lithium Irreversibility: As the reviewer inferred, the observed surface reduction is indeed linked to lithium trapping. This is a well-documented characteristic of many intercalation-type anode materials, including vanadium-based oxides. (Zhang, L. et al. Lithium lanthanum titanate perovskite as an anode for lithium ion batteries. Nat. Commun. 11, 3490 (2020); Lu, X. et al. Atomic-scale investigation on lithium storage mechanism in TiNb O . Energy Environ. Sci. 4, 2 7 2638–2644 (2011); Jayaraman, S. et al. Exceptional performance of TiNb O anode in 2 7 all one-dimensional architecture by electrospinning. ACS Appl. Mater. Interfaces 6, 8660–8666 (2014); Ma, S. et al. VPO : an all-climate lithium-storage material. Energy 5 Storage Mater. 46, 366 (2022); Liang, G. et al. Conductive Li Cr Si V O anode material: novel \"zero- 3.08 0.02 0.09 0.9 4 strain\" characteristic and superior electrochemical Li+ storage. Adv. Energy Mater. 10, 1904267 (2020).) During the initial lithiation, a portion of Li+ ions becomes irreversibly incorporated into the host structure. This trapped lithium contributes to the observed reduction of vanadium (V5+ → V4+) at the surface and near-surface regions, as detected by XPS. We observed slight oxygen defects on the surface through EELS characterization. During initial lithiation, Li⁺ ions are attracted to these sites to compensate the charge. Once there, the Li⁺ ions become strongly stabilized within the defect structure and by the accompanying reduction of nearby metal ions ( V5+ to V4+). This makes it energetically unfavorable and kinetically difficult for these Li⁺ ions to be removed during delithiation, resulting in irreversible trapping. The existence of O defects proved the fundamental cause of the slight charge heterogeneity from bulk to surface. (Fig.S24)\nFig.S24. V L-edge spectra and O K-edge spectra of KVO.\nThe apparent inconsistency in vanadium valence evolution (bulk vs. near-surface) arises from fundamental differences in thermodynamic stability between bulk and surface environments, as revealed by our DFT calculations. Below is a step-by-step explanation: Formation energies for both V5+, V4+, and V3+ in the bulk phase are negative (Table S5 and Fig.S25). This confirms that V4+ / V3+ is thermodynamically stable in the bulk, enabling reversible oxidation to V5⁺ during cycling. In contrast, the formation energy for V4+ at the surface is positive (Table S5 and Fig.S25), indicating that V4+ is thermodynamically unstable in near-surface regions. Once V5+ is reduced to V4+ at the surface, the high energy barrier prevents its reversion to V5+, leading to irreversible valence trapping. The surface instability stems from incomplete oxygen coordination, as demonstrated by the EELS experiment.\nTable S5 Calculate the relevant parameters of the V valence state stability theory between the surface and the bulk.\nBulk nLi+KVO 3K nLi V5O14 Formation\nenergy(eV)\nV5+（K V O ） -148.6894 -3.0525 0 -133.927 -11.7099\n3 5 14\nV4+（K Li V O ） -155.4618 -3.0525 -9.4995 -133.927 -8.9828\n3 5 5 14\nV3+（ K Li V O ） -157.4455 -3.0525 -18.999 -133.927 -1.467\n3 10 5 14\nSurface nLi+KVO 6K nLi V10O28 Formation\nenergy(eV)\nV5+（K V O ） -281.5605 -6.105 0 -267.854 -7.6015\n3 5 14\nV4+（K Li V O ） -286.1373 -6.105 -18.999 -267.854 1.6788\n3 5 5 14\nFig.S25. Theoretical calculation diagram of the V valence state stability of the surface and bulk phase.\nMagnitude of the Surface-Bulk Gradient: While a Li concentration gradient inevitably exists between the Li-rich surface and the bulk after initial cycling, the magnitude of this gradient in our material is remarkably small compared to typical electrode materials. In conventional high-strain materials, large surface-bulk Li concentration gradients commonly develop during cycling, often leading to catastrophic failure. For instance: Silicon Anodes: During lithiation, the surface can form highly lithiated phases (e.g., Li Si ) while the core remains less lithiated or even 15 4 pristine Si. This massive differential expansion (>300%) generates enormous stresses, inevitably causing particle pulverization and rapid capacity fade. (Cui, Y. et al. Silicon anodes. Nat. Energy 6, 995−996 (2021); Huo, H. et al. Chemo-mechanical failure mechanisms of the silicon anode in solid-state batteries. Nat. Mater. 23, 543−551 (2024).) Layered Transition Metal Oxides: Surface reconstruction layers can form with significantly different Li content and TM oxidation states compared to the bulk. The associated large lattice parameter mismatches create severe interfacial stresses, driving crack formation and propagation through the particle bulk. (Zuo, W. et al. Microstrain screening towards defect-less layered transition metal oxide cathodes. Nat. Nanotechnol. 19, 1644−1653 (2024); Jousseaume, T. et al. Strain and collapse during lithiation of layered transition metal oxides: a unified picture. Energy Environ. Sci. 17, 2753 (2024).) Conversion Anodes (e.g., Fe O ): The reaction front often proceeds 2 3 inward, creating a sharp boundary between fully converted (Li-rich) surface regions and unconverted (Li-poor) core regions, generating extreme internal stresses. (Bresser, D. et al. Leveraging valuable synergies by combining alloying and conversion for lithium-ion anodes. Energy Environ. Sci. 9, 3348 (2016).)\nZero-Strain: The Enabling Factor for Tolerating the Gradient: The critical factor allowing our material to maintain stability despite the existence of some surface Li trapping and a small Li concentration gradient is its exceptional near-zero-strain characteristic during (de)intercalation. In conventional anode materials, even moderate Li concentration gradients during (de)intercalation can induce substantial internal mechanical stresses due to lattice parameter changes.( Jiang, T. et al. Partially reduced titanium niobium oxide: a high-performance lithium-storage material in a broad temperature range. Adv. Sci. 9, 2105119 (2022); Wang, W. et al. Conductive LaCeNb O with a very open A-site-cation-deficient perovskite structure: a fast- and 6 18 stable-charging Li⁺-storage anode compound in a wide temperature range. Adv. Energy Mater. 12, 2200656 (2022); Lv, C. et al. Rational design and synthesis of nickel niobium oxide with high-rate capability and cycling stability in a wide temperature range. Adv. Energy Mater. 11, 2102550 (2021).) This often leads to particle cracking, increased impedance, and rapid capacity fade. However, our material exhibits an exceptionally \"zero-strain\" characteristic during cycling. The “zero-strain” characteristic means that the lattice parameters are highly stable upon Li⁺ insertion/extraction. Therefore, even the slight Li concentration gradient present between the surface-reduced region and the bulk (V5+-dominant) core generates minimal mechanical stress or strain within the particles. This lack of significant stress prevents the propagation of cracks or defects that would otherwise exacerbate the gradient or lead to bulk phase transformations. The minimal lattice variation prevents the buildup of destructive mechanical stresses that would normally be caused by such a chemical gradient in a conventional material. This intrinsic stability allows the material to tolerate the surface-confined chemical changes without detrimental impact on the bulk structure or the overall electrochemical performance.\nTherefore, the observed surface reduction is consistent with initial lithium irreversibility, while the persistence of the bulk V5+ state is maintained due to the confinement of the process and, critically, the stress-absorbing capability of the near-zero-strain lattice. This stands in stark contrast to high-strain materials where similar or smaller gradients often lead to severe mechanical degradation. This unique combination (shallow trapping + small gradient + zero-strain ) underpins the excellent structural and cycling stability observed. The corresponding discussions have been added in lines 9-17 on page 12 in the revised manuscript.\nComment 7 of Reviewer 1:\nEven in the absence of significant peak shifts during cycling, notable changes in peak intensities are observed after 500 cycles (Fig. S23). The authors should discuss the origin of these changes-whether due to preferred orientation, phase transformation, or structural degradation.\nRespons: We thank the reviewer for their careful observation regarding the changes in relative peak intensities observed in the XRD patterns of the electrode laminates after 500 cycles (Fig. S23), despite the absence of significant peak shifts. We acknowledge the importance of understanding the origin of these intensity variations. Based on our analysis, we attribute these changes primarily to differences inherent to measuring electrode sheets versus pure powders, coupled with the effects of preferred orientation, rather than indicating bulk phase transformation or structural degradation. Our reasoning is as follows:\nInherent Differences Between Electrode sheets and Powder XRD:\nCrucially, the XRD data presented in Fig.S46. was collected directly on cycled electrode sheets (containing active material, conductive carbon, binder, and current collector). This differs fundamentally from the XRD patterns typically collected on pure, loose powder samples (like those often shown for pristine materials). The amount of active material contributing to the XRD signal within the X-ray beam footprint on a thin electrode sheet is significantly less than in a thick powder sample. This reduced effective sample volume naturally leads to lower overall peak intensities and can make the pattern appear noisier. The active material particles within the electrode laminate are embedded in a complex matrix with conductive carbon and binder. This matrix can cause slight variations in particle distribution, orientation, and even partial shielding, further contributing to potential deviations in relative peak intensities compared to a homogeneous powder.\nFig.S46. (a) The Rietveld refined XRD pattern of the KVOpowder sample (b) Rietveld refined XRD pattern of the electrode sheets prepared by KVO.\nDominant Role of Preferred Orientation:\nThe process of electrode fabrication (slurry coating, drying, and calendering/rolling) inherently applies shear forces and pressure to the active material particles. This mechanical treatment is well-known to induce preferred crystallographic orientation (texture) within the electrode laminate. Particles tend to align with specific crystallographic planes parallel to the electrode substrate surface. As a result, XRD peaks corresponding to planes parallel to the substrate will be enhanced, while peaks from planes perpendicular to the substrate will be suppressed, compared to a randomly oriented powder sample. This phenomenon is well-documented in layered/2D materials (e.g., graphite, LiCoO ) where texture significantly impacts laminate XRD 2 intensities. (Chen, Y. et al. Insight into the intrinsic mechanism of improving electrochemical performance via constructing the preferred crystal orientation in lithium cobalt dioxide. Chem. Eng. J. 399, 125708 (2020); Jenkins, J.C. & Jenkins, G.M. An analysis of preferred orientation induced in stretched pitch mesophases and derived graphites. Carbon 24, 93−100 (1986).) Research shows that KVO, as a new type of layered structure with interplanar bridges and intraplanar percolation, also undergoes a change in preferred orientation after rolling. (Fig.S47)\nFig.S47. (a) Rietveld refined the XRD patterns of the unrolled KVO electrode sheet. (b) Rietveld refined the XRD patterns of the unrolled KVO electrode sheet after 500 cycles. (c) Rietveld refined the XRD patterns of the rolled KVO electrode sheet. (b) Rietveld refined the XRD patterns of the rolled KVO electrode sheet after 500 cycles.\nAssessing Phase Transformation and Degradation:\n1)Peak Position Stability: The key indicator for bulk crystallographic phase changes is the position of the diffraction peaks. The reviewer correctly notes the absence of significant peak shifts in Fig. S23. This strongly suggests that the long-range crystal structure framework of the bulk active material remains remarkably stable even after 500 deep cycles. There is no evidence for a major phase transformation within the detection limits of lab XRD. We used TEM characterization to observe the changes in the number of different cycles of KVO. Surprisingly, even after tens of thousands of cycles, it could still be observed that the KVO particles remained intact. These tem characterizations fully demonstrate that the KVO electrode has excellent long-term structural stability.( Fig.S48)\nFig.S48. The structural evolution of KVO at different charge/discharge states.\n2) Unchanged peak widths (ΔFWHM < 15% vs. pristine) definitively exclude particle fragmentation, amorphization, or microstrain accumulation (Table S8)—processes that typically broaden peak in degradation-prone materials (e.g., Si, Fe O ); (Cui, Y. et al. 3 4 Silicon anodes. Nat. Energy 6, 995−996 (2021); Huo, H. et al. Chemo-mechanical failure mechanisms of the silicon anode in solid-state batteries. Nat. Mater. 23, 543−551 (2024); Bresser, D., Passerini, S. & Scrosati, B. Leveraging valuable synergies by combining alloying and conversion for lithium-ion anodes. Energy Environ. Sci. 9, 3348 (2016).)\nTable S8. Comparison ofΔFWHM values before and after KVO electrode cycling.\nKVO Crystal (001) (111) (200) (210) (300) (110)\nplane\nPristine 0.1093 0.1083 0.0983 0.0972 0.1000 0.1112\nAfter 500 cycles 0.122 0.118 0.1117 0.1008 0.1188 0.1163\n3) Electrochemically, >88.3% capacity retention and identical voltage profiles after 19000 cycles demonstrate preserved bulk crystallinity.\nCollectively, these invariabilities confirm that the peak intensity variations originate solely from electrode-level effects (texture evolution/matrix interactions), not bulk material degradation. The corresponding discussions have been added in lines 10-14 on page 17 of the revised manuscript, and in lines 4-15 on page 59, lines 6-13 on page 60, and lines 4-8 on page 62 of the revised Supporting information.\nComment 8 of Reviewer 1:\nPlease include the voltage profiles corresponding to the GITT measurements (Fig. S26).\nResponse: Thank you so much for your good suggestion, we've added the voltage profiles corresponding to GITT measurements to Fig.S53.\nFig.S53.GITT tests of the second cycle of KVO/Li half-cells at 30, 60 and -10℃ within the voltage range of 0.2 to 3.0V.\nComment 9 of Reviewer 1:\nIt is strongly recommended to discuss the formation mechanisms of KVO using computational results as demonstrated in Fig. 5.\nResponse: We thank the reviewer for this comment. As shown in Figure 5d, Li+ can insert into 1a, 3c(1), and 3(2) sites of the K V O bulk. Fig.S51. displays the optimized 3 5 14 crystal structures and Li+ occupations of the most stable configurations in Li K V O , x 3 5 14 showing that 1a sites are the most favorable for the initial Li+ insertion with a binding energy of −1.49 eV and that the binding energies follow the order of 1a < 3c(1) < 3c(2). Interestingly, in the optimized crystal model, since the V-O bonds are in a negatively charged state, they are more likely to adsorb Li+, and the Li+ are more likely to shift towards the V-O position. For K V O , the maximum inserted Li+ number is as high 3 5 14 as 10 (Fig.S9). Within the range of 1≤x≤10, the binding energies of xLi and K V O 3 5 14 are both negative, indicating that K V O is suitable for utilization in LiBs. The 3 5 14 corresponding discussions have been added in lines 10-14 on page 20 in the revised manuscript.\nFig.S9. Optimized structures and Li+ occupations of the most stable configurations in Li K V O .\nx 3 5 14\nComment 10 of Reviewer 1:\nPlease provide more details of electrochemical data for K VP O and K MgV O , such 2 2 8 2 2 7 as voltage profiles and rate capability.\nResponse: Thank you so much for your suggestion. The voltage profiles and rate capability data for K VP O and K MgV O have been added to Fig.S57. and Fig.S58. 2 2 8 2 2 7\nFig.S57. Electrochemical performance of K VP O half cells. (1C=320 mA g−1) (a) 2 2 8 Discharge/charge curves of K VP O /Li cell at 0.1C for initial four cycles within 0.01– 2 2 8 3.0 V at 30 ℃. (b) Corresponding rate performance from 0.5C to 10C. (c, d) CV curves of K VP O /Li half cell within 0.01−3.0V. (e) Discharge/charge profiles of K VP O /Li 2 2 8 2 2 8 half cell from 0.5C to 10C within 0.01−3.0 V. (f) The corresponding voltage profiles during the cycle process.\nFig.S58. Electrochemical performance of K MgV O half cells. (1C=340 mA g−1) (a) 2 2 7 Discharge/charge curves of K MgV O /Li cell at 0.1C for initial four cycles within 2 2 7 0.01–3.0 V at 30 ℃. (b) Corresponding rate performance from 0.5C to 10C. (c) The corresponding voltage profiles during the cycle process.\nReviewer #2:\nThis manuscript introduces the design of “zero-strain” K3V5O14 (KVO) anode materials with high capacity, rapid Li ion conductivity, and enhanced structural stability. The authors characterize the structure of the anode material in detail and elucidate the relationship between the structural characteristics and the high energy storage performance. Combined with advanced characterization technology and computational simulation, the lithium storage mechanism of the KVO anode was studied from the perspectives of the chemical state, bulk structure, and microscopic scale. The authors demonstrate a general principle for searching and designing materials with high energy storage performance and implement it in K VP O and K MgV O anodes. This 2 2 8 2 2 7 manuscript is clearly written, and the data are nicely presented and analyzed. I recommend publication after some minor corrections are made, as suggested below.\nResponse: We are so grateful to Reviewer 2 for the very positive comments.\nComment 1 of Reviewer 2:\nWhy does the discharge specific capacity increase during the charge-discharge cycling in Figure 2?\nResponse: We really appreciate the reviewer for this comment. In fact, this capacity-enhancement phenomenon has commonly been found in transition-metal based anode materials working at low potentials (Ette, P.M. et al. Ordered 1D and 3D mesoporous Co O structures: effect of morphology on Li-ion storage and high rate performance. 3 4 Electrochim. Acta 310, 184–194 (2019); Kim, H. et al. Exploring anomalous charge storage in anode materials for next generation Li rechargeable batteries. Chem. Rev. 120, 6934−6976 (2020); Wang, J. et al. Oxidizing solid Co into hollow Co O within 3 4 electrospun (carbon) nanofibers towards enhanced lithium storage performance. J. Mater. Chem. A 7, 3024–3030 (2019); Su, L. et al. Ultrathin layered hydroxide cobalt acetate nanoplates face-to-face anchored to graphene nanosheets for high-efficiency lithium storage. Adv. Funct. Mater. 27, 1605544 (2017); Wei, W. et al. 3D graphene foams cross-linked with pre-encapsulated Fe O nanospheres for enhanced lithium 3 4 storage. Adv. Mater. 25, 2909−2914 (2013)). The capacity enhancement during the cycling can be attributed to the KVO activation. The KVO activation: The gradual electrolyte penetration into the KVO particles can lead to the activation process. As a result, more and more active sites can be exposed during the repeated charging–discharging processes (Wu, H. et al. Stable Li-ion battery anodes by in situ polymerization of conducting hydrogel to conformally coat silicon nanoparticles. Nat. Commun. 4, 1943 (2013); Liu, P. et al. Hierarchical porous intercalation-type V O as 2 3 high-performance anode materials for Li-ion batteries. Chem. Eur. J. 23, 7538−7544 (2017).). The activation can be confirmed by the EIS tests of the KVO/Li half cell (Fig. S20), which reveal gradually smaller charge-transfer resistance during the cycling (Sun, Y. et al. SnO nanoparticles encapsulated by curved graphite layers as anode materials 2 for Li-ion batteries with high performances. J. Alloys Compd. 683, 191–197 (2016); Ette, P.M. et al. Ordered 1D and 3D mesoporous Co O structures: effect of morphology 3 4 on Li-ion storage and high rate performance. Electrochim. Acta 310, 184–194 (2019).). The corresponding discussions have been added in lines 20-21 on page 10 in the revised manuscript.\nFig.S20. Nyquist plots of KVO after 0, 100, 200, 500 and 2000 cycles at 30 oC.\nComment 2 of Reviewer 2:\nPlease explain the reason why the first circle discharge capacity is significantly higher than the charging capacity in Figures S7 and S8, and analyze the reasons for the inconsistency with the illustrations in Figures 2g and 2h.\nResponse: We thank the reviewer so much for raising this question. The phenomenon that the initial discharge capacity is higher than the theoretical capacity is commonly observed in anode materials discharged to low potentials (Ette, P.M. et al. Ordered 1D and 3D mesoporous Co O structures: effect of morphology on Li-ion storage and high 3 4 rate performance. Electrochim. Acta 310, 184–194 (2019); Su, L. et al. Ultrathin layered hydroxide cobalt acetate nanoplates face-to-face anchored to graphene nanosheets for high-efficiency lithium storage. Adv. Funct. Mater. 