Vacancy-ordered perovskite superlattice in cerium titanate negative electrode for enhanced lithium-ion storage

Domain: physics
Published: 2025-12-12
Editorial outcome: minor_revision

Abstract
Commercial negative electrodes such as graphite and Li4Ti5O12 are fundamental to lithium-ion batteries but face inherent trade-offs among safety, energy density, rate performance, and cycling stability. In this work, we introduce structural ordering and vacancy engineering into a perovskite negative electrode Ce2/3TiO3 to tackle this dilemma, by creating highly ordered Ce vacancies that form a stable superlattice. As a result, micron-sized Ce2/3TiO3 achieves a high specific capacity (>200 mAh g-1) at an optimal operating potential (~0.8 V vs. Li+/Li), with fast-charging capability up to 50 C and stable cycling performance exceeding 10000 cycles at 20 C. Its electrochemical performance has the potential to overcome the shortcomings of graphite and Li4Ti5O12, comparable to many representative intercalation-type negative electrodes. In situ structural analysis and atomic-scale imaging reveal a reversible topological phase transition between long-range and short-range ordering, which preserves the lattice integrity while unlocking low-barrier Li+ diffusion pathways. Here, we show that vacancy ordering provides a compelling strategy for designing high-performance electrodes.

Introduction
With the rapid advancement of portable electronics, electric vehicles, and large-scale energy grids, lithium-ion batteries (LIBs) have emerged as the leading contender for energy storage due to their high energy density, lightweight design, and reliability1–3. Next-generation LIBs are increasingly expected to operate within safe potential regimes while delivering fast-charging capability and cycling stability, essential for minimizing charging times and extending service lifespan4,5. This goal critically relies on the rational compositional design and precise structural engineering of electrode materials6–8.

Graphite and Li4Ti5O12 (LTO) stand as the commercially successful intercalation-type negative electrodes for LIBs, extensively employed in numerous practical applications9,10. Graphite offers a high specific capacity of ~372 mAh g−1 and a low potential plateau of <0.1 V (vs. Li+/Li)11,12, enabling high energy density in full cells. However, its sluggish Li+ diffusion kinetics and inevitable structural degradation during cycling severely hinder the rate capability and long-term cycling stability13,14. Moreover, the extremely low potential renders graphite susceptible to lithium dendrite formation, significantly increasing the risk of short circuits and thermal runaway15. In contrast, LTO operates at a safer potential of ~1.55 V (vs. Li+/Li) and exhibits considerable structural stability owing to its zero-strain characteristic16–18. Nevertheless, its intrinsically low electrical conductivity and sluggish Li+ diffusion lead to a limited specific capacity of only 50–60 mAh g−1 at 10 C for micron-sized particles19,20, while the relatively high potential compromises the overall energy and power densities21. To overcome these drawbacks, strategies such as defect engineering and nanostructuring have been widely adopted to enhance Li+ transport kinetics and electronic conductivity22–25. For instance, Fu et al. fabricated thin LTO nanosheets that facilitated rapid Li+ diffusion, achieving a specific capacity of 146 mAh g−1 even at 100 C26. Similarly, Su et al. designed LTO nanocrystals with intrinsic oxygen vacancies and cation redistribution, delivering 176 mAh g−1 at 10 C owing to improved ionic/electronic conductivity27. However, such nanostructured designs often compromise volumetric energy density and cycling stability, due to increased surface reactivity, structural degradation, and pronounced volumetric effects. Thus, achieving intrinsic electrochemical performance with safety, high capacity, fast-charging capability, and cycling stability at the micrometer scale remains a critical challenge in the pursuit of advanced negative electrodes.

The degree of crystallographic ordering in electrode materials directly dictates their electrochemical performance. Structural disorder and insufficient vacancy concentration often lead to sluggish ion diffusion, electrochemically inactive behavior, and progressive structural degradation, thereby compromising capacity retention, rate capability, and long-term cycling stability28,29. Disorder disrupts ion transport pathways, elevates charge-transfer resistance, and accelerates undesirable phase transitions, ultimately diminishing overall electrochemical property30. In recent years, the concept of structural ordering design has found widespread applications in superconductivity, topological electronics, catalysis, and energy storage31–35. In these systems, structural ordering not only determines the intrinsic physical properties of materials but also governs the spatial arrangement of functional units, thus optimizing key performance metrics. For battery electrode materials, structural ordering plays a crucial role in facilitating ion diffusion, reducing charge-transfer resistance, and preserving structural integrity during cycling36,37. For example, precise control of lattice-ordering domains in Ni-rich layered positive electrodes enables efficient Li+ transport and mitigates degradation from phase transitions, thus improving both high-voltage stability and cycle life38. Inspired by these advances, structural ordering design emerges as a powerful strategy to simultaneously enhance electrochemical activity, rate capability, and cycling stability. Nevertheless, vacancy engineering remains a critical yet frequently neglected aspect of negative electrodes, as it is crucial to maximizing Li+ accommodation. Therefore, integrating structural ordering with rational vacancy engineering provides a promising route to optimize the electrochemical performance of negative electrodes.