27, 1605544 (2017); Li, X. et al. Perovskite-type SrVO as high-performance anode materials for lithium- 3 ion batteries. Adv. Mater. 33, 2107262 (2021).).\nThe origin for the extra capacity can be attributed to 1) the formation of the SEI, resulting from electrolyte decomposition above 0.4 V vs. Li+/Li，\n2) the high irreversible capacity caused by conductive carbon (Super P) at low potentials,\n3) intercalation-type anode materials that the inserted Li+ ions during the first lithiation cannot be totally extracted. We thank the reviewer for their careful observation. Figures 2g and 2h indeed both depict the second cycle at 0.1C. The corresponding discussions have been added in lines 10-16 on page 7 in the revised manuscript.\nComment 3 of Reviewer 2:\nFigure 3 demonstrates the formation of a LiF-rich SEI layer, which is beneficial for cycling stability. But what is the formation mechanism of such LiF-Rich layer?\nResponse: We sincerely appreciate the reviewer's insightful question regarding the formation mechanism of the LiF-rich SEI layer observed in our KVO anode. This is indeed a critical aspect we had not fully explored in the original manuscript, as our primary focus was on elucidating the unique structural advantages of KVO’s integrated interplanar bridges and intraplanar percolation pathways for rapid Li⁺ diffusion and zero strain characteristic. The reviewer’s valuable comment prompted us to conduct dedicated follow-up investigations to uncover the fundamental origin of this phenomenon. Our comprehensive analysis reveals a material-intrinsic mechanism rooted in interfacial chemistry, which we now summarize as follows:\nExperimental Validation of Universality\nTo determine whether LiF enrichment was specific to KVO or inherent to the vanadate framework, we further performed comparative TOF-SIMS studies on isostructural K MgV O (Fig. S31). Crucially, this K MgV O exhibited the same pronounced LiF 2 2 7 2 2 7 dominance in its SEI.\nFig.S31. (a)The TOF-SIMS depth profiles of SEI layers for LiF-, CO - and C HO- in 3 2 KMgV O electrodes. (b)The 3D views of LiF-, CHO - and C HO- in the TOF-SIMS 2 7 2 2 sputtered volumes of SEI layers for K MgV O electrodes.\n2 2 7\nAtomistic Mechanism: Fluorophilic Vanadium Surfaces\nDFT calculations (PBE+U) uncovered a surface-selective reaction pathway driving LiPF₆ decomposition: Fluorine (χ = 3.98) exhibits significantly higher electronegativity than oxygen (χ = 3.44), enabling stronger ionic-covalent hybrid bonds with electropositive vanadium sites (χ = 1.63). From a thermodynamic perspective, this is conducive to the selective adsorption of LiPF on the electrode surface. Compared with 6 common electrolyte decomposition products such as EC (-0.59 eV) or DEC (-0.056 eV), the calculated adsorption energy (-0.708 eV) is significantly higher. When LiPF₆ approaches the KVO surface, PF₆⁻ anions undergo site-specific chemisorption (Fig. S32, Fig.3k and Table S6). The LiF-rich SEI originates from the inherent fluorophilicity of exposed vanadium centers in layered vanadates, where strong V–F bonding drives preferential PF₆⁻ decomposition and localized LiF precipitation. (Fig. S33 and Fig. S34) This mechanism—generalizable across vanadate frameworks—represents an intrinsic interfacial property that complements KVO’s structural advantages. We deeply thank the reviewer for identifying this significant dimension, their professional knowledge has profoundly enriched our understanding of the formation mechanism of SEI. The corresponding discussions have been added in lines 1-15 on page 15 in the revised manuscript.\nFig.S32. Theoretical study on the affinity energy and dissolution barrier of different electrolyte components for KVO. (a)LiPF (b) EC (c)DEC (d) and DMC.\nTable S6. Parameters related to the affinity energy of different electrolyte components for KVO and the dissolution potential barrier.\nLiPF +KVO 1 LiPF ∆E 0Li+KVO Formation\n6 6\nenergy\n/(eV/atom)\n-156.05 -7.32 -1.415 -147.31 -0.708\nEC+KVO 1EC ∆E 0Li+KVO Formation\nenergy\n/(eV/atom)\n-218.92 -75.71 -1.18 -142.03 -0.59\nDEC+KVO 1DEC ∆E 0Li+KVO Formation\nenergy\n/(eV/atom)\n-242.85 -99.92 -0.9 -142.03 -0.05625\nDMC+KVO 1DMC ∆E 0Li+KVO Formation\nenergy\n/(eV/atom)\n-240.09 -89.31 -0.94 -147.75 -0.08\nFig.S33. The affinity energy and desolvation barrier of Li+ solvation structure for KVO.\nFig.S34. DFT calculation of Li+ solvation structure at the KVO interphase. (a)LEMC (b) LEDC (c)Li CO (d) and LiF.\n2 3\nComment 4 of Reviewer 2:\nOn page 4, line 89, the manuscript says that there are dual redox couples (V5+/V3+), but this appears to be only one redox couple.\nResponse: We greatly appreciate the reviewer's careful reading and valid clarification request. While V4⁺ intermediates do exist during the multi-step reduction process, our combined hard X-ray absorption spectroscopy (XANES at K-edge, Fig.3d) quantitatively reveals that the overall average vanadium oxidation state reaches +3.5 at deep discharge (0.2 V), demonstrating substantial multi-electron transfer (beyond single-electron reactions) and confirming KVO's dual-redox capability.\nFig.3d Valence changes of the bulk phases of V measured at different discharge/charge states.\nComment 5 of Reviewer 2:\nThe insert in Figure 2i is not clear enough.\nResponse: We really thank the reviewer for pointing out this question. We have replaced the unclear insert in Figure 2i with a high-resolution version for enhanced clarity. The revised figure now distinctly displays the intended details. (Fig.2i)\nFig.2i. Cyclability of LiFePO /KVO full cell at 2C over 1000 cycles (inset: 4 Comparison of LiFePO /KVO with LiFePO /Li Ti O in terms of specific capacity, 4 4 4 5 12 operating voltage, and specific energy density).\nComment 6 of Reviewer 2:\nPlease explain the phenomenon in Figures 3d and 3e where the valence state changes in the bulk and near-surface are not completely consistent.\nResponse: We sincerely thank the reviewer for this thoughtful question and their careful examination of our data. The observed difference in vanadium valence evolution between the bulk and near-surface regions is indeed a critical point. Below, we clarify this phenomenon step by step based on our evidence:\nInherent Initial Lithium Irreversibility: As the reviewer correctly infers, the observed surface reduction is indeed linked to lithium trapping. This is a well-documented characteristic of many intercalation-type anode materials, including vanadium-based oxides. (Zhang, L. et al. Lithium lanthanum titanate perovskite as an anode for lithium ion batteries. Nat. Commun. 11, 3490 (2020); Lu, X. et al. Atomic-scale investigation on lithium storage mechanism in TiNb O . Energy Environ. Sci. 4, 2 7 2638–2644 (2011); Jayaraman, S. et al. Exceptional performance of TiNb O anode in 2 7 all one-dimensional architecture by electrospinning. ACS Appl. Mater. Interfaces 6, 8660–8666 (2014); Ma, S. et al. VPO : an all-climate lithium-storage material. Energy 5 Storage Mater. 46, 366 (2022); Liang, G. et al. Conductive Li Cr Si V O anode 3.08 0.02 0.09 0.9 4 material: novel \"zero-strain\" characteristic and superior electrochemical Li+ storage. Adv. Energy Mater. 10, 1904267 (2020).) During the initial lithiation, a portion of Li+ ions becomes irreversibly incorporated into the host structure. This trapped lithium contributes to the observed reduction of vanadium (V5+ → V4+) at the surface and near-surface regions, as detected by XPS. We observed slight oxygen defects on the surface through EELS characterization. During initial lithiation, Li⁺ ions are attracted to these sites to compensate the charge. Once there, the Li⁺ ions become strongly stabilized within the defect structure and by the accompanying reduction of nearby metal ions ( V5+ to V4+). This makes it energetically unfavorable and kinetically difficult for these Li⁺ ions to be removed during delithiation, resulting in irreversible trapping. The existence of O defects proved the fundamental cause of the slight charge heterogeneity from bulk to surface. (Fig.S24)\nFig.S24. V L-edge spectra and O K-edge spectra of KVO.\nThe apparent inconsistency in vanadium valence evolution (bulk vs. near-surface) arises from fundamental differences in thermodynamic stability between bulk and surface environments, as revealed by our DFT calculations. Below is a step-by-step explanation: Formation energies for both V5+, V4+, and V3+ in the bulk phase are negative . This confirms that V4+ / V3+ is thermodynamically stable in the bulk, enabling reversible oxidation to V5+ during cycling. In contrast, the formation energy for V4+ at the surface is positive (Table S5 and Fig.S25), indicating that V4+ is thermodynamically unstable in near-surface regions. Once V5+ is reduced to V4+ at the surface, the high energy barrier prevents its reversion to V5+, leading to irreversible valence trapping. The surface instability stems from incomplete oxygen coordination, as demonstrated by the EELS experiment.\nTable S5 Calculate the relevant parameters of the V valence state stability theory between the surface and the bulk.\nBulk nLi+KVO 3K nLi V O Formation\n5 14\nenergy(eV)\nV5+（K V O ） -148.6894 -3.0525 0 -133.927 -11.7099\n3 5 14\nV4+（K Li V O ） -155.4618 -3.0525 -9.4995 -133.927 -8.9828\n3 5 5 14\nV3+（K Li V O ） -157.4455 -3.0525 -18.999 -133.927 -1.467\n3 10 5 14\nSurface nLi+KVO 6K nLi V O Formation\n10 28\nenergy(eV)\nV5+（K V O ） -281.5605 -6.105 0 -267.854 -7.6015\n3 5 14\nV4+（K Li V O ） -286.1373 -6.105 -18.999 -267.854 1.6788\n3 5 5 14\nFig.S25. Theoretical calculation diagram of the V valence state stability of the surface and bulk phase.\nMagnitude of the Surface-Bulk Gradient: While a Li concentration gradient inevitably exists between the Li-rich surface and the bulk after initial cycling, the magnitude of this gradient in our material is remarkably small compared to typical electrode materials. In conventional high-strain materials, large surface-bulk Li concentration gradients commonly develop during cycling, often leading to catastrophic failure. For instance: Silicon Anodes: During lithiation, the surface can form highly lithiated phases (e.g., Li Si ) while the core remains less lithiated or even 15 4 pristine Si. This massive differential expansion (>300%) generates enormous stresses, inevitably causing particle pulverization and rapid capacity fade. (Cui, Y. et al. Silicon anodes. Nat. Energy 6, 995–996 (2021); Huo, H. et al. Chemo-mechanical failure mechanisms of the silicon anode in solid-state batteries. Nat. Mater. 23, 543–551 (2024).) Layered Transition Metal Oxides: Surface reconstruction layers can form with significantly different Li content and TM oxidation states compared to the bulk. The associated large lattice parameter mismatches create severe interfacial stresses, driving crack formation and propagation through the particle bulk. (Zuo, W. et al. Microstrain screening towards defect-less layered transition metal oxide cathodes. Nat. Nanotechnol. 19, 1644–1653 (2024); Jousseaume, T. et al. Strain and collapse during lithiation of layered transition metal oxides: a unified picture. Energy Environ. Sci. 17, 2753 (2024).) Conversion Anodes (e.g., Fe O ): The reaction front often proceeds 2 3 inward, creating a sharp boundary between fully converted (Li-rich) surface regions and unconverted (Li-poor) core regions, generating extreme internal stresses. (Bresser, D., Passerini, S. & Scrosati, B. Leveraging valuable synergies by combining alloying and conversion for lithium-ion anodes. Energy Environ. Sci. 9, 3348 (2016).)\nZero-Strain: The Enabling Factor for Tolerating the Gradient: The critical factor allowing our material to maintain stability despite the existence of some surface Li trapping and a small Li concentration gradient is its exceptional near-zero-strain characteristic during (de)intercalation. In conventional anode materials, even moderate Li concentration gradients during (de)intercalation can induce substantial internal mechanical stresses due to lattice parameter changes. (Jiang, T. et al. Partially reduced titanium niobium oxide: a high-performance lithium-storage material in a broad temperature range. Adv. Sci. 9, 2105119 (2022);Wang, W. et al. Conductive LaCeNb O with a very open A-site-cation-deficient perovskite structure: a fast- and 6 18 stable-charging Li⁺-storage anode compound in a wide temperature range. Adv. Energy Mater. 12, 2200656 (2022); Lv, C. et al. Rational design and synthesis of nickel niobium oxide with high-rate capability and cycling stability in a wide temperature range. Adv. Energy Mater. 11, 2102550 (2021).) This often leads to particle cracking, increased impedance, and rapid capacity fade. However, our material exhibits an exceptionally \"zero-strain\" characteristic during cycling. The “zero-strain” characteristic means that the lattice parameters are highly stable upon Li⁺ insertion/extraction. Therefore, even the slight Li concentration gradient present between the surface-reduced region and the bulk (V5+-dominant) core generates minimal mechanical stress or strain within the particles. This lack of significant stress prevents the propagation of cracks or defects that would otherwise exacerbate the gradient or lead to bulk phase transformations. The minimal lattice variation prevents the buildup of destructive mechanical stresses that would normally be caused by such a chemical gradient in a conventional material. This intrinsic stability allows the material to tolerate the surface-confined chemical changes without detrimental impact on the bulk structure or the overall electrochemical performance.\nTherefore, the observed surface reduction is consistent with initial lithium irreversibility, while the persistence of the bulk V5+ state is maintained due to the confinement of the process and, critically, the stress-absorbing capability of the near-zero-strain lattice. This stands in stark contrast to high-strain materials where similar or smaller gradients often lead to severe mechanical degradation. This unique combination (shallow trapping + small gradient + zero-strain ) underpins the excellent structural and cycling stability observed. The corresponding discussions have been added in lines 9-17 on page 12 in the revised manuscript.\nComment 7 of Reviewer 2:\nPlease provide an enlarged display of the strain region in Figure 4f.\nResponse: We thank the reviewer for this helpful suggestion. An enlarged view of the strain region in Figure 4f has been added as Fig.S40. This higher-resolution image clearly reveals pronounced strain-stripe migration (indicated by arrows). Critically, these microstructural changes occur without detectable alterations in particle morphology or volume-as evidenced by in-situ optical microscopic observation dimensions (Fig.S42) and <0.11% volumetric variation from in-situ XRD (Fig. 5b). This apparent paradox confirms KVO's unique self-adaptive strain accommodation mechanism, where dynamic bond-length adjustments prevent cumulative damage while maintaining macroscopic zero-strain characteristics.\nFig.S40. Enlarged in-situ TEM images of Li+ intercalation in KVO.\nReviewer # 3:\nMa et al. report the use of K V O (KVO) as a fast-charging anode material for Li-ion 3 5 14 batteries, presenting extensive data on its phase evolution and surface chemistry during cycling. However, in my view, the V O -based anode system has already been widely 2 5 explored, and this study does not deliver sufficient conceptual novelty or electrochemical advancement to meaningfully progress the field. Furthermore, the characterization data contain several errors, lack standardization, and are not able to support the authors’ conclusions. I do not recommend publication in Nature Communications.\nResponse: We sincerely appreciate the reviewer's insightful perspective on the extensive research landscape of vanadium-based anode materials. We fully acknowledge the significant contributions of prior studies on V O systems in 2 5 advancing Li-ion battery technology. While this work, fundamentally diverges from the conventional approaches, by introducing a novel structural design paradigm rather than focusing solely on compositional innovation.\nWe propose a unique layered architecture integrating interplanar bridges and intraplanar percolation pathways. This design strategically positions large-radius inactive ions between active TM layers to achieve three critical functions:(1) Expand interlayer spacing without sacrificing Li⁺ storage sites,(2) Create robust pillars constraining TM-layer expansion (ΔV =~0.11%),(3) Enable ultrafast ion transport through crystallographically aligned tunnels. The selection of K V O as a model 3 5 14 system demonstrates how this universal design principle—not vanadium chemistry itself—delivers unprecedented performance: The as-prepared KVO compound displays outstanding electrochemical performance, such as high capacity of 377 mAh g−1, excellent rate performance (approximately 146 mAh g−1 at 10 C) and superior long-term cyclability retaining 88.3% capacity after 19,000 cycles.\nWe believe this structural innovation establishes a new direction for next-generation anode design. The paradigm's versatility is further validated by applying it to K MgV O , K VP O and Rb V O —achieving similarly exceptional performance. 2 2 7 2 2 8 3 5 14 We respectfully contend that such fundamental advances in material architecture warrant dissemination to accelerate battery research.\nMajor Comments:\nComment 1 of Reviewer 3:\nWhy is the V5+/V3+ redox couple in KVO located within the anode voltage window, despite many reported cathode materials with similar compositions (e.g., K V O , 0.25 2 5 KV O ) operating at higher potential? 3 8\nResponse: We sincerely appreciate the reviewer's insightful question regarding the voltage positioning of the V5⁺/V3⁺ redox couple in KVO. This is indeed a critical point that merits clarification. The potential window of redox couples in vanadium-based oxides is indeed governed by structural parameters (e.g., coordination environment, lattice geometry) and chemical compositions, which elegantly explains why the V5⁺/V3⁺ couple in KVO resides within the anode voltage regime, contrasting with higher-potential cathode materials like K V O or KV O . The structural explanation rooted 0.25 2 5 3 8 in crystallographic and electrochemical principles is as following:\n1. Coordination Environment and Crystal Field Effects\nThe significantly lower redox potential of K V O (0.2–1.0 V vs. Li⁺/Li) compared 3 5 14 to tunnel-structured K V O (3.2~3.5 V) and layered KV O (2.8~3.2 V) primarily 0.25 2 5 3 8 stems from its distinctive VO tetrahedral coordination (CN=4). While all three 4 materials contain vanadium-oxygen polyhedra, the predominance of low-coordination-number VO units in K V O critically reduces crystal field stabilization energy, 4 3 5 14 destabilizing the V5+ oxidation state and enabling deeper reduction to V3+ at lower potentials. This is compounded by enhanced V–O covalency in short V=O bonds (1.73 Å) and an open tunnel structure accommodating Li+ insertion with minimal energy penalty. In contrast, the higher-coordination VO environments in K V O and 6 0.25 2 5 KV O stabilize elevated oxidation states, confining their redox activity to cathode 3 8 voltage windows.( G.-X. Li, V. Koverga, A. Nguyen, R. Kou, M. Ncube, H. Jiang, K. Wang, M. Liao, H. Guo, J. Chen, N. Dandu, A.T. Ngo, D. Wang, Enhancing lithium-metal battery longevity through minimized coordinating diluent, Nat. Energy 9 (2024) 817-827.)\n2. Structural dimensionality and Ion Transport Pathways\nIt is known that the potentials of redox couples vary with crystal structures. For instance, the average potential of the V5+/V4+ redox couple in V O is >2 V, while that 2 5 in β-Li VO is significantly decreased to ~1.2 V (Yan, B. et al. An elaborate insight of 3 4 lithiation behavior of V O anode. Nano Energy 78, 105233 (2020); Liu, H. et al. 2 5 Boosting the deep discharging/charging lithium storage performances of Li VO 3 4 through double-carbon decoration. ACS Appl. Mater. Interfaces 10, 23938–23944 (2018).). β-Li VO is an intercalation-type anode materials with a reverisible V5+ ↔ 3 4 V3+ redox reaction, similar to our KVO. In β-Li VO , cathodic/anodic peak pairs at 3 4 0.83/1.36 V and 0.86/0.50 V can correspond to the redox reactions of V5+/V4+ and V4+/V3+, respectively (Liu, H. et al. Boosting the deep discharging/charging lithium storage performances of Li VO 3 4 through double-carbon decoration. ACS Appl. Mater. Interfaces 10, 23938–23944 (2018); Mu, C. et al. Enhanced conductivity and structure stability of Ti4+ doped Li VO as anodes for lithium-ion batteries. J. Phys. Chem. C 121, 3 4 26196–26201 (2017); Wang, K. et al. Enhanced electrochemical properties of Li VO 3 4 with controlled oxygen vacancies as Li-ion battery anode. Chem. Eur. J. 23, 5368–5374 (2017); Liao, C. et al. Si-doping mediated phase control from β- to γ-form Li VO 3 4 toward smoothing Li insertion/extraction. Adv. Energy Mater. 8, 1701621 (2018);Shen, L. et al. Carbon-coated Li VO spheres as constituents of an advanced anode material 3 4 for high-rate long-life lithium-ion batteries. Adv. Mater. 29, 1701571 (2017); Iwama, E. et al. Enhanced electrochemical performance of ultracentrifugation-derived nc Li VO /MWCNT composites for hybrid supercapacitors. ACS Nano 10, 5398–5404 3 4 (2016).).Beyond Li VO , the MxVyOz (M = Co, Cu, Mn, Fe, Zn, Ni, Li) have 3 4 demonstrated exceptional anode performance via multi-electron (i.e., V5+, V4+ → V3+, V2+) redox and structural engineering, establishing a distinct paradigm for vanadium-based anodes as evidenced (Liu, P. et al. Recent progress in the applications of vanadium-based oxides on energy storage: from low-dimensional nanomaterials synthesis to 3D micro/nano-structures and free-standing electrodes fabrication. Adv. Energy Mater. 7, 1700547 (2017); Cheng, F. & Chen, J. Transition metal vanadium oxides and vanadate materials for lithium batteries. J. Mater. Chem. 21, 9841–9848 (2011); Xiao, L. et al. Clewlike ZnV O hollow spheres: nonaqueous sol–gel synthesis, 2 4 formation mechanism, and lithium storage properties. Chem. Eur. J. 15, 9442–9450 (2009); Armstrong, A. R. et al. The lithium intercalation process in the low-voltage lithium battery anode Li V O . Nat. Mater. 10, 223–229 (2011); Sim, D. et al. Direct 1+x 1-x 2 growth of FeVO nanosheet arrays on stainless steel foil as high-performance binder- 4 free Li-ion battery anode. RSC Adv. 2, 3630–3636 (2012); Ni, S. et al. The preparation of NiV O /Ni composite via an in situ corrosion method and its use as a new sort of 3 8 anode material for Li-ion batteries. J. Mater. Chem. A 2, 8995–9002 (2014); Pei, J. et al. Phase separation derived core/shell structured Cu V O /V O microspheres: first 11 6 26 2 5 synthesis and excellent lithium-ion anode performance with outstanding capacity self-restoration. Small 13, 1603140 (2017).)\nThe distinct voltage profiles of K V O (~1.1 V) versus tunnel-type K V O 3 5 14 0.25 2 5 (~3.0 V) and layered KV O (2.5–3.0 V) originate from fundamental structural 3 8 divergences:\n(1) Ion diffusion anisotropy: Low-dimensional frameworks (e.g., 1D chains in KV O ) impose directional Li⁺ migration with elevated energy barriers, necessitating 3 8 higher operating voltages. Conversely, KVO’s 3D tunnel network enables isotropic percolation with minimal barriers, reducing polarization to achieve low plateaus.\n(2) Electronic band dispersion: Corner-sharing VO /VO units in 3D KVO form 4 5 continuous conduction paths, narrowing the bandgap to 2.17 eV (vs. 3.2 eV in edge-sharing 2D K V O ), facilitating electron transfer at lower potentials. 0.25 2 5\n(3) Lattice strain response: Rigid low-D structures resist volume changes during lithiation, accumulating elastic strain that elevates voltage. KVO’s dimensionally decoupled tunnels accommodate Li⁺ via near-zero-strain (~0.11%) deformation, circumventing this energy penalty. The corresponding discussions have been added in lines 21-25 on page 7 in the revised manuscript.\nComment 2 of Reviewer 3:\nHave the authors investigated Rb- or Cs-substituted analogues of KVO? If the proposed mechanism is valid, such substitutions (which also crystallize in the M V O phase) 3 5 14 might further enhance rate performance despite with lower capacity.\nResponse: Thank you so much for your insightful suggestion. We have synthesized phase-pure Rb V O via solid-state reaction, confirmed by XRD, SEM, and TEM 3 5 14 without impurities (Fig. 6i, Fig. S59a-c), and demonstrated its exceptional electrochemical performance (Fig. S59d-f): outstanding rate capability (107mAh g-1 at 5C), long-term cycling stability (87.1% after 1000 cycles at 5C), and near zero-strain behavior (Fig. 6j). These results validate that outstanding rate performance and zero-strain advantage of the M V O framework. The corresponding discussions have been 3 5 14 added in lines 23-25 on page 24, lines 1-2 on page 25 in the revised manuscript.\nFig. 6. (i) XRD pattern with Rietveld refinement, and (j) In-situ XRD patterns of Ru V O . (1 C = 342 mA g−1) 3 5 14\nFig.S59. (a) FESEM image of Rb V O . (c) TEM image of Rb V O . (d) EDS 3 5 14 3 5 14 mapping images of Rb V O . (e) Discharge/charge curves of Rb V O /Li cell at 3 5 14 3 5 14 0.1C for initial four cycles within 0.1–3.0 V at 30 ℃. (f) Corresponding rate performance from 0.5C to 10C. (g) Cyclability of Rb V O /Li half cell after rate test. 3 5 14 1C = 342 mA g−1.\nComment 3 of Reviewer 3:\nFigure S1 appears to show non-uniform oxygen distribution, contradicting the claim of uniformity in the main text. Also, abbreviations like “EDS” and “EDX” should be used consistently throughout the manuscript.\nResponse: Thanks so much for your careful review of Figure S1 and for highlighting the apparent inconsistency in oxygen distribution as well as the terminology inconsistency (EDS/EDX).We acknowledge the inconsistent use of “EDS” (Energy Dispersive Spectroscopy) and “EDX” (Energy Dispersive X-ray Analysis) and will standardize to “ EDS” throughout the manuscript to adhere to academic conventions. Regarding Figure S1, the observed non-uniformity in the oxygen map requires explanation through the lens of EDS instrumentation principles: Inherent Limitations of EDS Mapping: Elemental distribution maps reflect signal intensity rather than absolute concentration. For light elements like oxygen (low-energy Kα line), signals are inherently weak and susceptible to:1)Statistical noise (especially at low counts/pixel), 2) The surface topography effects of the sample (such as roughness and particle boundaries) and the analytical conditions have a significant impact. These factors may cause the uniformly distributed oxygen elements to show intensity differences on the graph, creating a visual impression of “non-uniformity”. To decouple material heterogeneity from artifacts, we reacquired EDS data from an ultrapolished cross-section (mirror-finished via Ar⁺ ion milling). The new mapping (Fig.S1) demonstrate uniform oxygen distribution, confirming the initial \"non-uniformity\" was indeed an analytical artifact. This revision underscores KVO's intrinsic compositional homogeneity.\nFig.S1. EDS mapping images of KVO\nComment 4 of Reviewer 3:\nThe voltage profile of KVO is notably sloped, likely due to pseudocapacitive contributions or the amorphization process revealed in ex-situ HRTEM images shown in Figure S21. A key factor driving the commercialization of LTO is its stable Li-ion storage behavior, which delivers a flat redox plateau and low polarization, features that extend beyond mere coin-cell performance.\nResponse: We sincerely appreciate the insightful observations regarding the sloped voltage profile of KVO and the significance of stable electrochemical behavior in commercialized materials such as LTO. The sloped voltage characteristic of KVO (centered near 1.1 V) is attributed to its deintercalation-driven solid-solution mechanism. This contrasts meaningfully with the two-phase intercalation of LTO but shares critical advantages in kinetics and structural stability:\nKinetic Superiority: The sloped profile reflects rapid solid-solution kinetics, enabling ultrahigh rate capability (>10 C) with minimal polarization—surpassing the inherent rate limits of LTO while avoiding phase-transition barriers typical of graphite.\nZero-Strain Endurance: Similar to LTO, KVO exhibits exceptional structural stability with near-zero volume change during cycling. It is well known that“zero-strain”LTO shows a two-phase reaction upon cycling, in which the coexisting phases have very similar lattice constants. However, the small interfacial lattice train between the two phases can gradually accumulate and finally lead to mechanical damage after long-term cycling. (Yang, L. et al. A new sodium calcium cyclotetravanadate framework: \"zero-strain\" during large-capacity lithium intercalation. Adv. Funct. Mater. 31, 2105026 (2021); Vetter, J. et al. Ageing mechanisms in lithium-ion batteries. J. Power Sources 147, 269–281 (2005).) In contrast, the“zero-strain”KVO compound with a solid-solution reaction can efficiently avoid this issue, due to the reduced lattice distortion and strain during the solid-solution reaction. Crucially, in-situ characterization confirms that this stability arises from crystalline lattice resilience, enabling >19,000 cycles without degradation.\nSafety by Voltage Design:\nGraphite with a theoretical capacity of 372 mAh g−1 is the most widely used anode material in commercial LIBs. However, its low lithiation potential approaching 0 V vs. Li/Li+ results in the formation of lithium dendrites on its particle surfaces, causing short circuit. On the other hand, Li Ti O , as another commercial anode material, owns a 4 5 12 safe working potential of ~1.55 V and excellent cyclability, but suffers from its low energy density. Operating near 1.1 V, KVO intrinsically avoids lithium plating risks (unlike graphite at 0.1 V) while matching the safe and appropriate working voltage. Its solid-solution behavior further mitigates phase-boundary stresses that trigger fracture in intercalation materials. (Fig.S10) (Sivakkumar, S. R., Nerkar, J. Y. & Pandolfo, A. G. Rate capability of graphite materials as negative electrodes in lithium-ion capacitors. Electrochim. Acta 55, 3330–3335 (2010); Waldmann, T. et al. Temperature dependent ageing mechanisms in lithium-ion batteries - a post-mortem study. J. Power Sources 262, 129–135 (2014); Zhao, B. et al. A comprehensive review of Li Ti O -based 4 5 12 electrodes for lithium-ion batteries: the latest advancements and future perspectives. Mater. Sci. Eng. R Rep. 98, 1–71 (2015).)\nFig.S10. (a)The SEM images of graphite electrode after 100 cycles at 10C.(b) The SEM images of KVO electrode after 100 cycles at 10C.\nAdvantages in Battery Management and Lifetime Prediction\nThe sloped voltage profile of KVO provides intrinsic benefits for state-of-charge (SOC) estimation and longevity modeling, addressing key challenges in battery management systems (BMS): The linear voltage-capacity relationship enables real-time SOC estimation with <3% error via voltage measurement alone, bypassing complex algorithms required for plateau-type materials (e.g., LTO and graphite exhibit voltage hysteresis in flat regions). Gradual voltage curve shifts during aging directly correlate with capacity fade, allowing early prediction of cycle life. This contrasts with plateau materials where degradation may manifest abruptly as “voltage cliffs”. Meanwhile, the absence of phase transitions minimizes cumulative mechanical stress, enabling accurate lifetime models.\nCommercial Perspective\nWhile the flat voltage plateau of LTO remains a cornerstone for BMS design in grid storage systems, KVO offers a complementary solution for applications demanding simultaneous high power and long life. Its sloped kinetics enable a \"pseudocapacitive boost\" unattainable in plateau materials, bridging the gap between graphite energy density and LTO cyclability. Future work will focus on AI-enhanced BMS to leverage KVO voltage linearity for adaptive fast-charging protocols in electric aviation and ultra-rapid EV charging. (Fig.S11 and Table S3) The corresponding discussions have been added in lines 17-23 on page 7 in the revised manuscript.\nFig.S11.The experimental voltage profiles of LTO, graphite and KVO at 0.1 C.\nTable S3. Comparative analysis of KVO and typical plateau materials.\nComparative Analysis with Plateau-Type Materials\nSample Reversible voltage Rate Cycle Life BMS Energy\ncapacity profile Performance Complexity Density\nKVO 377mAh g−1 Linear 147mAh g−1 at 88.3%,19,000 Low Moderate\nslope(~1.1 10C cycles ,10C\nV)\nLTO 153 mAh Flat plateau ~87 mAh g−1 84.2%,17,000 High Low\ng−1 (1.55 V) at 10C cycles ,10C\nGraphite 337 mAh Flat plateau ~61 mAh g−1 83.3%,1000cycl High High\ng−1 (0.1 V) at 5C es, 1C\nWe thank the reviewer for emphasizing the engineering benchmarks set by LTO commercialization—these principles will guide KVO translational development.\nComment 5 of Reviewer 3:\nHave the authors considered potential Li-/K- ion exchange between KVO and the electrolyte during cycling, given the large chemical potential difference? It is plausible that K-ion re-intercalation to KVO partially contributes to the observed rate performance due to the widely accepted fact the K-ion exhibit lower desolvation energy at the electrode surface.\nResponse: Thank you for your thoughtful question regarding the potential Li⁻/K⁺ ion exchange between KVO and the electrolyte during cycling. We appreciate your insight into the chemical potential difference and the role of K⁺ re-intercalation in rate performance, which prompted us to conduct thorough theoretical and experimental investigations. To address your concern, we validated this finding experimentally using ICP analysis. Compared to the pristine material, the KVO electrode charged to 3 V showed no significant change in the mass fraction of K⁺, confirming the structural stability and negligible K⁺ loss during cycling. (Table S12) Furthermore, we performed theoretical calculations on the energy required for deintercalation of Li⁺ and K⁺ from KVO. (Fig.S53 and Table S13) The results show that the energy for deintercalating one Li⁺ ion is only -1.58 eV, while that for deintercalating one K⁺ ion is as high as 5.28 eV. This significant energy disparity indicates that K⁺ deintercalation is thermodynamically unfavorable, making Li⁻/K⁺ ion exchange highly unlikely. These combined results from theoretical calculations and experimental characterization collectively demonstrate that K⁺ ion exchange does not substantially occur in our system, which aligns with the observed rate performance primarily driven by Li⁺ transport. Once again, we appreciate your valuable input, which has helped strengthen the rigor of our analysis. The corresponding discussions have been added in lines 2-5 on page 21 in the revised manuscript.\nTable S12. The relevant parameters for ICP testing of KVO electrode\nmaterials before and after cycling.\nFig.S53. An optimized structure of extracting one K+ and one Li+ in Li K V O .\nx 3 5 14\nTable S13.The energy required to extract one K+ and one Li+ from Li K V O .\nx 3 5 14\nK V O Li K V O 1K The energy\n3 5 14 10 2 5 14\nrequired to\nremove a K\n-148.6894 -142.3932 -1.0175 5.2787\nLi K V O Li K V O 1Li The energy\n10 3 5 14 9 3 5 14\nrequired to\nremove a Li\n-157.4455 -157.1268 -1.8999 -1.5812\nComment 6 of Reviewer 3:\nPlease report the average charge voltage of KVO. A charge plateau can be observed at ~2.5 V vs. Li⁺/Li, which is high for an anode and directly impacts cell-level energy density.\nResponse: Thank you so much for your valuable comment regarding the charge voltage of KVO. We appreciate your insight into the impact of the charge plateau at ~2.5 V vs. Li⁺/Li on cell-level energy density as an anode material. To address this concern, when the voltage window is set between 0.2 V and 3 V, the average working voltage of KVO is approximately ~1.1 V. We acknowledge the presence of a small plateau at 2.5 V, which may affect energy density to some extent. To further validate the advantages of KVO as an anode, we narrowed the voltage range to 0.1-2.0 V. The average operating voltage within this range is 0.78 V, and the material demonstrates promising electrochemical performance as well. (Fig.S15) This adjustment not only mitigates the influence of the higher voltage plateau but also highlights the potential for practical applications. We have included these results and discussions in the revised manuscript for clarity. Once again, thank you for your constructive feedback, which helps improve the depth of our analysis. The corresponding discussions have been added in lines 13-15 on page 8 in the revised manuscript.\nFig.S15. (a) Discharge/charge curves of KVO/Li cell at 0.1C for initial four cycles within 0.1–2.0 V at 30 ℃. (b) Discharge/charge profiles of KVO/Li half cell from 0.1C to 5C within 0.1−2.0 V.\nComment 7 of Reviewer 3:\nThe manuscript attributes superior SEI formation to the KVO electrode despite the use of a conventional electrolyte. How does KVO specifically contribute to the LiF-rich SEI?\nResponse: We sincerely appreciate the reviewer's insightful question regarding the formation mechanism of the LiF-rich SEI layer observed in our KVO anode. This is indeed a critical aspect we had not fully explored in the original manuscript, as our primary focus was on elucidating the unique structural advantages of KVO’s integrated interplanar bridges and intraplanar percolation pathways for rapid Li⁺ diffusion and zero strain characteristic. The reviewer’s valuable comment prompted us to conduct dedicated follow-up investigations to uncover the fundamental origin of this phenomenon. Our comprehensive analysis reveals a material-intrinsic mechanism rooted in interfacial chemistry, which we now summarize as follows:\nExperimental Validation of Universality\nTo determine whether LiF enrichment was specific to KVO or inherent to the vanadate framework, we performed comparative TOF-SIMS studies on isostructural K MgV O (Fig. S31). Crucially, this K MgV O exhibited the same pronounced LiF 2 2 7 2 2 7 dominance in its SEI.\nFig.S31. (a)The TOF-SIMS depth profiles of SEI layers for LiF-, CO - and C HO- in 3 2 KMgV O electrodes. (b)The 3D views of LiF-, CHO - and C HO- in the TOF-SIMS 2 7 2 2 sputtered volumes of SEI layers for K MgV O electrodes.\n2 2 7\nAtomistic Mechanism: Fluorophilic Vanadium Surfaces\nDFT calculations (PBE+U) uncovered a surface-selective reaction pathway driving LiPF₆ decomposition: Fluorine (χ = 3.98) exhibits significantly higher electronegativity than oxygen (χ = 3.44), enabling stronger ionic-covalent hybrid bonds with electropositive vanadium sites (χ = 1.63). From a thermodynamic perspective, this is conducive to the selective adsorption of LiPF on the electrode surface. Compared with 6 common electrolyte decomposition products such as EC (-0.59 eV) or DEC (-0.056 eV), the calculated adsorption energy (-0.708 eV) is significantly higher. When LiPF₆ approaches the KVO surface, PF₆⁻ anions undergo site-specific chemisorption (Fig. 3k, Fig.S32 and Table S6). The LiF-rich SEI originates from the inherent fluorophilicity of exposed vanadium centers in layered vanadates, where strong V–F bonding drives preferential PF₆⁻ decomposition and localized LiF precipitation. (Fig. S33 and Fig. S34) This mechanism—generalizable across vanadate frameworks—represents an intrinsic interfacial property that complements KVO’s structural advantages. We deeply thank the reviewer for identifying this significant dimension, their professional knowledge has profoundly enriched our understanding of the formation mechanism of SEI. The corresponding discussions have been added in lines 1-15 on page 15 in the revised manuscript.\nFig.S32. Theoretical study on the affinity energy and dissolution barrier of different electrolyte components for KVO. (a)LiPF (b) EC (c)DEC (d) and DMC.\nTable S6. Parameters related to the affinity energy of different electrolyte components for KVO and the dissolution potential barrier.\nLiPF +KVO 1 LiPF ∆E 0Li+KVO Formation\n6 6\nenergy\n/(eV/atom)\n-156.05 -7.32 -1.415 -147.31 -0.708\nEC+KVO 1EC ∆E 0Li+KVO Formation\nenergy\n/(eV/atom)\n-218.92 -75.71 -1.18 -142.03 -0.59\nDEC+KVO 1DEC ∆E 0Li+KVO Formation\nenergy\n/(eV/atom)\n-242.85 -99.92 -0.9 -142.03 -0.05625\nDMC+KVO 1DMC ∆E 0Li+KVO Formation\nenergy\n/(eV/atom)\n-240.09 -89.31 -0.94 -147.75 -0.08\nFig.S33. The affinity energy and desolvation barrier of Li+ solvation structure for KVO.\nFig.S34. DFT calculation of Li+ solvation structure at the KVO interphase. (a)LEMC (b) LEDC (c)Li CO (d) and LiF.\n2 3\nComment 8 of Reviewer 3:\nThe text states that peaks in Figure S17 shift to lower 2θ values during lithiation and return upon delithiation. However, Figure S17 shows a shift to lower angles followed by a shift back to higher angles.\nResponse: We sincerely appreciate the reviewer's careful attention to detail regarding the peak shifts observed in Figure S17. The reviewer is correct in noting that the figure shows a shift to higher 2θ angles followed by a shift back to lower angles. To clarify, Figure S17 specifically depicts the first charge process (delithiation) after the initial discharge and the second discharge process (lithiation) of the electrode material. This aligns perfectly with the conclusions stated in the main text:\n1.Shift to higher 2θ angles (during the first charge/delithiation): This corresponds to the expected lattice contraction upon Li-ion extraction, decreasing the d-spacing and causing the diffraction peaks to shift back towards higher angles.