In this context, A-site deficient perovskites with the general formula Ln2/3TiO3 (Ln=La, Ce, Pr, Nd) serve as ideal structural prototypes for intercalation-type negative electrodes, intrinsically coupling ordered cation vacancies with robust TiO6 octahedral frameworks39–43. The vacancy ordering not only modulates the Ti-O-Ti connectivity to promote Li+ transport but also mitigates volume fluctuations during cycling to enhance structural stability. Several derivatives of this family, such as Li0.27La0.54TiO3, La0.5Li0.5TiO3, Li0.35Nd0.55TiO3, and Li0.38Pr0.54TiO3, have been explored as the negative electrodes for LIBs44–47. However, they usually suffer from limited rate capability and insufficient long-term cycling stability, while their underlying Li+ storage mechanisms at the atomic scale remain poorly understood. Therefore, exploring alternative A-site deficient perovskites is essential to overcome current performance constraints and to establish universal structure-property relationships for high-performance intercalation-type negative electrodes.

In this study, we integrate structural ordering and vacancy engineering within the perovskite negative electrode Ce2/3TiO3 (CTO), which possesses a superlattice structure with intrinsically ordered Ce vacancies (Fig. 1a). This vacancy-ordered superlattice endows CTO with great electrochemical performance in LIBs. Experimental results reveal that micron-sized CTO operates at a low yet safe potential of ~0.8 V (vs. Li+/Li), intermediate between graphite and LTO, while delivering a higher specific capacity of 221 mAh g−1 at 0.1 C. Furthermore, CTO exhibits high-rate capability up to 50 C and stable cycling performance exceeding 10000 cycles at 20 C, highlighting its comprehensive Li+ storage performance. To elucidate the underlying mechanism, in situ X-ray diffraction (XRD) and atomic-scale imaging were employed to directly visualize and quantify the structural evolution of CTO during Li+ insertion and extraction. First-principles calculations further confirm the low-barrier Li+ diffusion within the vacancy-ordered superlattice, consistent with its rapid electrochemical kinetics. This study proves CTO as a promising alternative to conventional graphite and LTO negative electrodes, while providing fundamental insight into the pivotal role of vacancy ordering in electrochemical performance.

Results
Structural properties of vacancy-ordered perovskite superlattice CTO

As illustrated in Fig. 1b, CTO possesses a characteristic A-site deficient perovskite structure, which consists of corner-sharing TiO6 octahedra interspersed with partially occupied Ce sites. These Ce sites are arranged in an orderly manner, forming alternating Ce-rich and Ce-poor layers with distinct occupancy ratios that collectively construct a vacancy-ordered superlattice. The pristine CTO was synthesized via a high-temperature solid-state route, yielding micron-sized particles with sizes ranging from 5 to 20 μm (Supplementary Fig. 1). Rietveld refinement of the XRD pattern (Fig. 1c) confirms that the synthesized CTO exhibits high crystallinity and phase purity, adopting a tetragonal structure with the P4/mmm space group. The refinement results (Supplementary Table 1) show that the average occupancy of the Ce1 site is 0.889, which is higher than that of the Ce2 site (0.444). Accordingly, Ce1 sites with fewer vacancies form Ce-rich layers, while Ce2 sites with more vacancies constitute Ce-poor layers. This structural configuration is consistent with the ideal model of a vacancy-ordered superlattice.