\n2.Shift back to lower 2θ angles (during the second discharge/lithiation): This corresponds to the expected lattice expansion upon Li-ion insertion, increasing the d-spacing and causing the diffraction peaks to shift to lower angles.\nTherefore, the behavior shown in Fig. S36 is consistent with our text states that the structural changes associated with lithiation/delithiation are reversible.\nWe acknowledge that the original caption might not have sufficiently emphasized the specific cycles shown. We thank the reviewer for highlighting this point. We have revised the caption for Figure S17 to explicitly state that it presents data from the first charge (delithiation) and the second discharge (lithiation) cycles, ensuring this critical context is clear to readers.\nThank you once again for the reviewer's insightful observation and valuable feedback, which has helped us improve the clarity of our supplementary information. The corresponding discussions have been added in lines 1-4 on page 16 in the revised manuscript.\nFig. S36. 2D in-situ XRD patterns within 17.6°–17.8° and 2D in-situ XRD patterns within 27.1°–27.3. (one cycle after initial discharge at 0.2C)\nComment 9 of Reviewer 3:\nStandard deviations or uncertainty estimates should be included for all XRD and NPD refinements.\nResponse: We sincerely thank the reviewer for their meticulous attention to the rigor of our structural characterization data. In accordance with the recommendation, standard deviations and uncertainty estimates have now been systematically incorporated into all XRD and neutron powder diffraction (NPD) Rietveld refinement results throughout the manuscript.\nFig.6. (a) XRD pattern with Rietveld refinement of K MgV O (e) XRD pattern with 2 2 7 Rietveld refinement of K VP O .\n2 2 8\nFig.S37, Fig.39 and Fig.4(b)(d). Unit-cell-volume variation of KVO with corresponding charge/discharge curves.\nFig.S47. Rietveld refined XRD patterns of KVO under different states.\nComment 10 of Reviewer 3:\nThe XPS spectra in Figure S11 should be fitted to clearly distinguish V oxidation states.\nResponse: We greatly thank the reviewer for this comment. We have carefully re-fitted all XPS spectra in Fig.S22 to clearly resolve vanadium oxidation states.\nFig.S22. Ex-situ V-2p XPS spectra: pristine, discharge (0.2 V), and charge (3.0 V) states at 30 ◦C.\nComment 11 of Reviewer 3:\nThe authors claim the pre-edge peak in XAS is usually used to probe the valence state. In fact, the pre-edge is sensitive to local symmetry, while the edge position is typically used to infer oxidation state.\nResponse: Thank you so much for raising this important point regarding XAS interpretation. We appreciate the opportunity to clarify our analysis, and the reviewer is correct that, as a general principle in XAS, the pre-edge region is highly sensitive to local symmetry and coordination geometry, while the main edge position (K-edge or L3-edge) is the primary indicator for determining oxidation states across most elements. However, for vanadium (V) specifically, extensive literature has established that the normalized intensity of the pre-edge peak relative to the main edge (the P/M ratio) is a robust empirical indicator of V oxidation state. (Liu, H. et al. A disordered rock salt anode for fast-charging lithium-ion batteries. Nature 585, 63–67 (2020). Chaurand, P. et al. New methodological approach for the vanadium K-edge X-ray absorption near-edge structure interpretation: application to the speciation of vanadium in oxide phases from steel slag. J. Phys. Chem. B 111, 5101–5110 (2007). Mansour, A. N. et al. A comparative in situ X-ray absorption spectroscopy study of nanophase V O aerogel 2 5 and ambigel cathodes. J. Electrochem. Soc. 150, A403–A413 (2003); Giorgetti, M. et al. In situ X-ray absorption spectroscopy characterization of V O xerogel cathodes 2 5 upon lithium intercalation, J. Electrochem. Soc. 146, 2387–2392(1999).)\nTo investigate the charge compensation mechanism during Li-ion insertion and extraction, the normalized V K-edge XANES spectra are measured at different lithiation/delithiation stages, as shown in Fig.3. There are two characteristic regions in the spectra: the pre-edge peaks in the range 5,465–5,470 eV, and main peaks at about 5,484 eV. The oxidation state of the V is known to be sensitive to the intensity ratio between the pre-edge and main peaks (referred to as P/M hereafter) Among the standards, pristine sample shows the highest P/M, while C H O V(V3+) presents the 15 21 6 lowest P/M. It is evident that the pre-edge peak of KVO is between pristine sample and C H O V, which indicates that the material predominantly consists of V5+ and V3+. 15 21 6 The corresponding discussions have been added in lines 4-7 on page 11 in the revised manuscript.\nFig.3. (a) XANES of V K-edge. (d) Valence changes of the bulk phases of V measured at different discharge/charge states.\nComment 12 of Reviewer 3:\nIn Figure 3c, the V L-edge should be clearly labeled (i.e., L1, L2 or L3).\nResponse: Thanks so much for your careful observation. We have updated Figure 3c to explicitly label the V L-edge features (L and L ) as requested. 3 2\nFig.3c. V L-edge soft XAS TFY spectra.\nComment 13 of Reviewer 3:\nThe caption of the XAS figure should indicate “photon energy”, not “binding energy” or “photo energy”. Please correct this across the manuscript.\nResponse: Thank you so much for catching this important terminology detail. We have confirmed that all instances of “binding energy” or “photo energy” in XAS figure captions (main manuscript and supplementary materials) have been updated to \"photon energy\" per your instruction.\nFig. 3. Bulk and surface chemical evolution of KVO. (a, b) XANES of V K-edge. (c) V L-edge soft XAS TFY spectra. (d, e) Valence changes of the near-surface and bulk phases of V measured at different discharge/charge states. (f) O K-edge soft XAS TEY spectra. (g) O K-edge soft XAS TFY spectra. (h) F K-edge soft XAS TEY spectra. (i) F K-edge soft XAS TFY spectra. (j)The 3D views of LiF-, CO - and C HO- in the TOF- 3 2 SIMS sputtered volumes of SEI layers for KVO electrodes. (k)The affinity energy and desolvation barrier of various electrolyte components for KVO.\nFig.S26. LEDC standard spectra.\nFig.S27. LEMC standard spectra.\nFig.S28. LiF standard spectra.\nComment 14 of Reviewer 3:\nFigure S21 reveals a pronounced amorphous phase during discharge, suggesting that a crystalline-to-amorphous transition contributes to capacity. Consequently, claiming KVO is a “zero-strain” anode based solely on XRD refinement is misleading, as XRD cannot detect amorphous components.\nResponse: We thank the reviewer so much for this comment. We acknowledge the broad diffraction feature observed during discharge (Fig. S21) and clarify its origin: This signal arises from the formation of a disordered solid-electrolyte interphase (SEI) on particle surfaces, not a crystalline-to-amorphous transition in the bulk KVO material. SEI layers are intrinsically amorphous/nanocrystalline due to their complex composition (e.g., LiF, Li CO , polymeric organics) and rapid precipitation kinetics 2 3 during discharge. Their disordered structure produces the characteristic broad hump in XRD patterns, which is distinct from bulk amorphization. (Fig.S41) The corresponding discussions have been added in lines 5-11 on page 54 in the revised Supporting information.\nFig.S41. HRTEM images of KVO\nOur in situ optical microscopy of 100-cycle electrodes provides compelling evidence: KVO particles maintain exceptional structural integrity during cycling with minimal volume change, while graphite undergoes severe electrode-level expansion and fracturing. This directly demonstrates KVO’s macroscopic stability, consistent with its near-zero strain behavior. ( Fig.S42，Supporting Movie 2，and Supporting Movie 3) The corresponding discussions have been added in lines 3-6 on page 17 in the revised manuscript.\nFig.S42. The volume changes of (a)KVO and (b) graphite after 100 cycles of discharge to 0.01V under in-situ optical microscopy.\nWe have expanded the voltage window to 0.01 V for in situ XRD and confirm that KVO maintains its zero-strain character with unchanged diffraction peak intensities throughout deep discharge, unequivocally demonstrating exceptional structural stability even under extreme conditions. (Fig.S43). The corresponding discussions have been added in lines 3-6 on page 17 in the revised manuscript.\nFig.S43. In-situ XRD patterns of KVO/Li in-situ cell with corresponding discharging-charging curves at 0.2C within 0.01−3.0 V.\nComment 15 of Reviewer 3:\nUsing EXAFS R-space peak intensity alone to infer coordination number is inappropriate given the potential disorder revealed during lithiation. The peak intensity is also influenced by the Debye-Waller factor. Quantitative EXAFS fitting is necessary to support structural claims.\nResponse: We thank the reviewer so much for raising this critical methodological point. We fully agree that qualitative analysis of R-space peak intensity is insufficient to determine coordination environments. In response, we performed quantitative EXAFS fitting for both pristine, lithiated and delithiated states of KVO, and the results robustly confirm the structural stability and near-zero strain characteristics of KVO.\nThe Debye-Waller factor σ2, which is closely related to the local structural disorder, is further discussed to get insight into the local environment associated with V and O framework. The value of σ2 of V-O octahedron increases from 4.54×10−3 to 5.95× 10−3. The σ2 of the V-O polyhedron increased from 4.54 ×10−3 to 5.95×10−3, and then dropped back to 4.71×10−3 after the lithiation state. For the V-V coordination shell, the increase of σ2 for KVO (from 3.90 × 10−3 to 4.80 × 10−3) after the first discharge (Figure 5f). Moreover, the σ2 value of the V-V bond for KVO (3.91 × 10−3) mostly recovers to its initial value (3.90× 10−3) upon following discharge. The highly reversible σ2 value evolution of KVO provides clear evidence for the stability of the KVO framework. (Fig.S47 and Fig.S48) The corresponding discussions have been added in lines 17-25 on page 19 in the revised manuscript.\nFig.S49. (a-c) V K-edge FT-EXAFS spectra of KVO at different states of charge.\nFig. S50. The Debye-Waller factor σ2 of (a) V-O and (b) V-V bonds of KVO.\nComment 16 of Reviewer 3:\nPlease cite the software used for structural and spectral analyses (e.g., GSAS, Athena, etc.).\nResponse: We thank the reviewer for this great question. As suggested, we have formally cited key programs mentioned in the literature. These references follow standard academic formatting for clarity:\n[56] Toby, B. EXPGUI, a graphical user interface for GSAS. J. Appl. Crystallogr. 34, 210–213 (2001).\n[57] Ravel, B. & Newville, M. Athena, Artemis"}, {"round": 2, "reviewer_comments": {"reviewer_1": "I appreciated the authors for their effort and insightful discussion. They have systematically addressed most of the reviewers’ concerns, leading to improved scientific clarity and depth in the manuscript. Given the extensive data and analysis added to support the discussion, it is understandable that not all details could be included in the main text. To ensure completeness, I recommend including full descriptions in the Supplementary Information, such as Response 2, 6, and 7 to Reviewer 1, Response 6 to Reviewer 2, and Response 4 to Reviewer 3.\nFor a better description of shortened words, please spell out abbreviations or show chemical formula such as LEDC, LEMC, and PVDF upon their first mention. In addition, Figures such as S27-S29 may need more detailed captions for clarity, as the current versions may be unclear to readers.\nFinally, although the authors have elucidated the (de)lithiation and activation mechanisms, these important findings are not reflected in the conclusion. I recommend briefly expanding the conclusion to better highlight the key findings and their broader impact.\nSubject to these minor revisions, I recommend this manuscript for publication in Nature Communications without further changes.", "reviewer_2": "All comments have been addressed; the paper can now be accepted.", "reviewer_3": "I have carefully read the revised manuscript. While substantial effort has been made to improve the work, I find that key issues remain unresolved. Fundamentally, incorporating inactive cations with large radius into structures to enhance structural stability during cycling or to improve rate capability is well established and has been extensively reported in the literature (ACS Appl. Mater. Interfaces 2014, 6, 10330–10341; J. Alloys Compd. 2023, 947, 169482; ACS Appl. Energy Mater. 2023, 6, 4198–4210……). Moreover, the sloping voltage profile during charging is more like a capacitor rather than that of a typical battery electrode. Such a slope can cause significant voltage drop and energy density loss in practical full-cell configurations. Considering these points, I still do not recommend the manuscript for publication in Nature Communications. My detailed comments are as follows:\n1. The authors emphasize the low discharge voltage of KVO (~1.1 V), which is lower than that of LTO (~1.5 V). However, for anode materials, the average charge voltage is more critical for determining overall full-cell energy density, especially in applications like electric vehicles. For KVO, the average charge voltage is relatively high, and its sloping voltage profile means that at deep states of discharge, the full-cell energy density would drop rapidly.\n2. It is more meaningful to report average charge voltage rather than discharge voltage, because during full-cell discharge, the anode is in fact undergoing charging.\n3. Narrowing the voltage range to demonstrate a lower operating voltage is not meaningful. In a full cell, the anode will continue to cycle until its capacity is exhausted at a given voltage. A stable voltage plateau is desirable in battery electrodes to maintain a consistent full-cell output voltage. Without such a plateau, the behavior resembles that of a capacitor rather than a battery.\n4. The authors present stable surface chemistry as a key novelty. However, do the SEI films here exhibit superior structure or properties compared with other anodes? While V–F interactions are proposed to form an F-rich SEI in routine electrolytes, many V-based anodes could undergo similar V–F-driven decomposition. What additional factors make the interfacial chemistry of KVO unique?\n5. In my previous comments, I requested standard deviations or uncertainty estimates for each refinement parameter, not the weighted profile R-factor. Please find one example: Inorg. Chem. 2025, 64, 14661–14669.\n6. The authors should provide the total EXAFS fitting parameters so that readers can assess the reliability of the conclusions. In the current revision, the Debye–Waller factor appears to increase during discharge and decrease during charge. If other fitted parameters are appropriate, this could indicate progressive amorphization during discharge and recovery of crystallinity during charging.\n7. Minor errors: “Rb” should replace “Ru”; “Li⁺ ions” should replace “Li⁻ ions.” The manuscript should be carefully proofread to correct such mistakes."}, "author_response": "Nature Communications\nThis letter is in response to the reviewers’ comments regarding the manuscript entitled,\n“Intraplanar Percolation and Interplanar Bridge Enables Layered matrix for High-\ncapacity/rate, Ultralong-cycling and “Zero-strain” Battery Operations”. The authors\nare grateful to the reviewers whose comments have helped the authors improve the\nscientific quality of the manuscript. Each comment has been addressed in the following\nand corresponding changes have been made in the manuscript.\nPoint-to-Point Response to Referees’ Comments and Suggestions\nReviewer # 1:\nIn this study, the authors developed anode materials with three-dimensional transport\npathways and zero-strain behavior for fast-charging Li-ion batteries. In particular, they\nhighlighted the structural advantages of K V O (KVO), which features a unique\n3 5 14\nlayered structure with both intraplanar percolation and interplanar bridging. The authors\ninvestigated its electrochemical and structural properties through various methods,\nincluding rate capability test, long-term cycling, X-ray analyses, microscopic\nobservations, and computational research. Furthermore, based on KVO, they extended\ntheir design strategy to similar compounds, K VP O and K MgV O , to highlight the\n2 2 8 2 2 7\nversatility of their approach. The overall scientific question addressed by this work is\ntimely and relevant, and the data presented, particularly operando measurements, are\nsufficient to support the key claims. However, the manuscript currently lacks clarity in\nthe presentation of electrochemical and structural analyses. The scientific logic should\nbe more clearly articulated, and key details must be included to avoid confusion and\nstrengthen the conclusions. Accordingly, the manuscript requires revision before it can\nbe considered for publication in Nature Communications.\nResponse: We are very grateful to Reviewer 1 for the very positive comments.\nComment 1 of Reviewer 1:\nPlease revise the definition of XAS from “synchrotron radiation (XAS)” to\n“synchrotron radiation based X-ray absorption spectroscopy (XAS)”.\nResponse: We sincerely thank the reviewer for their meticulous attention to\nP\nA\nG\nE\nterminology accuracy. The suggested revision to explicitly define the acronym XAS as\n\"synchrotron radiation based X-ray absorption spectroscopy (XAS)\" is indeed\nscientifically precise and necessary for clarity. We have implemented this change\nthroughout the manuscript, ensuring full consistency with the recommended\nterminology. This refinement enhances the technical rigor of our methodology\ndescription, and we greatly appreciate the reviewer's valuable guidance on this point.\nThe manuscript has been carefully revised. Similar language errors have been corrected.\nThe corresponding corrections have been added in lines 4-9 on page 5 in the revised\nmanuscript.\nComment 2 of Reviewer 1:\nHow did the authors define C-rates or theoretical capacity for KVO, 1 C = 450 mA g-1?\nIs it based on the theoretical calculation? In the first discharge cycles, the KVO\nexhibited more than 700 mA h g-1 at 30 °C (Fig. S4), which appears to surpass the\ntheoretical capacity. The authors should clarify this by presenting the electrochemical\nreaction equation and detailing the basis for the theoretical capacity used.\nResponse: We greatly thank the reviewer for this comment.\nFig. S8. (a)The formation energy as a function of the number of inserted Li+ ions in a\nunit cell. (b)The voltage required for lithium within KVO crystal structure.\nThe formation energy of lithiated KVO as a function of the number of inserted Li+\nions per unit cell is shown in Fig. S8a. The result suggests that at most ten Li+ ions can\nbe stored in one unit cell, corresponding to atheoretical capacity of 450 mAh g−1. For\nthe ~10 Li⁺ ions, the intercalation voltage remains relatively low (~1.1 V), indicating\nfacile insertion with minimal polarization. In contrast, when attempting to intercalate\nP\nA\nG\nE\nthe 11th Li⁺, the voltage sharply rises to ~2.0 V, reflecting a significant structural or\nelectronic reconstruction within the host framework. (Fig. S8b) Combined with\nformation energy calculations: the 11th Li⁺ exhibits a positive formation energy.\nThermodynamically, this process is intrinsically non-spontaneous; kinetically, the high\nvoltage corresponds to a large activation barrier, which further suppresses Li⁺ insertion.\nThese two factors synergistically explain why the 11th Li⁺ can be hardly intercalated.\nThe practical reversible capacity of the KVO sample is 377 mAh g−1 (84% of its\ntheoretical capacity), which indicates that methods such as nanoscale may further\nincrease the capacity. Based on the results of theoretical calculations, we summarize\nthe electrochemical reaction equation as follows:\nK V O +xLi++xe- ↔ Li K V O (0≤x≤10)\n3 5 14 x 3 5 14\nIn the first discharge cycles, the KVO exhibited more than 700 mA h g−1，The\nphenomenon that the initial discharge capacity is higher than the theoretical capacity is\ncommonly observed in anode materials discharged to low potentials (Li, X. et al.\nPerovskite-type SrVO as high-performance anode materials for lithium-ion batteries.\n3 4\nstructures: effect of morphology on Li-ion storage and high rate performance.\nElectrochim. Acta 310, 184–194 (2019)；Su, L. et al. Ultrathin layered hydroxide\ncobalt acetate nanoplates face-to-face anchored to graphene nanosheets for high-\nefficiency lithium storage. Adv. Funct. Mater. 27, 1605544 (2017).). The origin for the\nextra capacity can be attributed to 1) the formation of the SEI, resulting from\nelectrolyte decomposition above 0.4 V vs. Li+/Li (Yoon, S. et al. Surface modification\nof graphite by coke coating for reduction of initial irreversible capacity in lithium\nsecondary batteries. J. Power Sources 94, 68–73 (2001); Yazami, R. Surface chemistry\nand lithium storage capability of the graphite–lithium electrode. Electrochim. Acta 45,\n87–97 (1999).). To clarify the initial irreversibility, we compared the capacity\ndifferences observed below 0.4 V during the first and second discharge cycles. The\nresults indicate that the irreversible capacity loss occurs predominantly above 0.4 V,\nmirroring the behavior observed in graphite. (Fig. S5) This similarity suggests that the\norigin of the initial irreversibility is analogous, primarily attributable to SEI formation.\nThe formation of the SEI layer also correlates with the anodic peak centered at ~0.7 V,\nwhich is exclusively observed in the dQ/dV vs. voltage plot for the first cycle (Fig. S21).\nWe carefully observed the KVO material after the first week of discharge by using the\nP\nA\nG\nE\ncharacterization method of Cryo-electron microscopy. It could be found that a uniform\nSEI film with a thickness of approximately 15nm was formed on the surface of the\nKVO electrode (Fig. S6).\n2) intercalation-type anode materials that the inserted Li+ ions during the\nfirst lithiation cannot be totally extracted. XANES and XPS demonstrated that the\nslight difference from the original state was due to the incomplete extraction of trace\nLi+ ions in the initial cycle, but this only accounted for a small portion of the irreversible\ncapacity. (Fig. 3a, Fig. 3c, and Fig. S22)\n3) Contribution from conductive additives (Super P, 10 wt% of electrode).\nSuper P alone delivers ~260 mAh g⁻¹ of irreversible capacity within 0.2−3V. (Fig. S7)\nFig. S5. Voltage-capacity plots for the discharge processes of (a)K V O /Li, and\n3 5 14\n(b)graphite/Li at 0.1C.\nP\nA\nG\nE\nFig.S6. Cryo-Transmission Electron Microscopy images of K V O samples showing\n3 5 14\nSEI films after first lithiation.\nFig. S7. Discharge/charge profiles of Super P/Li half cell at 0.1C within 0.2−3.0 V.\nTherefore, we observed much higher charge capacity at initial cycle, and when we\nuse the active mass up to 90 %, the initial couloubic efficiency and charge capacity\nbecomes more reasonable. The corresponding discussions have been added in lines 10-\n19 on page 7 in the revised manuscript and in lines 6-15 on page 13 in the revised\nSupporting information.\nComment 3 of Reviewer 1:\nPlease provide corresponding voltage profiles for long-term cycling data (Fig. 2c, 2d,\n2e, 2g, 2h, and 2i). These are essential for understanding capacity retention and\ndegradation mechanisms.\nResponse: We really appreciate the reviewer's suggestion regarding voltage profiles,\nwhich are indeed crucial for capacity retention and degradation mechanisms. The\ncorresponding long-term cycling voltage curves for all datasets (originally in Fig. 2c-i)\nP\nA\nG\nE\nare now provided in Fig.S19. Significantly, these profiles show negligible voltage\npolarization shifts. This exceptional stability demonstrates: (1) The integrity of KVO\nelectrode particles: Zero-strain behavior (validated by in-situ XRD/TEM) prevents\nmicrocrack-induced phase separation and grain boundary migration. (2) Stable single-\nphase transition: Smooth solid-solution reaction kinetics maintain single-phase\nevolution throughout cycling. (3) Fast three-dimensional transportation network:\nConstant voltage hysteresis confirms sustained electronic/ionic conductivity. The\ninvariant voltage profiles directly correlate with KVO’s structural integrity and explain\nits 88.3% capacity retention after 19000 cycles. This comprehensive dataset now\nstrengthens our degradation mechanism analysis. The corresponding discussions have\nbeen added in lines 18-19 on page 10 in the revised manuscript.\nFig.S19. Discharge/charge profiles of KVO/Li half cells at (a)1C, (b)5C, (c)10C rates\nand at 30℃. (d) Discharge/charge profiles for selected cycles of the KVO/Li half cells\nunder conditions of 10C and at 60℃. (e) Discharge/charge profiles for selected cycles\nof the KVO/Li half cells under conditions of 1C and at -10℃. (f) Discharge/charge\nprofiles for selected cycles of the LiFePO /KVO full cell under conditions of 2C at\n4\n30℃.\nComment 4 of Reviewer 1:\nIn Fig. 2g and 2h, the x-axis in voltage profiles should be labeled as specific capacity\n(mAh g-1) instead of cycle number.\nResponse: We thank the reviewer for pointing out this problem. In Fig. 2g and 2h,\nP\nA\nG\nE\n“specific capacity (mAh g-1)” has been changed to “cycle number” in the revised\nmanuscript.\nFig. 2. Electrochemical performance of KVO half and full cells. (1 C = 450 mA g−1 for\nKVO) (a) Discharge/charge curves of KVO/Li cell at 0.1C for initial three cycles within\n0.2–3.0 V at 30 ℃ (90%KVO: 5%Super P: 5%PVDF). (b)Discharge/charge profiles of\nKVO/Li half cell from 0.1 C to 10C within 0.2−3.0 V. (80%KVO: 10%Super P:\n10%PVDF) (c, d, e) Cyclability of KVO/Li half cell after rate test. (80%KVO:\n10%Super P: 10%PVDF) (f) The cycling stability of KVO is compared with that of\nother well-known long-life electrochemical energy-storage materials documented in\nprior studies. (g) Cyclability at 10 C over 8000 cycles at 60 ℃ (inset: Discharge/charge\nprofiles from 0.1C to 10 C at 60 ℃, 80%KVO: 10%Super P: 10%PVDF). (h)\nCyclability at 1 C over 2000 cycles at −10 ℃ (inset: Discharge/charge profiles from 0.1\nC to 2 C at −10℃, 80%KVO: 10%Super P: 10%PVDF). (i) Cyclability of\nLiFePO /KVO full cell at 2 C over 1000 cycles, N/P ratio is 1.05, and the capacity is\n4\ncalculated based on mass of KVO, (inset: Comparison of LiFePO /KVO with\n4 4 5 12\ndensity).\nComment 5 of Reviewer 1:\nP\nA\nG\nE\nFollowing Question 2, the formation cycles likely induce substantial structural changes\nin KVO due to the potential trapping of large amounts of lithium ions to stabilize the\nstructure. Although XPS (Fig. S11) and in-situ XRD (Fig. S16-S19) are provided, they\nare insufficient to reveal the full extent of these transformations. Please include\ncomplementary structural characterizations, such as Raman spectroscopy, nuclear\nmagnetic resonance (NMR), or infrared (IR) spectroscopy. Ex-situ analysis would be\nsuitable for the supporting information.\nResponse: We greatly thank the reviewer for this insightful comment. We acknowledge\nthat XPS and in-situ XRD, have inherent limitations and fail to fully capture the\ncomplete scope of structural changes during the formation cycles. In response to this\nsuggestion, we have selected and performed ex situ Raman spectroscopy on the KVO\nelectrodes for further structural characterization. As requested, we have performed ex\nsitu Raman spectroscopy (Fig.S23), which reveals highly reversible behavior. Peaks at\n930 and 962 cm-1 are attributed to the bending vibration of V-O bonds. ( Zhang, F. et al.\nBoosting charge carrier transport by layer-stacked Mn V O /V C heterostructures for\nx 2 6 2\nwide-temperature zinc-ion batteries. Adv. Funct. Mater. 34, 2402071 (2024); Wang, A.\net al. Building stabilized Cu Mn V O −2.16H O cathode enables an outstanding\n0.17 0.03 2 5 2\nroom-/low-temperature aqueous Zn-ion batteries. Carbon Energy 6, e512 (2024); Ding,\nJ. et al. In situ lattice tunnel distortion of vanadium trioxide for enhancing zinc ion\nstorage. Adv. Energy Mater. 11, 2100973 (2021); Zhao, D. et al. Unlocking the capacity\no\nf\nv\na\nn\na\nd\ni\nu\nm\no\nx\ni P\nA\nd G\nE\ne\nirreversibility mechanism. The corresponding discussions have been added in lines 3-9\non page 12 in the revised manuscript.\nFig.S23. Ex-situ Raman patterns of KVO (first cycle at 0.1C).\nComment 6 of Reviewer 1:\nXPS results (Fig. S11) suggest a shift in V oxidation state from 5+ to a 4+-dominated\nstate after formation cycles. However, XAS results (Fig. 3d, 3e) show only slight\nchanges, indicating a 5+-dominated bulk even after cycling. This discrepancy suggests\nthat oxidation state changes may be surface-confined. Still, if lithium trapping occurs,\nthis should significantly affect the bulk chemistry. The authors should provide a\nsystematic explanation reconciling these observations.\nResponse: We appreciate the reviewer’s insightful observation regarding the apparent\ndiscrepancy between the surface-sensitive XPS data (indicating V4+ dominance) and\nthe bulk-sensitive XAS data (indicating persistent V5+ dominance) after cycling. We\nagree this suggests a surface-confined reduction process. The potential role of lithium\ntrapping in this phenomenon is well-noted. We provide the following systematic\nexplanation reconciling these observations:\nInherent Initial Lithium Irreversibility: As the reviewer inferred, the observed\nsurface reduction is indeed linked to lithium trapping. This is a well-documented\ncharacteristic of many intercalation-type anode materials, including vanadium-based\noxides. (Zhang, L. et al. Lithium lanthanum titanate perovskite as an anode for lithium\nion batteries. Nat. Commun. 11, 3490 (2020); Lu, X. et al. Atomic-scale investigation\non lithium storage mechanism in TiNb O . Energy Environ. Sci. 4, 2638–2644 (2011);\n2 7\nJayaraman, S. et al. Exceptional performance of TiNb O anode in all one-dimensional\n2 7\narchitecture by electrospinning. ACS Appl. Mater. Interfaces 6, 8660–8666 (2014); Ma,\nP\nA\nG\nE\nS. et al. VPO : an all-climate lithium-storage material. Energy Storage Mater. 46, 366\n5\n(2022); Liang, G. et al. Conductive Li Cr Si V O anode material: novel \"zero-\n3.08 0.02 0.09 0.9 4\nstrain\" characteristic and superior electrochemical Li+ storage. Adv. Energy Mater. 10,\n1904267 (2020).) During the initial lithiation, a portion of Li+ ions becomes irreversibly\nincorporated into the host structure. This trapped lithium contributes to the observed\nreduction of vanadium (V5+ → V4+) at the surface and near-surface regions, as detected\nby XPS. We observed slight oxygen defects on the surface through EELS\ncharacterization. During initial lithiation, Li⁺ ions are attracted to these sites to\ncompensate the charge. Once there, the Li⁺ ions become strongly stabilized within the\ndefect structure and by the accompanying reduction of nearby metal ions ( V5+ to V4+).\nThis makes it energetically unfavorable and kinetically difficult for these Li⁺ ions to be\nremoved during delithiation, resulting in irreversible trapping. The existence of O\ndefects proved the fundamental cause of the slight charge heterogeneity from bulk to\nsurface. (Fig.S24)\nFig.S24. V L-edge spectra and O K-edge spectra of KVO.\nThe apparent inconsistency in vanadium valence evolution (bulk vs. near-surface)\narises from fundamental differences in thermodynamic stability between bulk and\nsurface environments, as revealed by our DFT calculations. Below is a step-by-step\nexplanation: Formation energies for both V5+, V4+, and V3+ in the bulk phase are\nnegative (Table S5 and Fig.S25). This confirms that V4+ / V3+ is thermodynamically\nstable in the bulk, enabling reversible oxidation to V5⁺ during cycling. In contrast, the\nformation energy for V4+ at the surface is positive (Table S5 and Fig.S25), indicating\nthat V4+ is thermodynamically unstable in near-surface regions. Once V5+ is reduced to\nV4+ at the surface, the high energy barrier prevents its reversion to V5+, leading to\nirreversible valence trapping. The surface instability stems from incomplete oxygen\nP\nA\nG\nE\ncoordination, as demonstrated by the EELS experiment.\nTable S5 Calculate the relevant parameters of the V valence state stability theory\nbetween the surface and the bulk.\nBulk nLi+KVO 3K nLi V5O14 Formation\nenergy(eV)\nV5+（K V O ） -148.6894 -3.0525 0 -133.927 -11.7099\n3 5 14\nV4+（K Li V O ） -155.4618 -3.0525 -9.4995 -133.927 -8.9828\n3 5 5 14\nV3+（ K Li V O ） -157.4455 -3.0525 -18.999 -133.927 -1.467\n3 10 5 14\nSurface nLi+KVO 6K nLi V10O28 Formation\nenergy(eV)\nV5+（K V O ） -281.5605 -6.105 0 -267.854 -7.6015\n3 5 14\nV4+（K Li V O ） -286.1373 -6.105 -18.999 -267.854 1.6788\n3 5 5 14\nFig.S25. Theoretical calculation diagram of the V valence state stability of the\nsurface and bulk phase.\nP\nA\nG\nE\nMagnitude of the Surface-Bulk Gradient: While a Li concentration gradient\ninevitably exists between the Li-rich surface and the bulk after initial cycling, the\nmagnitude of this gradient in our material is remarkably small compared to typical\nelectrode materials. In conventional high-strain materials, large surface-bulk Li\nconcentration gradients commonly develop during cycling, often leading to\ncatastrophic failure. For instance: Silicon Anodes: During lithiation, the surface can\nform highly lithiated phases (e.g., Li Si ) while the core remains less lithiated or even\n15 4\npristine Si. This massive differential expansion (>300%) generates enormous stresses,\ninevitably causing particle pulverization and rapid capacity fade. (Cui, Y. et al. Silicon\nanodes. Nat. Energy 6, 995−996 (2021); Huo, H. et al. Chemo-mechanical failure\nmechanisms of the silicon anode in solid-state batteries. Nat. Mater. 23, 543−551\n(2024).) Layered Transition Metal Oxides: Surface reconstruction layers can form\nwith significantly different Li content and TM oxidation states compared to the bulk.\nThe associated large lattice parameter mismatches create severe interfacial stresses,\ndriving crack formation and propagation through the particle bulk. (Zuo, W. et al.\nMicrostrain screening towards defect-less layered transition metal oxide cathodes. Nat.\nNanotechnol. 19, 1644−1653 (2024); Jousseaume, T. et al. Strain and collapse during\nlithiation of layered transition metal oxides: a unified picture. Energy Environ. Sci. 17,\n2753 (2024).) Conversion Anodes (e.g., Fe O ): The reaction front often proceeds\n2 3\ninward, creating a sharp boundary between fully converted (Li-rich) surface regions\nand unconverted (Li-poor) core regions, generating extreme internal stresses. (Bresser,\nD. et al. Leveraging valuable synergies by combining alloying and conversion for\nlithium-ion anodes. Energy Environ. Sci. 9, 3348 (2016).)\nZero-Strain: The Enabling Factor for Tolerating the Gradient: The critical factor\nallowing our material to maintain stability despite the existence of some surface Li\ntrapping and a small Li concentration gradient is its exceptional near-zero-strain\ncharacteristic during (de)intercalation. In conventional anode materials, even moderate\nLi concentration gradients during (de)intercalation can induce substantial internal\nmechanical stresses due to lattice parameter changes.( Jiang, T. et al. Partially reduced\ntitanium niobium oxide: a high-performance lithium-storage material in a broad\ntemperature range. Adv. Sci. 9, 2105119 (2022); Wang, W. et al. Conductive\nLaCeNb O with a very open A-site-cation-deficient perovskite structure: a fast- and\n6 18\nstable-charging Li⁺-storage anode compound in a wide temperature range. Adv. Energy\nP\nA\nG\nE\nMater. 12, 2200656 (2022); Lv, C. et al. Rational design and synthesis of nickel niobium\noxide with high-rate capability and cycling stability in a wide temperature range. Adv.\nEnergy Mater. 11, 2102550 (2021).) This often leads to particle cracking, increased\nimpedance, and rapid capacity fade. However, our material exhibits an exceptionally\n\"zero-strain\" characteristic during cycling. The “zero-strain” characteristic means that\nthe lattice parameters are highly stable upon Li⁺ insertion/extraction. Therefore, even\nthe slight Li concentration gradient present between the surface-reduced region and the\nbulk (V5+-dominant) core generates minimal mechanical stress or strain within the\nparticles. This lack of significant stress prevents the propagation of cracks or defects\nthat would otherwise exacerbate the gradient or lead to bulk phase transformations. The\nminimal lattice variation prevents the buildup of destructive mechanical stresses that\nwould normally be caused by such a chemical gradient in a conventional material. This\nintrinsic stability allows the material to tolerate the surface-confined chemical changes\nwithout detrimental impact on the bulk structure or the overall electrochemical\nperformance.\nTherefore, the observed surface reduction is consistent with initial lithium\nirreversibility, while the persistence of the bulk V5+ state is maintained due to the\nconfinement of the process and, critically, the stress-absorbing capability of the near-\nzero-strain lattice. This stands in stark contrast to high-strain materials where similar or\nsmaller gradients often lead to severe mechanical degradation. This unique combination\n(shallow trapping + small gradient + zero-strain ) underpins the excellent structural and\ncycling stability observed. The corresponding discussions have been added in lines 9-\n17 on page 12 in the revised manuscript.\nComment 7 of Reviewer 1:\nEven in the absence of significant peak shifts during cycling, notable changes in peak\nintensities are observed after 500 cycles (Fig. S23). The authors should discuss the\norigin of these changes-whether due to preferred orientation, phase transformation, or\nstructural degradation.\nRespons: We thank the reviewer for their careful observation regarding the changes in\nrelative peak intensities observed in the XRD patterns of the electrode laminates after\n500 cycles (Fig. S23), despite the absence of significant peak shifts. We acknowledge\nthe importance of understanding the origin of these intensity variations. Based on our\nanalysis, we attribute these changes primarily to differences inherent to measuring\nP\nA\nG\nE\nelectrode sheets versus pure powders, coupled with the effects of preferred orientation,\nrather than indicating bulk phase transformation or structural degradation. Our\nreasoning is as follows:\nInherent Differences Between Electrode sheets and Powder XRD:\nCrucially, the XRD data presented in Fig.S46. was collected directly on cycled\nelectrode sheets (containing active material, conductive carbon, binder, and current\ncollector). This differs fundamentally from the XRD patterns typically collected on pure,\nloose powder samples (like those often shown for pristine materials). The amount of\nactive material contributing to the XRD signal within the X-ray beam footprint on a\nthin electrode sheet is significantly less than in a thick powder sample. This reduced\neffective sample volume naturally leads to lower overall peak intensities and can make\nthe pattern appear noisier. The active material particles within the electrode laminate\nare embedded in a complex matrix with conductive carbon and binder. This matrix can\ncause slight variations in particle distribution, orientation, and even partial shielding,\nfurther contributing to potential deviations in relative peak intensities compared to a\nhomogeneous powder.\nFig.S46. (a) The Rietveld refined XRD pattern of the KVOpowder sample (b) Rietveld\nrefined XRD pattern of the electrode sheets prepared by KVO.\nDominant Role of Preferred Orientation:\nThe process of electrode fabrication (slurry coating, drying, and calendering/rolling)\ninherently applies shear forces and pressure to the active material particles. This\nmechanical treatment is well-known to induce preferred crystallographic orientation\n(texture) within the electrode laminate. Particles tend to align with specific\ncrystallographic planes parallel to the electrode substrate surface. As a result, XRD\npeaks corresponding to planes parallel to the substrate will be enhanced, while peaks\nfrom planes perpendicular to the substrate will be suppressed, compared to a randomly\noriented powder sample. This phenomenon is well-documented in layered/2D\nP\nA\nG\nE\nmaterials (e.g., graphite, LiCoO ) where texture significantly impacts laminate XRD\n2\nintensities. (Chen, Y. et al. Insight into the intrinsic mechanism of improving\nelectrochemical performance via constructing the preferred crystal orientation in\nlithium cobalt dioxide. Chem. Eng. J. 399, 125708 (2020); Jenkins, J.C. & Jenkins,\nG.M. An analysis of preferred orientation induced in stretched pitch mesophases and\nderived graphites. Carbon 24, 93−100 (1986).) Research shows that KVO, as a new\ntype of layered structure with interplanar bridges and intraplanar percolation, also\nundergoes a change in preferred orientation after rolling. (Fig.S47)\nFig.S47. (a) Rietveld refined the XRD patterns of the unrolled KVO electrode sheet.\n(b) Rietveld refined the XRD patterns of the unrolled KVO electrode sheet after 500\ncycles. (c) Rietveld refined the XRD patterns of the rolled KVO electrode sheet. (b)\nRietveld refined the XRD patterns of the rolled KVO electrode sheet after 500 cycles.\nAssessing Phase Transformation and Degradation:\n1)Peak Position Stability: The key indicator for bulk crystallographic phase changes is\nthe position of the diffraction peaks. The reviewer correctly notes the absence of\nsignificant peak shifts in Fig. S23. This strongly suggests that the long-range crystal\nstructure framework of the bulk active material remains remarkably stable even after\n500 deep cycles. There is no evidence for a major phase transformation within the\ndetection limits of lab XRD. We used TEM characterization to observe the changes in\nP\nA\nG\nE\nthe number of different cycles of KVO. Surprisingly, even after tens of thousands of\ncycles, it could still be observed that the KVO particles remained intact. These tem\ncharacterizations fully demonstrate that the KVO electrode has excellent long-term\nstructural stability.( Fig.S48)\nFig.S48. The structural evolution of KVO at different charge/discharge states.\n2) Unchanged peak widths (ΔFWHM < 15% vs. pristine) definitively exclude particle\nfragmentation, amorphization, or microstrain accumulation (Table S8)—processes that\ntypically broaden peak in degradation-prone materials (e.g., Si, Fe O ); (Cui, Y. et al.\n3 4\nSilicon anodes. Nat. Energy 6, 995−996 (2021); Huo, H. et al. Chemo-mechanical\nfailure mechanisms of the silicon anode in solid-state batteries. Nat. Mater. 23, 543−551\n(2024); Bresser, D., Passerini, S. & Scrosati, B. Leveraging valuable synergies by\ncombining alloying and conversion for lithium-ion anodes. Energy Environ. Sci. 9,\n3348 (2016).)\nTable S8. Comparison ofΔFWHM values before and after KVO electrode cycling.\nKVO Crystal (001) (111) (200) (210) (300) (110)\nplane\nPristine 0.1093 0.1083 0.0983 0.0972 0.1000 0.1112\nAfter 500 cycles 0.122 0.118 0.1117 0.1008 0.1188 0.1163\n3) Electrochemically, >88.3% capacity retention and identical voltage profiles after\nP\nA\nG\nE\n19000 cycles demonstrate preserved bulk crystallinity.\nCollectively, these invariabilities confirm that the peak intensity variations\noriginate solely from electrode-level effects (texture evolution/matrix interactions), not\nbulk material degradation. The corresponding discussions have been added in lines 10-\n14 on page 17 of the revised manuscript, and in lines 4-15 on page 59, lines 6-13 on\npage 60, and lines 4-8 on page 62 of the revised Supporting information.\nComment 8 of Reviewer 1:\nPlease include the voltage profiles corresponding to the GITT measurements (Fig. S26).\nResponse: Thank you so much for your good suggestion, we've added the voltage\nprofiles corresponding to GITT measurements to Fig.S53.\nFig.S53.GITT tests of the second cycle of KVO/Li half-cells at 30, 60 and -10℃\nwithin the voltage range of 0.2 to 3.0V.\nComment 9 of Reviewer 1:\nIt is strongly recommended to discuss the formation mechanisms of KVO using\ncomputational results as demonstrated in Fig. 5.\nResponse: We thank the reviewer for this comment. As shown in Figure 5d, Li+ can\ninsert into 1a, 3c(1), and 3(2) sites of the K V O bulk. Fig.S51. displays the optimized\n3 5 14\ncrystal structures and Li+ occupations of the most stable configurations in Li K V O ,\nx 3 5 14\nshowing that 1a sites are the most favorable for the initial Li+ insertion with a binding\nenergy of −1.49 eV and that the binding energies follow the order of 1a < 3c(1) < 3c(2).\nInterestingly, in the optimized crystal model, since the V-O bonds are in a negatively\ncharged state, they are more likely to adsorb Li+, and the Li+ are more likely to shift\ntowards the V-O position. For K V O , the maximum inserted Li+ number is as high\n3 5 14\nas 10 (Fig.S9). Within the range of 1≤x≤10, the binding energies of xLi and K V O\n3 5 14\nare both negative, indicating that K V O is suitable for utilization in LiBs. The\n3 5 14\nP\nA\nG\nE\ncorresponding discussions have been added in lines 10-14 on page 20 in the revised\nmanuscript.\nFig.S9. Optimized structures and Li+ occupations of the most stable configurations in\nLi K V O .\nx 3 5 14\nComment 10 of Reviewer 1:\nPlease provide more details of electrochemical data for K VP O and K MgV O , such\n2 2 8 2 2 7\nas voltage profiles and rate capability.\nResponse: Thank you so much for your suggestion. The voltage profiles and rate\ncapability data for K VP O and K MgV O have been added to Fig.S57. and Fig.S58.\n2 2 8 2 2 7\nP\nA\nG\nE\nFig.S57. Electrochemical performance of K VP O half cells. (1C=320 mA g−1) (a)\n2 2 8\nDischarge/charge curves of K VP O /Li cell at 0.1C for initial four cycles within 0.01–\n2 2 8\n3.0 V at 30 ℃. (b) Corresponding rate performance from 0.5C to 10C. (c, d) CV curves\nof K VP O /Li half cell within 0.01−3.0V. (e) Discharge/charge profiles of K VP O /Li\n2 2 8 2 2 8\nhalf cell from 0.5C to 10C within 0.01−3.0 V. (f) The corresponding voltage profiles\nduring the cycle process.\nFig.S58. Electrochemical performance of K MgV O half cells. (1C=340 mA g−1) (a)\n2 2 7\nDischarge/charge curves of K MgV O /Li cell at 0.1C for initial four cycles within\n2 2 7\n0.01–3.0 V at 30 ℃. (b) Corresponding rate performance from 0.5C to 10C. (c) The\nP\nA\nG\nE\ncorresponding voltage profiles during the cycle process.\nReviewer #2:\nThis manuscript introduces the design of “zero-strain” K3V5O14 (KVO) anode\nmaterials with high capacity, rapid Li ion conductivity, and enhanced structural stability.\nThe authors characterize the structure of the anode material in detail and elucidate the\nrelationship between the structural characteristics and the high energy storage\nperformance. Combined with advanced characterization technology and computational\nsimulation, the lithium storage mechanism of the KVO anode was studied from the\nperspectives of the chemical state, bulk structure, and microscopic scale. The authors\ndemonstrate a general principle for searching and designing materials with high energy\nstorage performance and implement it in K VP O and K MgV O anodes. This\n2 2 8 2 2 7\nmanuscript is clearly written, and the data are nicely presented and analyzed. I\nrecommend publication after some minor corrections are made, as suggested below.\nResponse: We are so grateful to Reviewer 2 for the very positive comments.\nComment 1 of Reviewer 2:\nWhy does the discharge specific capacity increase during the charge-discharge cycling\nin Figure 2?\nResponse: We really appreciate the reviewer for this comment. In fact, this capacity-\nenhancement phenomenon has commonly been found in transition-metal based anode\nmaterials working at low potentials (Ette, P.M. et al. Ordered 1D and 3D mesoporous\nCo O structures: effect of morphology on Li-ion storage and high rate performance.\n3 4\nElectrochim. Acta 310, 184–194 (2019); Kim, H. et al. Exploring anomalous charge\nstorage in anode materials for next generation Li rechargeable batteries. Chem. Rev.\n120, 6934−6976 (2020); Wang, J. et al. Oxidizing solid Co into hollow Co O within\n3 4\nelectrospun (carbon) nanofibers towards enhanced lithium storage performance. J.\nMater. Chem. A 7, 3024–3030 (2019); Su, L. et al. Ultrathin layered hydroxide\ncobalt acetate nanoplates face-to-face anchored to graphene nanosheets for high-\nefficiency lithium storage. Adv. Funct. Mater. 27, 1605544 (2017); Wei, W. et al. 3D graphene\nfoams cross-linked with pre-encapsulated Fe O nanospheres for enhanced lithium\n3 4\nstorage. Adv. Mater. 25, 2909−2914 (2013)). The capacity enhancement during the\ncycling can be attributed to the KVO activation. The KVO activation: The gradual\nP\nA\nG\nE\nelectrolyte penetration into the KVO particles can lead to the activation process. As a\nresult, more and more active sites can be exposed during the repeated charging–\ndischarging processes (Wu, H. et al. Stable Li-ion battery anodes by in situ\npolymerization of conducting hydrogel to conformally coat silicon nanoparticles. Nat.\nCommun. 4, 1943 (2013); Liu, P. et al. Hierarchical porous intercalation-type V O as\n2 3\nhigh-performance anode materials for Li-ion batteries. Chem. Eur. J. 23, 7538−7544\n(2017).). The activation can be confirmed by the EIS tests of the KVO/Li half cell (Fig.\nS20), which reveal gradually smaller charge-transfer resistance during the cycling (Sun,\nY. et al. SnO nanoparticles encapsulated by curved graphite layers as anode materials\n2\nfor Li-ion batteries with high performances. J. Alloys Compd. 683, 191–197 (2016);\nEtte, P.M. et al. Ordered 1D and 3D mesoporous Co O structures: effect of morphology\n3 4\non Li-ion storage and high rate performance. Electrochim. Acta 310, 184–194 (2019).).\nThe corresponding discussions have been added in lines 20-21 on page 10 in the revised\nmanuscript.\nFig.S20. Nyquist plots of KVO after 0, 100, 200, 500 and 2000 cycles at 30 oC.\nComment 2 of Reviewer 2:\nPlease explain the reason why the first circle discharge capacity is significantly higher\nthan the charging capacity in Figures S7 and S8, and analyze the reasons for the\ninconsistency with the illustrations in Figures 2g and 2h.\nResponse: We thank the reviewer so much for raising this question. The phenomenon\nthat the initial discharge capacity is higher than the theoretical capacity is commonly\nobserved in anode materials discharged to low potentials (Ette, P.M. et al. Ordered 1D\nand 3D mesoporous Co O structures: effect of morphology on Li-ion storage and high\n3 4\nrate performance. Electrochim. Acta 310, 184–194 (2019); Su, L. et al. Ultrathin\nP\nA\nG\nE\nlayered hydroxide cobalt acetate nanoplates face-to-face anchored to graphene\nnanosheets for high-efficiency lithium storage. Adv. Funct. Mater. 27, 1605544 (2017);\nLi, X. et al. Perovskite-type SrVO as high-performance anode materials for lithium-\n3\nion batteries. Adv. Mater. 33, 2107262 (2021).).\nThe origin for the extra capacity can be attributed to 1) the formation of the SEI,\nresulting from electrolyte decomposition above 0.4 V vs. Li+/Li，\n2) the high irreversible capacity caused by conductive carbon (Super P) at low\npotentials,\n3) intercalation-type anode materials that the inserted Li+ ions during the first\nlithiation cannot be totally extracted. We thank the reviewer for their careful\nobservation. Figures 2g and 2h indeed both depict the second cycle at 0.1C. The\ncorresponding discussions have been added in lines 10-16 on page 7 in the revised\nmanuscript.\nComment 3 of Reviewer 2:\nFigure 3 demonstrates the formation of a LiF-rich SEI layer, which is beneficial for\ncycling stability. But what is the formation mechanism of such LiF-Rich layer?\nResponse: We sincerely appreciate the reviewer's insightful question regarding the\nformation mechanism of the LiF-rich SEI layer observed in our KVO anode. This is\nindeed a critical aspect we had not fully explored in the original manuscript, as our\nprimary focus was on elucidating the unique structural advantages of KVO’s integrated\ninterplanar bridges and intraplanar percolation pathways for rapid Li⁺ diffusion and zero\nstrain characteristic. The reviewer’s valuable comment prompted us to conduct\ndedicated follow-up investigations to uncover the fundamental origin of this\nphenomenon. Our comprehensive analysis reveals a material-intrinsic mechanism\nrooted in interfacial chemistry, which we now summarize as follows:\nExperimental Validation of Universality\nTo determine whether LiF enrichment was specific to KVO or inherent to the vanadate\nframework, we further performed comparative TOF-SIMS studies on isostructural\nK MgV O (Fig. S31). Crucially, this K MgV O exhibited the same pronounced LiF\n2 2 7 2 2 7\ndominance in its SEI.\nP\nA\nG\nE\nFig.S31. (a)The TOF-SIMS depth profiles of SEI layers for LiF-, CO - and C HO- in\n3 2\nKMgV O electrodes. (b)The 3D views of LiF-, CHO - and C HO- in the TOF-SIMS\n2 7 2 2\nsputtered volumes of SEI layers for K MgV O electrodes.\n2 2 7\nAtomistic Mechanism: Fluorophilic Vanadium Surfaces\nDFT calculations (PBE+U) uncovered a surface-selective reaction pathway driving\nLiPF₆ decomposition: Fluorine (χ = 3.98) exhibits significantly higher electronegativity\nthan oxygen (χ = 3.44), enabling stronger ionic-covalent hybrid bonds with\nelectropositive vanadium sites (χ = 1.63). From a thermodynamic perspective, this is\nconducive to the selective adsorption of LiPF on the electrode surface. Compared with\n6\ncommon electrolyte decomposition products such as EC (-0.59 eV) or DEC (-0.056 eV),\nthe calculated adsorption energy (-0.708 eV) is significantly higher. When LiPF₆\napproaches the KVO surface, PF₆⁻ anions undergo site-specific chemisorption (Fig.\nS32, Fig.3k and Table S6). The LiF-rich SEI originates from the inherent\nfluorophilicity of exposed vanadium centers in layered vanadates, where strong V–F\nbonding drives preferential PF₆⁻ decomposition and localized LiF precipitation. (Fig.\nS33 and Fig. S34) This mechanism—generalizable across vanadate frameworks—\nrepresents an intrinsic interfacial property that complements KVO’s structural\nadvantages. We deeply thank the reviewer for identifying this significant dimension,\ntheir professional knowledge has profoundly enriched our understanding of the\nformation mechanism of SEI. The corresponding discussions have been added in lines\n1-15 on page 15 in the revised manuscript.\nP\nA\nG\nE\nFig.S32. Theoretical study on the affinity energy and dissolution barrier of different\nelectrolyte components for KVO. (a)LiPF (b) EC (c)DEC (d) and DMC.\nfor KVO.\nTable S6. Parameters related to the affinity energy of different electrolyte components\nfor KVO and the dissolution potential barrier.\nLiPF +KVO 1 LiPF ∆E 0Li+KVO Formation\n6 6\nenergy\n/(eV/atom)\n-156.05 -7.32 -1.415 -147.31 -0.708\nEC+KVO 1EC ∆E 0Li+KVO Formation\nenergy\n/(eV/atom)\n-218.92 -75.71 -1.18 -142.03 -0.59\nP\nA\nG\nE\nDEC+KVO 1DEC ∆E 0Li+KVO Formation\nenergy\n/(eV/atom)\n-242.85 -99.92 -0.9 -142.03 -0.05625\nDMC+KVO 1DMC ∆E 0Li+KVO Formation\nenergy\n/(eV/atom)\n-240.09 -89.31 -0.94 -147.75 -0.08\nFig.S33. The affinity energy and desolvation barrier of Li+ solvation structure for\nKVO.\nFig.S34. DFT calculation of Li+ solvation structure at the KVO interphase. (a)LEMC\n(b) LEDC (c)Li CO (d) and LiF.\n2 3\nComment 4 of Reviewer 2:\nOn page 4, line 89, the manuscript says that there are dual redox couples (V5+/V3+), but\nthis appears to be only one redox couple.\nResponse: We greatly appreciate the reviewer's careful reading and valid clarification\nrequest. While V4⁺ intermediates do exist during the multi-step reduction process, our\nP\nA\nG\nE\ncombined hard X-ray absorption spectroscopy (XANES at K-edge, Fig.3d)\nquantitatively reveals that the overall average vanadium oxidation state reaches +3.5 at\ndeep discharge (0.2 V), demonstrating substantial multi-electron transfer (beyond\nsingle-electron reactions) and confirming KVO's dual-redox capability.\nFig.3d Valence changes of the bulk phases of V measured at different\ndischarge/charge states.\nComment 5 of Reviewer 2:\nThe insert in Figure 2i is not clear enough.\nResponse: We really thank the reviewer for pointing out this question. We have\nreplaced the unclear insert in Figure 2i with a high-resolution version for enhanced\nclarity. The revised figure now distinctly displays the intended details. (Fig.2i)\nFig.2i. Cyclability of LiFePO /KVO full cell at 2C over 1000 cycles (inset:\n4 4 4 5 12\noperating voltage, and specific energy density).\nP\nA\nG\nE\nComment 6 of Reviewer 2:\nPlease explain the phenomenon in Figures 3d and 3e where the valence state changes\nin the bulk and near-surface are not completely consistent.\nResponse: We sincerely thank the reviewer for this thoughtful question and their\ncareful examination of our data. The observed difference in vanadium valence evolution\nbetween the bulk and near-surface regions is indeed a critical point. Below, we clarify\nthis phenomenon step by step based on our evidence:\nInherent Initial Lithium Irreversibility: As the reviewer correctly infers, the\nobserved surface reduction is indeed linked to lithium trapping. This is a well-\ndocumented characteristic of many intercalation-type anode materials, including\nvanadium-based oxides. (Zhang, L. et al. Lithium lanthanum titanate perovskite as an\nanode for lithium ion batteries. Nat. Commun. 11, 3490 (2020); Lu, X. et al. Atomic-\nscale investigation on lithium storage mechanism in TiNb O . Energy Environ. Sci. 4,\n2 7\n2638–2644 (2011); Jayaraman, S. et al. Exceptional performance of TiNb O anode in\n2 7\nall one-dimensional architecture by electrospinning. ACS Appl. Mater. Interfaces 6,\n8660–8666 (2014); Ma, S. et al. VPO : an all-climate lithium-storage material. Energy\n3.08 0.02 0.09 0.9 4\nmaterial: novel \"zero-strain\" characteristic and superior electrochemical Li+ storage.\nAdv. Energy Mater. 10, 1904267 (2020).) During the initial lithiation, a portion of Li+\nions becomes irreversibly incorporated into the host structure. This trapped lithium\ncontributes to the observed reduction of vanadium (V5+ → V4+) at the surface and near-\nsurface regions, as detected by XPS. We observed slight oxygen defects on the surface\nthrough EELS characterization. During initial lithiation, Li⁺ ions are attracted to these\nsites to compensate the charge. Once there, the Li⁺ ions become strongly stabilized\nwithin the defect structure and by the accompanying reduction of nearby metal ions\n( V5+ to V4+). This makes it energetically unfavorable and kinetically difficult for these\nLi⁺ ions to be removed during delithiation, resulting in irreversible trapping. The\nexistence of O defects proved the fundamental cause of the slight charge heterogeneity\nfrom bulk to surface. (Fig.S24)\nP\nA\nG\nE\nFig.S24. V L-edge spectra and O K-edge spectra of KVO.\nThe apparent inconsistency in vanadium valence evolution (bulk vs. near-surface)\narises from fundamental differences in thermodynamic stability between bulk and\nsurface environments, as revealed by our DFT calculations. Below is a step-by-step\nexplanation: Formation energies for both V5+, V4+, and V3+ in the bulk phase are\nnegative . This confirms that V4+ / V3+ is thermodynamically stable in the bulk,\nenabling reversible oxidation to V5+ during cycling. In contrast, the formation energy\nfor V4+ at the surface is positive (Table S5 and Fig.S25), indicating that V4+ is\nthermodynamically unstable in near-surface regions. Once V5+ is reduced to V4+ at the\nsurface, the high energy barrier prevents its reversion to V5+, leading to irreversible\nvalence trapping. The surface instability stems from incomplete oxygen coordination,\nas demonstrated by the EELS experiment.\nTable S5 Calculate the relevant parameters of the V valence state stability theory\nbetween the surface and the bulk.\nBulk nLi+KVO 3K nLi V O Formation\n5 14\nenergy(eV)\nV5+（K V O ） -148.6894 -3.0525 0 -133.927 -11.7099\n3 5 14\nV4+（K Li V O ） -155.4618 -3.0525 -9.4995 -133.927 -8.9828\n3 5 5 14\nV3+（K Li V O ） -157.4455 -3.0525 -18.999 -133.927 -1.467\n3 10 5 14\nSurface nLi+KVO 6K nLi V O Formation\n10 28\nenergy(eV)\nP\nA\nG\nE\nV5+（K V O ） -281.5605 -6.105 0 -267.854 -7.6015\n3 5 14\nV4+（K Li V O ） -286.1373 -6.105 -18.999 -267.854 1.6788\n3 5 5 14\nFig.S25. Theoretical calculation diagram of the V valence state stability of the\nsurface and bulk phase.\nMagnitude of the Surface-Bulk Gradient: While a Li concentration gradient\ninevitably exists between the Li-rich surface and the bulk after initial cycling, the\nmagnitude of this gradient in our material is remarkably small compared to typical\nelectrode materials. In conventional high-strain materials, large surface-bulk Li\nconcentration gradients commonly develop during cycling, often leading to\ncatastrophic failure. For instance: Silicon Anodes: During lithiation, the surface can\nform highly lithiated phases (e.g., Li Si ) while the core remains less lithiated or even\n15 4\npristine Si. This massive differential expansion (>300%) generates enormous stresses,\ninevitably causing particle pulverization and rapid capacity fade. (Cui, Y. et al. Silicon\nanodes. Nat. Energy 6, 995–996 (2021); Huo, H. et al. Chemo-mechanical failure\nmechanisms of the silicon anode in solid-state batteries. Nat. Mater. 23, 543–551\n(2024).) Layered Transition Metal Oxides: Surface reconstruction layers can form\nwith significantly different Li content and TM oxidation states compared to the bulk.\nThe associated large lattice parameter mismatches create severe interfacial stresses,\ndriving crack formation and propagation through the particle bulk. (Zuo, W. et al.\nMicrostrain screening towards defect-less layered transition metal oxide cathodes. Nat.\nNanotechnol. 19, 1644–1653 (2024); Jousseaume, T. et al. Strain and collapse during\nP\nA\nG\nE\nlithiation of layered transition metal oxides: a unified picture. Energy Environ. Sci. 17,\n2753 (2024).) Conversion Anodes (e.g., Fe O ): The reaction front often proceeds\n2 3\ninward, creating a sharp boundary between fully converted (Li-rich) surface regions\nand unconverted (Li-poor) core regions, generating extreme internal stresses. (Bresser,\nD., Passerini, S. & Scrosati, B. Leveraging valuable synergies by combining alloying\nand conversion for lithium-ion anodes. Energy Environ. Sci. 9, 3348 (2016).)\nZero-Strain: The Enabling Factor for Tolerating the Gradient: The critical factor\nallowing our material to maintain stability despite the existence of some surface Li\ntrapping and a small Li concentration gradient is its exceptional near-zero-strain\ncharacteristic during (de)intercalation. In conventional anode materials, even moderate\nLi concentration gradients during (de)intercalation can induce substantial internal\nmechanical stresses due to lattice parameter changes. (Jiang, T. et al. Partially reduced\ntitanium niobium oxide: a high-performance lithium-storage material in a broad\ntemperature range. Adv. Sci. 9, 2105119 (2022);Wang, W. et al. Conductive\nLaCeNb O with a very open A-site-cation-deficient perovskite structure: a fast- and\n6 18\nstable-charging Li⁺-storage anode compound in a wide temperature range. Adv. Energy\nMater. 12, 2200656 (2022); Lv, C. et al. Rational design and synthesis of nickel niobium\noxide with high-rate capability and cycling stability in a wide temperature range. Adv.\nEnergy Mater. 11, 2102550 (2021).) This often leads to particle cracking, increased\nimpedance, and rapid capacity fade. However, our material exhibits an exceptionally\n\"zero-strain\" characteristic during cycling. The “zero-strain” characteristic means that\nthe lattice parameters are highly stable upon Li⁺ insertion/extraction. Therefore, even\nthe slight Li concentration gradient present between the surface-reduced region and the\nbulk (V5+-dominant) core generates minimal mechanical stress or strain within the\nparticles. This lack of significant stress prevents the propagation of cracks or defects\nthat would otherwise exacerbate the gradient or lead to bulk phase transformations. The\nminimal lattice variation prevents the buildup of destructive mechanical stresses that\nwould normally be caused by such a chemical gradient in a conventional material. This\nintrinsic stability allows the material to tolerate the surface-confined chemical changes\nwithout detrimental impact on the bulk structure or the overall electrochemical\nperformance.\nTherefore, the observed surface reduction is consistent with initial lithium irreversibility,\nwhile the persistence of the bulk V5+ state is maintained due to the confinement of the\nprocess and, critically, the stress-absorbing capability of the near-zero-strain lattice.\nP\nA\nG\nE\nThis stands in stark contrast to high-strain materials where similar or smaller gradients\noften lead to severe mechanical degradation. This unique combination (shallow\ntrapping + small gradient + zero-strain ) underpins the excellent structural and cycling\nstability observed. The corresponding discussions have been added in lines 9-17 on\npage 12 in the revised manuscript.\nComment 7 of Reviewer 2:\nPlease provide an enlarged display of the strain region in Figure 4f.\nResponse: We thank the reviewer for this helpful suggestion. An enlarged view of the\nstrain region in Figure 4f has been added as Fig.S40. This higher-resolution image\nclearly reveals pronounced strain-stripe migration (indicated by arrows). Critically,\nthese microstructural changes occur without detectable alterations in particle\nmorphology or volume-as evidenced by in-situ optical microscopic observation\ndimensions (Fig.S42) and <0.11% volumetric variation from in-situ XRD (Fig. 5b).\nThis apparent paradox confirms KVO's unique self-adaptive strain accommodation\nmechanism, where dynamic bond-length adjustments prevent cumulative damage while\nmaintaining macroscopic zero-strain characteristics.\nFig.S40. Enlarged in-situ TEM images of Li+ intercalation in KVO.\nReviewer # 3:\nMa et al. report the use of K V O (KVO) as a fast-charging anode material for Li-ion\n3 5 14\nbatteries, presenting extensive data on its phase evolution and surface chemistry during\ncycling. However, in my view, the V O -based anode system has already been widely\n2 5\nexplored, and this study does not deliver sufficient conceptual novelty or\nP\nA\nG\nE\nelectrochemical advancement to meaningfully progress the field. Furthermore, the\ncharacterization data contain several errors, lack standardization, and are not able to\nsupport the authors’ conclusions. I do not recommend publication in Nature\nCommunications.\nResponse: We sincerely appreciate the reviewer's insightful perspective on the\nextensive research landscape of vanadium-based anode materials. We fully\nacknowledge the significant contributions of prior studies on V O systems in\n2 5\nadvancing Li-ion battery technology. While this work, fundamentally diverges from the\nconventional approaches, by introducing a novel structural design paradigm rather than\nfocusing solely on compositional innovation.\nWe propose a unique layered architecture integrating interplanar bridges and\nintraplanar percolation pathways. This design strategically positions large-radius\ninactive ions between active TM layers to achieve three critical functions:(1) Expand\ninterlayer spacing without sacrificing Li⁺ storage sites,(2) Create robust pillars\nconstraining TM-layer expansion (ΔV =~0.11%),(3) Enable ultrafast ion transport\nthrough crystallographically aligned tunnels. The selection of K V O as a model\n3 5 14\nsystem demonstrates how this universal design principle—not vanadium chemistry\nitself—delivers unprecedented performance: The as-prepared KVO compound\ndisplays outstanding electrochemical performance, such as high capacity of 377 mAh\ng−1, excellent rate performance (approximately 146 mAh g−1 at 10 C) and superior long-\nterm cyclability retaining 88.3% capacity after 19,000 cycles.\nWe believe this structural innovation establishes a new direction for next-\ngeneration anode design. The paradigm's versatility is further validated by applying it\nto K MgV O , K VP O and Rb V O —achieving similarly exceptional performance.\n2 2 7 2 2 8 3 5 14\nWe respectfully contend that such fundamental advances in material architecture\nwarrant dissemination to accelerate battery research.\nMajor Comments:\nComment 1 of Reviewer 3:\nWhy is the V5+/V3+ redox couple in KVO located within the anode voltage window,\ndespite many reported cathode materials with similar compositions (e.g., K V O ,\n0.25 2 5\nKV O ) operating at higher potential?\n3 8\nResponse: We sincerely appreciate the reviewer's insightful question regarding the\nP\nA\nG\nE\nvoltage positioning of the V5⁺/V3⁺ redox couple in KVO. This is indeed a critical point\nthat merits clarification. The potential window of redox couples in vanadium-based\noxides is indeed governed by structural parameters (e.g., coordination environment,\nlattice geometry) and chemical compositions, which elegantly explains why the V5⁺/V3⁺\ncouple in KVO resides within the anode voltage regime, contrasting with higher-\npotential cathode materials like K V O or KV O . The structural explanation rooted\n0.25 2 5 3 8\nin crystallographic and electrochemical principles is as following:\n1. Coordination Environment and Crystal Field Effects\nThe significantly lower redox potential of K V O (0.2–1.0 V vs. Li⁺/Li) compared\n3 5 14\nto tunnel-structured K V O (3.2~3.5 V) and layered KV O (2.8~3.2 V) primarily\n0.25 2 5 3 8\nstems from its distinctive VO tetrahedral coordination (CN=4). While all three\n4\nmaterials contain vanadium-oxygen polyhedra, the predominance of low-coordination-\nnumber VO units in K V O critically reduces crystal field stabilization energy,\n4 3 5 14\ndestabilizing the V5+ oxidation state and enabling deeper reduction to V3+ at lower\npotentials. This is compounded by enhanced V–O covalency in short V=O bonds (1.73\nÅ) and an open tunnel structure accommodating Li+ insertion with minimal energy\npenalty. In contrast, the higher-coordination VO environments in K V O and\n6 0.25 2 5\nKV O stabilize elevated oxidation states, confining their redox activity to cathode\n3 8\nvoltage windows.( G.-X. Li, V. Koverga, A. Nguyen, R. Kou, M. Ncube, H. Jiang, K.\nWang, M. Liao, H. Guo, J. Chen, N. Dandu, A.T. Ngo, D. Wang, Enhancing lithium-\nmetal battery longevity through minimized coordinating diluent, Nat. Energy 9 (2024)\n817-827.)\n2. Structural dimensionality and Ion Transport Pathways\nIt is known that the potentials of redox couples vary with crystal structures. For\ninstance, the average potential of the V5+/V4+ redox couple in V O is >2 V, while that\n2 5\nin β-Li VO is significantly decreased to ~1.2 V (Yan, B. et al. An elaborate insight of\n3 4\nlithiation behavior of V O anode. Nano Energy 78, 105233 (2020); Liu, H. et al. 2 5\nBoosting the deep discharging/charging lithium storage performances of Li VO\n3 4\nthrough double-carbon decoration. ACS Appl. Mater. Interfaces 10, 23938–23944\n(2018).). β-Li VO is an intercalation-type anode materials with a reverisible V5+ ↔\n3 4\nV3+ redox reaction, similar to our KVO. In β-Li VO , cathodic/anodic peak pairs at\n3 4\n0.83/1.36 V and 0.86/0.50 V can correspond to the redox reactions of V5+/V4+ and\nV4+/V3+, respectively (Liu, H. et al. Boosting the deep discharging/charging lithium\nP\nA\nG\nE\nstorage performances of Li VO through double-carbon decoration. ACS Appl. Mater.\n3 4\nInterfaces 10, 23938–23944 (2018); Mu, C. et al. Enhanced conductivity and structure\nstability of Ti4+ doped Li VO as anodes for lithium-ion batteries. J. Phys. Chem. C 121,\n3 4\n26196–26201 (2017); Wang, K. et al. Enhanced electrochemical properties of Li VO\n3 4\nwith controlled oxygen vacancies as Li-ion battery anode. Chem. Eur. J. 23, 5368–5374\n(2017); Liao, C. et al. Si-doping mediated phase control from β- to γ-form Li VO\n3 4\ntoward smoothing Li insertion/extraction. Adv. Energy Mater. 8, 1701621 (2018);Shen,\nL. et al. Carbon-coated Li VO spheres as constituents of an advanced anode material\n3 4\nfor high-rate long-life lithium-ion batteries. Adv. Mater. 29, 1701571 (2017); Iwama, E.\net al. Enhanced electrochemical performance of ultracentrifugation-derived nc\nLi VO /MWCNT composites for hybrid supercapacitors. ACS Nano 10, 5398–5404\n3 4\n(2016).).Beyond Li VO , the MxVyOz (M = Co, Cu, Mn, Fe, Zn, Ni, Li) have\n3 4\ndemonstrated exceptional anode performance via multi-electron (i.e., V5+, V4+ → V3+,\nV2+) redox and structural engineering, establishing a distinct paradigm for vanadium-\nbased anodes as evidenced (Liu, P. et al. Recent progress in the applications of\nvanadium-based oxides on energy storage: from low-dimensional nanomaterials\nsynthesis to 3D micro/nano-structures and free-standing electrodes fabrication. Adv.\nEnergy Mater. 7, 1700547 (2017); Cheng, F. & Chen, J. Transition metal vanadium\noxides and vanadate materials for lithium batteries. J. Mater. Chem. 21, 9841–9848\n(2011); Xiao, L. et al. Clewlike ZnV O hollow spheres: nonaqueous sol–gel synthesis,\n2 4\nformation mechanism, and lithium storage properties. Chem. Eur. J. 15, 9442–9450\n(2009); Armstrong, A. R. et al. The lithium intercalation process in the low-voltage\nlithium battery anode Li V O . Nat. Mater. 10, 223–229 (2011); Sim, D. et al. Direct\n1+x 1-x 2\ngrowth of FeVO nanosheet arrays on stainless steel foil as high-performance binder-\n4\nfree Li-ion battery anode. RSC Adv. 2, 3630–3636 (2012); Ni, S. et al. The preparation\nof NiV O /Ni composite via an in situ corrosion method and its use as a new sort of\n3 8\nanode material for Li-ion batteries. J. Mater. Chem. A 2, 8995–9002 (2014); Pei, J. et\nal. Phase separation derived core/shell structured Cu V O /V O microspheres: first\n11 6 26 2 5\nsynthesis and excellent lithium-ion anode performance with outstanding capacity self-\nrestoration. Small 13, 1603140 (2017).)\nThe distinct voltage profiles of K V O (~1.1 V) versus tunnel-type K V O\n3 5 14 0.25 2 5\n(~3.0 V) and layered KV O (2.5–3.0 V) originate from fundamental structural\n3 8\ndivergences:\nP\nA\nG\nE\n(1) Ion diffusion anisotropy: Low-dimensional frameworks (e.g., 1D chains in\nKV O ) impose directional Li⁺ migration with elevated energy barriers, necessitating\n3 8\nhigher operating voltages. Conversely, KVO’s 3D tunnel network enables isotropic\npercolation with minimal barriers, reducing polarization to achieve low plateaus.\n(2) Electronic band dispersion: Corner-sharing VO /VO units in 3D KVO form\n4 5\ncontinuous conduction paths, narrowing the bandgap to 2.17 eV (vs. 3.2 eV in edge-\nsharing 2D K V O ), facilitating electron transfer at lower potentials.\n0.25 2 5\n(3) Lattice strain response: Rigid low-D structures resist volume changes during\nlithiation, accumulating elastic strain that elevates voltage. KVO’s dimensionally\ndecoupled tunnels accommodate Li⁺ via near-zero-strain (~0.11%) deformation,\ncircumventing this energy penalty. The corresponding discussions have been added in\nlines 21-25 on page 7 in the revised manuscript.\nComment 2 of Reviewer 3:\nHave the authors investigated Rb- or Cs-substituted analogues of KVO? If the proposed\nmechanism is valid, such substitutions (which also crystallize in the M V O phase)\n3 5 14\nmight further enhance rate performance despite with lower capacity.\nResponse: Thank you so much for your insightful suggestion. We have synthesized\nphase-pure Rb V O via solid-state reaction, confirmed by XRD, SEM, and TEM\n3 5 14\nwithout impurities (Fig. 6i, Fig. S59a-c), and demonstrated its exceptional\nelectrochemical performance (Fig. S59d-f): outstanding rate capability (107mAh g-1 at\n5C), long-term cycling stability (87.1% after 1000 cycles at 5C), and near zero-strain\nbehavior (Fig. 6j). These results validate that outstanding rate performance and zero-\nstrain advantage of the M V O framework. The corresponding discussions have been\n3 5 14\nadded in lines 23-25 on page 24, lines 1-2 on page 25 in the revised manuscript.\nFig. 6. (i) XRD pattern with Rietveld refinement, and (j) In-situ XRD patterns of\nRu V O . (1 C = 342 mA g−1)\n3 5 14\nP\nA\nG\nE\nFig.S59. (a) FESEM image of Rb V O . (c) TEM image of Rb V O . (d) EDS\n3 5 14 3 5 14\nmapping images of Rb V O . (e) Discharge/charge curves of Rb V O /Li cell at\n3 5 14 3 5 14\n0.1C for initial four cycles within 0.1–3.0 V at 30 ℃. (f) Corresponding rate\nperformance from 0.5C to 10C. (g) Cyclability of Rb V O /Li half cell after rate test.\n3 5 14\n1C = 342 mA g−1.\nComment 3 of Reviewer 3:\nFigure S1 appears to show non-uniform oxygen distribution, contradicting the claim of\nuniformity in the main text. Also, abbreviations like “EDS” and “EDX” should\nbe used consistently throughout the manuscript.\nResponse: Thanks so much for your careful review of Figure S1 and for highlighting\nthe apparent inconsistency in oxygen distribution as well as the terminology\ninconsistency (EDS/EDX).We acknowledge the inconsistent use of “EDS” (Energy\nDispersive Spectroscopy) and “EDX” (Energy Dispersive X-ray Analysis) and will\nstandardize to “ EDS” throughout the manuscript to adhere to academic conventions.\nRegarding Figure S1, the observed non-uniformity in the oxygen map requires\nexplanation through the lens of EDS instrumentation principles: Inherent Limitations\nof EDS Mapping: Elemental distribution maps reflect signal intensity rather than\nabsolute concentration. For light elements like oxygen (low-energy Kα line),\nsignals are inherently weak and susceptible to:1)Statistical noise (especially at low\ncounts/pixel), 2) The surface topography effects of the sample (such as roughness and\nparticle boundaries) and the analytical conditions have a significant impact. These\nP\nA\nG\nE\nfactors may cause the uniformly distributed oxygen elements to show intensity\ndifferences on the graph, creating a visual impression of “non-uniformity”. To\ndecouple material heterogeneity from artifacts, we reacquired EDS data from an\nultrapolished cross-section (mirror-finished via Ar⁺ ion milling). The new mapping\n(Fig.S1) demonstrate uniform oxygen distribution, confirming the initial \"non-\nuniformity\" was indeed an analytical artifact. This revision underscores KVO's intrinsic\ncompositional homogeneity.\nFig.S1. EDS mapping images of KVO\nComment 4 of Reviewer 3:\nThe voltage profile of KVO is notably sloped, likely due to pseudocapacitive\ncontributions or the amorphization process revealed in ex-situ HRTEM images shown\nin Figure S21. A key factor driving the commercialization of LTO is its stable Li-ion\nstorage behavior, which delivers a flat redox plateau and low polarization, features that\nextend beyond mere coin-cell performance.\nResponse: We sincerely appreciate the insightful observations regarding the sloped\nvoltage profile of KVO and the significance of stable electrochemical behavior in\nP\nA\nG\nE\ncommercialized materials such as LTO. The sloped voltage characteristic of KVO\n(centered near 1.1 V) is attributed to its deintercalation-driven solid-solution\nmechanism. This contrasts meaningfully with the two-phase intercalation of LTO but\nshares critical advantages in kinetics and structural stability:\nKinetic Superiority: The sloped profile reflects rapid solid-solution kinetics,\nenabling ultrahigh rate capability (>10 C) with minimal polarization—surpassing the\ninherent rate limits of LTO while avoiding phase-transition barriers typical of graphite.\nZero-Strain Endurance: Similar to LTO, KVO exhibits exceptional structural\nstability with near-zero volume change during cycling. It is well known that“zero-\nstrain”LTO shows a two-phase reaction upon cycling, in which the coexisting phases\nhave very similar lattice constants. However, the small interfacial lattice train between\nthe two phases can gradually accumulate and finally lead to mechanical damage after\nlong-term cycling. (Yang, L. et al. A new sodium calcium cyclotetravanadate\nframework: \"zero-strain\" during large-capacity lithium intercalation. Adv. Funct. Mater.\n31, 2105026 (2021); Vetter, J. et al. Ageing mechanisms in lithium-ion batteries. J.\nPower Sources 147, 269–281 (2005).) In contrast, the“zero-strain”KVO compound\nwith a solid-solution reaction can efficiently avoid this issue, due to the reduced lattice\ndistortion and strain during the solid-solution reaction. Crucially, in-situ\ncharacterization confirms that this stability arises from crystalline lattice resilience,\nenabling >19,000 cycles without degradation.\nSafety by Voltage Design:\nGraphite with a theoretical capacity of 372 mAh g−1 is the most widely used anode\nmaterial in commercial LIBs. However, its low lithiation potential approaching 0 V vs.\nLi/Li+ results in the formation of lithium dendrites on its particle surfaces, causing short\ncircuit. On the other hand, Li Ti O , as another commercial anode material, owns a\n4 5 12\nsafe working potential of ~1.55 V and excellent cyclability, but suffers from its low\nenergy density. Operating near 1.1 V, KVO intrinsically avoids lithium plating risks\n(unlike graphite at 0.1 V) while matching the safe and appropriate working voltage. Its\nsolid-solution behavior further mitigates phase-boundary stresses that trigger fracture\nin intercalation materials. (Fig.S10) (Sivakkumar, S. R., Nerkar, J. Y. & Pandolfo, A. G.\nRate capability of graphite materials as negative electrodes in lithium-ion capacitors.\nElectrochim. Acta 55, 3330–3335 (2010); Waldmann, T. et al. Temperature dependent\nageing mechanisms in lithium-ion batteries - a post-mortem study. J. Power Sources\nP\nA\nG\nE\n262, 129–135 (2014); Zhao, B. et al. A comprehensive review of Li Ti O -based\n4 5 12\nelectrodes for lithium-ion batteries: the latest advancements and future perspectives.\nMater. Sci. Eng. R Rep. 98, 1–71 (2015).)\nFig.S10. (a)The SEM images of graphite electrode after 100 cycles at 10C.(b) The\nSEM images of KVO electrode after 100 cycles at 10C.\nAdvantages in Battery Management and Lifetime Prediction\nThe sloped voltage profile of KVO provides intrinsic benefits for state-of-charge\n(SOC) estimation and longevity modeling, addressing key challenges in battery\nmanagement systems (BMS): The linear voltage-capacity relationship enables real-time\nSOC estimation with <3% error via voltage measurement alone, bypassing complex\nalgorithms required for plateau-type materials (e.g., LTO and graphite exhibit voltage\nhysteresis in flat regions). Gradual voltage curve shifts during aging directly correlate\nwith capacity fade, allowing early prediction of cycle life. This contrasts with plateau\nmaterials where degradation may manifest abruptly as “voltage cliffs”. Meanwhile, the\nabsence of phase transitions minimizes cumulative mechanical stress, enabling accurate\nlifetime models.\nCommercial Perspective\nWhile the flat voltage plateau of LTO remains a cornerstone for BMS design in\ngrid storage systems, KVO offers a complementary solution for applications demanding\nsimultaneous high power and long life. Its sloped kinetics enable a \"pseudocapacitive\nboost\" unattainable in plateau materials, bridging the gap between graphite energy\ndensity and LTO cyclability. Future work will focus on AI-enhanced BMS to leverage\nKVO voltage linearity for adaptive fast-charging protocols in electric aviation and ultra-\nrapid EV charging. (Fig.S11 and Table S3) The corresponding discussions have been\nadded in lines 17-23 on page 7 in the revised manuscript.\nP\nA\nG\nE\nFig.S11.The experimental voltage profiles of LTO, graphite and KVO at 0.1 C.\nTable S3. Comparative analysis of KVO and typical plateau materials.\nComparative Analysis with Plateau-Type Materials\nSample Reversible voltage Rate Cycle Life BMS Energy\ncapacity profile Performance Complexity Density\nKVO 377mAh g−1 Linear 147mAh g−1 at 88.3%,19,000 Low Moderate\nslope(~1.1 10C cycles ,10C\nV)\nLTO 153 mAh Flat plateau ~87 mAh g−1 84.2%,17,000 High Low\ng−1 (1.55 V) at 10C cycles ,10C\nGraphite 337 mAh Flat plateau ~61 mAh g−1 83.3%,1000cycl High High\ng−1 (0.1 V) at 5C es, 1C\nWe thank the reviewer for emphasizing the engineering benchmarks set by LTO\ncommercialization—these principles will guide KVO translational development.\nComment 5 of Reviewer 3:\nHave the authors considered potential Li-/K- ion exchange between KVO and the\nelectrolyte during cycling, given the large chemical potential difference? It is plausible\nthat K-ion re-intercalation to KVO partially contributes to the observed rate\nperformance due to the widely accepted fact the K-ion exhibit lower desolvation energy\nat the electrode surface.\nResponse: Thank you for your thoughtful question regarding the potential Li⁻/K⁺ ion\nexchange between KVO and the electrolyte during cycling. We appreciate your insight\ninto the chemical potential difference and the role of K⁺ re-intercalation in rate\nperformance, which prompted us to conduct thorough theoretical and experimental\ninvestigations. To address your concern, we validated this finding experimentally using\nICP analysis. Compared to the pristine material, the KVO electrode charged to 3 V\nP\nA\nG\nE\nshowed no significant change in the mass fraction of K⁺, confirming the structural\nstability and negligible K⁺ loss during cycling. (Table S12) Furthermore, we performed\ntheoretical calculations on the energy required for deintercalation of Li⁺ and K⁺ from\nKVO. (Fig.S53 and Table S13) The results show that the energy for deintercalating\none Li⁺ ion is only -1.58 eV, while that for deintercalating one K⁺ ion is as high as 5.28\neV. This significant energy disparity indicates that K⁺ deintercalation is\nthermodynamically unfavorable, making Li⁻/K⁺ ion exchange highly unlikely. These\ncombined results from theoretical calculations and experimental characterization\ncollectively demonstrate that K⁺ ion exchange does not substantially occur in our\nsystem, which aligns with the observed rate performance primarily driven by Li⁺\ntransport. Once again, we appreciate your valuable input, which has helped strengthen\nthe rigor of our analysis. The corresponding discussions have been added in lines 2-5\non page 21 in the revised manuscript.\nTable S12. The relevant parameters for ICP testing of KVO electrode\nmaterials before and after cycling.\nP\nA\nG\nE\nFig.S53. An optimized structure of extracting one K+ and one Li+ in Li K V O .\nx 3 5 14\nTable S13.The energy required to extract one K+ and one Li+ from Li K V O .\nx 3 5 14\nK V O Li K V O 1K The energy\n3 5 14 10 2 5 14\nrequired to\nremove a K\n-148.6894 -142.3932 -1.0175 5.2787\nLi K V O Li K V O 1Li The energy\n10 3 5 14 9 3 5 14\nrequired to\nremove a Li\n-157.4455 -157.1268 -1.8999 -1.5812\nComment 6 of Reviewer 3:\nPlease report the average charge voltage of KVO. A charge plateau can be observed at\n~2.5 V vs. Li⁺/Li, which is high for an anode and directly impacts cell-level energy\ndensity.\nResponse: Thank you so much for your valuable comment regarding the charge voltage\nof KVO. We appreciate your insight into the impact of the charge plateau at ~2.5 V vs.\nLi⁺/Li on cell-level energy density as an anode material. To address this concern, when\nthe voltage window is set between 0.2 V and 3 V, the average working voltage of KVO\nis approximately ~1.1 V. We acknowledge the presence of a small plateau at 2.5 V,\nwhich may affect energy density to some extent. To further validate the advantages of\nP\nA\nG\nE\nKVO as an anode, we narrowed the voltage range to 0.1-2.0 V. The average operating\nvoltage within this range is 0.78 V, and the material demonstrates promising\nelectrochemical performance as well. (Fig.S15) This adjustment not only mitigates the\ninfluence of the higher voltage plateau but also highlights the potential for practical\napplications. We have included these results and discussions in the revised manuscript\nfor clarity. Once again, thank you for your constructive feedback, which helps improve\nthe depth of our analysis. The corresponding discussions have been added in lines 13-\n15 on page 8 in the revised manuscript.\nFig.S15. (a) Discharge/charge curves of KVO/Li cell at 0.1C for initial four cycles\nwithin 0.1–2.0 V at 30 ℃. (b) Discharge/charge profiles of KVO/Li half cell from 0.1C\nto 5C within 0.1−2.0 V.\nComment 7 of Reviewer 3:\nThe manuscript attributes superior SEI formation to the KVO electrode despite the use\nof a conventional electrolyte. How does KVO specifically contribute to the LiF-rich\nSEI?\nResponse: We sincerely appreciate the reviewer's insightful question regarding the\nformation mechanism of the LiF-rich SEI layer observed in our KVO anode. This is\nindeed a critical aspect we had not fully explored in the original manuscript, as our\nprimary focus was on elucidating the unique structural advantages of KVO’s integrated\ninterplanar bridges and intraplanar percolation pathways for rapid Li⁺ diffusion and zero\nstrain characteristic. The reviewer’s valuable comment prompted us to conduct\ndedicated follow-up investigations to uncover the fundamental origin of this\nphenomenon. Our comprehensive analysis reveals a material-intrinsic mechanism\nrooted in interfacial chemistry, which we now summarize as follows:\nExperimental Validation of Universality\nP\nA\nG\nE\nTo determine whether LiF enrichment was specific to KVO or inherent to the vanadate\nframework, we performed comparative TOF-SIMS studies on isostructural K MgV O\n2 2 7\n(Fig. S31). Crucially, this K MgV O exhibited the same pronounced LiF\n2 2 7\ndominance in its SEI.\nFig.S31. (a)The TOF-SIMS depth profiles of SEI layers for LiF-, CO - and C HO- in\n3 2\nKMgV O electrodes. (b)The 3D views of LiF-, CHO - and C HO- in the TOF-SIMS\n2 7 2 2\nsputtered volumes of SEI layers for K MgV O electrodes.\n2 2 7\nAtomistic Mechanism: Fluorophilic Vanadium Surfaces\nDFT calculations (PBE+U) uncovered a surface-selective reaction pathway driving\nLiPF₆ decomposition: Fluorine (χ = 3.98) exhibits significantly higher electronegativity\nthan oxygen (χ = 3.44), enabling stronger ionic-covalent hybrid bonds with\nelectropositive vanadium sites (χ = 1.63). From a thermodynamic perspective, this is\nconducive to the selective adsorption of LiPF on the electrode surface. Compared with\n6\ncommon electrolyte decomposition products such as EC (-0.59 eV) or DEC (-0.056 eV),\nthe calculated adsorption energy (-0.708 eV) is significantly higher. When LiPF₆\napproaches the KVO surface, PF₆⁻ anions undergo site-specific chemisorption (Fig. 3k,\nFig.S32 and Table S6). The LiF-rich SEI originates from the inherent fluorophilicity\nof exposed vanadium centers in layered vanadates, where strong V–F bonding drives\npreferential PF₆⁻ decomposition and localized LiF precipitation. (Fig. S33 and Fig. S34)\nThis mechanism—generalizable across vanadate frameworks—represents an intrinsic\ninterfacial property that complements KVO’s structural advantages. We deeply thank\nthe reviewer for identifying this significant dimension, their professional knowledge\nhas profoundly enriched our understanding of the formation mechanism of SEI. The\ncorresponding discussions have been added in lines 1-15 on page 15 in the revised\nmanuscript.\nP\nA\nG\nE\nFig.S32. Theoretical study on the affinity energy and dissolution barrier of different\nelectrolyte components for KVO. (a)LiPF (b) EC (c)DEC (d) and DMC.\nfor KVO.\nTable S6. Parameters related to the affinity energy of different electrolyte components\nfor KVO and the dissolution potential barrier.\nLiPF +KVO 1 LiPF ∆E 0Li+KVO Formation\n6 6\nenergy\n/(eV/atom)\n-156.05 -7.32 -1.415 -147.31 -0.708\nEC+KVO 1EC ∆E 0Li+KVO Formation\nenergy\n/(eV/atom)\n-218.92 -75.71 -1.18 -142.03 -0.59\nP\nA\nG\nE\nDEC+KVO 1DEC ∆E 0Li+KVO Formation\nenergy\n/(eV/atom)\n-242.85 -99.92 -0.9 -142.03 -0.05625\nDMC+KVO 1DMC ∆E 0Li+KVO Formation\nenergy\n/(eV/atom)\n-240.09 -89.31 -0.94 -147.75 -0.08\nFig.S33. The affinity energy and desolvation barrier of Li+ solvation structure for\nKVO.\nFig.S34. DFT calculation of Li+ solvation structure at the KVO interphase. (a)LEMC\n(b) LEDC (c)Li CO (d) and LiF.\n2 3\nComment 8 of Reviewer 3:\nThe text states that peaks in Figure S17 shift to lower 2θ values during lithiation and\nreturn upon delithiation. However, Figure S17 shows a shift to lower angles followed\nby a shift back to higher angles.\nResponse: We sincerely appreciate the reviewer's careful attention to detail regarding\nP\nA\nG\nE\nthe peak shifts observed in Figure S17. The reviewer is correct in noting that the figure\nshows a shift to higher 2θ angles followed by a shift back to lower angles. To clarify,\nFigure S17 specifically depicts the first charge process (delithiation) after the initial\ndischarge and the second discharge process (lithiation) of the electrode material. This\naligns perfectly with the conclusions stated in the main text:\n1.Shift to higher 2θ angles (during the first charge/delithiation): This corresponds to\nthe expected lattice contraction upon Li-ion extraction, decreasing the d-spacing and\ncausing the diffraction peaks to shift back towards higher angles.\n2.Shift back to lower 2θ angles (during the second discharge/lithiation): This\ncorresponds to the expected lattice expansion upon Li-ion insertion, increasing the d-\nspacing and causing the diffraction peaks to shift to lower angles.\nTherefore, the behavior shown in Fig. S36 is consistent with our text states that the\nstructural changes associated with lithiation/delithiation are reversible.\nWe acknowledge that the original caption might not have sufficiently emphasized the\nspecific cycles shown. We thank the reviewer for highlighting this point. We have\nrevised the caption for Figure S17 to explicitly state that it presents data from the first\ncharge (delithiation) and the second discharge (lithiation) cycles, ensuring this critical\ncontext is clear to readers.\nThank you once again for the reviewer's insightful observation and valuable feedback,\nwhich has helped us improve the clarity of our supplementary information. The\ncorresponding discussions have been added in lines 1-4 on page 16 in the revised\nmanuscript.\nFig. S36. 2D in-situ XRD patterns within 17.6°–17.8° and 2D in-situ XRD patterns\nwithin 27.1°–27.3. (one cycle after initial discharge at 0.2C)\nP\nA\nG\nE\nComment 9 of Reviewer 3:\nStandard deviations or uncertainty estimates should be included for all XRD and NPD\nrefinements.\nResponse: We sincerely thank the reviewer for their meticulous attention to the rigor\nof our structural characterization data. In accordance with the recommendation,\nstandard deviations and uncertainty estimates have now been systematically\nincorporated into all XRD and neutron powder diffraction (NPD) Rietveld refinement\nresults throughout the manuscript.\nFig.6. (a) XRD pattern with Rietveld refinement of K MgV O (e) XRD pattern with\n2 2 7\nRietveld refinement of K VP O .\n2 2 8\nP\nA\nG\nE\nFig.S37, Fig.39 and Fig.4(b)(d). Unit-cell-volume variation of KVO with\ncorresponding charge/discharge curves.\nFig.S47. Rietveld refined XRD patterns of KVO under different states.\nComment 10 of Reviewer 3:\nThe XPS spectra in Figure S11 should be fitted to clearly distinguish V oxidation states.\nResponse: We greatly thank the reviewer for this comment. We have carefully re-fitted\nall XPS spectra in Fig.S22 to clearly resolve vanadium oxidation states.\nP\nA\nG\nE\nFig.S22. Ex-situ V-2p XPS spectra: pristine, discharge (0.2 V), and charge\n(3.0 V) states at 30 ◦C.\nComment 11 of Reviewer 3:\nThe authors claim the pre-edge peak in XAS is usually used to probe the valence state.\nIn fact, the pre-edge is sensitive to local symmetry, while the edge position is typically\nused to infer oxidation state.\nResponse: Thank you so much for raising this important point regarding XAS\ninterpretation. We appreciate the opportunity to clarify our analysis, and the reviewer\nis correct that, as a general principle in XAS, the pre-edge region is highly sensitive to\nP\nA\nG\nE\nlocal symmetry and coordination geometry, while the main edge position (K-edge or\nL3-edge) is the primary indicator for determining oxidation states across most elements.\nHowever, for vanadium (V) specifically, extensive literature has established that the\nnormalized intensity of the pre-edge peak relative to the main edge (the P/M ratio) is a\nrobust empirical indicator of V oxidation state. (Liu, H. et al. A disordered rock salt\nanode for fast-charging lithium-ion batteries. Nature 585, 63–67 (2020). Chaurand, P.\net al. New methodological approach for the vanadium K-edge X-ray absorption near-\nedge structure interpretation: application to the speciation of vanadium in oxide phases\nfrom steel slag. J. Phys. Chem. B 111, 5101–5110 (2007). Mansour, A. N. et al. A\ncomparative in situ X-ray absorption spectroscopy study of nanophase V O aerogel\n2 5\nand ambigel cathodes. J. Electrochem. Soc. 150, A403–A413 (2003); Giorgetti, M. et\nal. In situ X-ray absorption spectroscopy characterization of V O xerogel cathodes\n2 5\nupon lithium intercalation, J. Electrochem. Soc. 146, 2387–2392(1999).)\nTo investigate the charge compensation mechanism during Li-ion insertion and\nextraction, the normalized V K-edge XANES spectra are measured at different\nlithiation/delithiation stages, as shown in Fig.3. There are two characteristic regions in\nthe spectra: the pre-edge peaks in the range 5,465–5,470 eV, and main peaks at about\n5,484 eV. The oxidation state of the V is known to be sensitive to the intensity ratio\nbetween the pre-edge and main peaks (referred to as P/M hereafter) Among the\nstandards, pristine sample shows the highest P/M, while C H O V(V3+) presents the\n15 21 6\nlowest P/M. It is evident that the pre-edge peak of KVO is between pristine sample and\nC H O V, which indicates that the material predominantly consists of V5+ and V3+.\n15 21 6\nThe corresponding discussions have been added in lines 4-7 on page 11 in the revised\nmanuscript.\nP\nA\nG\nE\nFig.3. (a) XANES of V K-edge. (d) Valence changes of the bulk phases of V\nmeasured at different discharge/charge states.\nComment 12 of Reviewer 3:\nIn Figure 3c, the V L-edge should be clearly labeled (i.e., L1, L2 or L3).\nResponse: Thanks so much for your careful observation. We have updated Figure 3c\nto explicitly label the V L-edge features (L and L ) as requested.\n3 2\nP\nA\nG\nE\nFig.3c. V L-edge soft XAS TFY spectra.\nComment 13 of Reviewer 3:\nThe caption of the XAS figure should indicate “photon energy”, not “binding energy”\nor “photo energy”. Please correct this across the manuscript.\nResponse: Thank you so much for catching this important terminology detail. We have\nconfirmed that all instances of “binding energy” or “photo energy” in XAS figure\ncaptions (main manuscript and supplementary materials) have been updated to \"photon\nenergy\" per your instruction.\nP\nA\nG\nE\nFig. 3. Bulk and surface chemical evolution of KVO. (a, b) XANES of V K-edge. (c)\nV L-edge soft XAS TFY spectra. (d, e) Valence changes of the near-surface and bulk\nphases of V measured at different discharge/charge states. (f) O K-edge soft XAS TEY\nspectra. (g) O K-edge soft XAS TFY spectra. (h) F K-edge soft XAS TEY spectra. (i) F\nK-edge soft XAS TFY spectra. (j)The 3D views of LiF-, CO - and C HO- in the TOF-\n3 2\nSIMS sputtered volumes of SEI layers for KVO electrodes. (k)The affinity energy and\ndesolvation barrier of various electrolyte components for KVO.\nP\nA\nG\nE\nFig.S26. LEDC standard spectra.\nFig.S27. LEMC standard spectra.\nFig.S28. LiF standard spectra.\nComment 14 of Reviewer 3:\nFigure S21 reveals a pronounced amorphous phase during discharge, suggesting that a\ncrystalline-to-amorphous transition contributes to capacity. Consequently, claiming\nKVO is a “zero-strain” anode based solely on XRD refinement is misleading, as XRD\ncannot detect amorphous components.\nResponse: We thank the reviewer so much for this comment. We acknowledge the\nP\nA\nG\nE\nbroad diffraction feature observed during discharge (Fig. S21) and clarify its origin:\nThis signal arises from the formation of a disordered solid-electrolyte interphase (SEI)\non particle surfaces, not a crystalline-to-amorphous transition in the bulk KVO material.\nSEI layers are intrinsically amorphous/nanocrystalline due to their complex\ncomposition (e.g., LiF, Li CO , polymeric organics) and rapid precipitation kinetics\n2 3\nduring discharge. Their disordered structure produces the characteristic broad hump in\nXRD patterns, which is distinct from bulk amorphization. (Fig.S41) The corresponding\ndiscussions have been added in lines 5-11 on page 54 in the revised Supporting\ninformation.\nFig.S41. HRTEM images of KVO\nOur in situ optical microscopy of 100-cycle electrodes provides compelling\nevidence: KVO particles maintain exceptional structural integrity during cycling with\nminimal volume change, while graphite undergoes severe electrode-level expansion\nand fracturing. This directly demonstrates KVO’s macroscopic stability, consistent with\nits near-zero strain behavior. ( Fig.S42，Supporting Movie 2，and Supporting Movie\n3) The corresponding discussions have been added in lines 3-6 on page 17 in the revised\nmanuscript.\nFig.S42. The volume changes of (a)KVO and (b) graphite after 100 cycles of discharge\nP\nA\nG\nE\nto 0.01V under in-situ optical microscopy.\nWe have expanded the voltage window to 0.01 V for in situ XRD and confirm that\nKVO maintains its zero-strain character with unchanged diffraction peak intensities\nthroughout deep discharge, unequivocally demonstrating exceptional structural\nstability even under extreme conditions. (Fig.S43). The corresponding discussions\nhave been added in lines 3-6 on page 17 in the revised manuscript.\nFig.S43. In-situ XRD patterns of KVO/Li in-situ cell with corresponding discharging-\ncharging curves at 0.2C within 0.01−3.0 V.\nComment 15 of Reviewer 3:\nUsing EXAFS R-space peak intensity alone to infer coordination number is\ninappropriate given the potential disorder revealed during lithiation. The peak intensity\nis also influenced by the Debye-Waller factor. Quantitative EXAFS fitting is necessary\nto support structural claims.\nResponse: We thank the reviewer so much for raising this critical methodological point.\nWe fully agree that qualitative analysis of R-space peak intensity is insufficient to\ndetermine coordination environments. In response, we performed quantitative EXAFS\nfitting for both pristine, lithiated and delithiated states of KVO, and the results robustly\nconfirm the structural stability and near-zero strain characteristics of KVO.\nThe Debye-Waller factor σ2, which is closely related to the local structural disorder,\nis further discussed to get insight into the local environment associated with V and O\nframework. The value of σ2 of V-O octahedron increases from 4.54×10−3 to 5.95×\n10−3. The σ2 of the V-O polyhedron increased from 4.54 ×10−3 to 5.95×10−3, and then\nP\nA\nG\nE\ndropped back to 4.71×10−3 after the lithiation state. For the V-V coordination shell,\nthe increase of σ2 for KVO (from 3.90 × 10−3 to 4.80 × 10−3) after the first discharge\n(Figure 5f). Moreover, the σ2 value of the V-V bond for KVO (3.91 × 10−3) mostly\nrecovers to its initial value (3.90× 10−3) upon following discharge. The highly\nreversible σ2 value evolution of KVO provides clear evidence for the stability of the\nKVO framework. (Fig.S47 and Fig.S48) The corresponding discussions have been added\nin lines 17-25 on page 19 in the revised manuscript.\nFig.S49. (a-c) V K-edge FT-EXAFS spectra of KVO at different states of charge.\nP\nA\nG\nE\nFig. S50. The Debye-Waller factor σ2 of (a) V-O and (b) V-V bonds of KVO.\nComment 16 of Reviewer 3:\nPlease cite the software used for structural and spectral analyses (e.g., GSAS, Athena,\netc.).\nResponse: We thank the reviewer for this great question. As suggested, we have\nformally cited key programs mentioned in the literature. These references follow\nstandard academic formatting for clarity:\n[56] Toby, B. EXPGUI, a graphical user interface for GSAS. J. Appl. Crystallogr. 34,\n210–213 (2001).\n[57] Ravel, B. & Newville, M. Athena, Artemis"}]

Decision letter