To directly visualize the vacancy-ordered superlattice structure of CTO, high-angle annular dark-field scanning transmission electron microscopy (HAADF-STEM) images with atomic resolution were acquired along the [100] zone axis. As shown in Fig. 1d, the large-scale atomic arrangement exhibits a well-defined periodic pattern with alternating bright and dark atomic columns. A magnified region aligns perfectly with the theoretical crystal structure of CTO. The brighter atomic columns correspond to Ce1 sites with higher occupancy, while the darker atoms represent Ce2 sites with a higher vacancy concentration. Smaller Ti atoms occupy positions between the Ce1 and Ce2 layers. The line intensity profile along the vertical direction further confirm the distinct contrast differences between Ce1 and Ce2 sites, consistent with the structural model of CTO. The selected area electron diffraction (SAED) pattern shows bright diffraction spots that align with the simulated results (Supplementary Fig. 2), while weak superlattice reflections in the interstitial regions indicate long-range structural modulation arising from the ordered Ce distribution. To further verify the ordered arrangement of Ce vacancies, integrated differential phase contrast scanning transmission electron microscopy (iDPC–STEM) imaging was conducted, which enhances contrasts for both heavy and light elements. As shown in Fig. 1e, the Ce-rich layers exhibit higher contrast owing to more complete Ce occupancy, whereas the Ce-poor layers display pronounced vacancy features. This alternating distribution of Ce-rich and Ce-poor layers constitutes the unique vacancy-ordered perovskite superlattice. In addition, atomic-resolution energy dispersive X-ray spectroscopy (EDX) mappings (Fig. 1f) validate the high phase purity and uniform atomic arrangement of CTO. The periodic fluctuation in Ce signal intensity suggests the different occupancy of Ce sites, and the nearly uniform Ti signal reflects the homogeneous distribution of Ti atoms. Furthermore, the electron energy loss spectroscopy (EELS) spectrum (Supplementary Fig. 3) identifies the presence of Ce3+ and Ti4+. Specifically, the Ti L2,3-edge peaks at 458.1 eV and 463.2 eV correspond to Ti4+ 2p3/2 and 2p1/2 transitions48, and the Ce M4,5-edge peaks at 880.2 eV and 897.8 eV are attributed to Ce3+ 3d5/2 and 3d3/2 transitions49,50. Collectively, these results confirm the successful synthesis of a vacancy-ordered perovskite superlattice CTO, in great agreement with the predicted crystallographic model.

Electrochemical Li+ storage performance of micron-sized CTO

Given the structural advantages of CTO in accommodating Li+, the electrochemical performance of the prepared micron-sized CTO was systematically evaluated in half-cell configurations. The initial discharge/charge cycle (Supplementary Fig. 4) delivered specific capacities of 343 and 221 mAh g−1 for Li+ insertion and extraction, respectively, corresponding to an initial coulombic efficiency of 64.4%, as corroborated by the cyclic voltammetry (CV) curves in Fig. 2a. The initial capacity loss is primarily attributed to the formation of the solid electrolyte interphase (SEI). As shown in Supplementary Fig. 5, transmission electron microscopy (TEM) images reveal progressive thickening of the SEI layer on CTO particle surfaces during the initial discharge. High-resolution Li 1s X-ray photoelectron spectroscopy (XPS) spectra indicate that the SEI consists of both organic (ROCO2Li) and inorganic (Li2CO3 and LiF) components, which are from different interfacial reactions. ROCO2Li and Li2CO3 arises from the reductive decomposition of carbonate solvents and side reactions with conductive carbon, while LiF forms through the decomposition of LiPF6 salt. As lithiation proceeds, the relative intensities of Li2CO3 and LiF signals increase while those of organic species decrease, suggesting a gradual enrichment of inorganic SEI components at low potentials and corresponding irreversible lithium consumption. After prolonged cycling, the SEI maintains a similar chemical composition but becomes slightly thicker, with further accumulation of inorganic components such as LiF and Li2CO3 (Supplementary Fig. 6). This dense, inorganic-rich SEI enhances interfacial mechanical and chemical stability, effectively suppressing continuous electrolyte decomposition and enabling long-term cycling durability. Beyond the initial cycle, subsequent CV curves at 0.2 mV s−1 exhibit nearly consistent profiles, indicating the high electrochemical reversibility of CTO. Three distinct pairs of redox peaks are observed. The pairs at 1.45 V/1.24 V and 0.81 V/0.66 V correspond to the two-step Ti4+/Ti3+ redox process46, while the pair at 0.30 V/0.10 V is attributed to the Ti3+/Ti2+ transition during deep lithiation51. These assignments are corroborated by ex situ XPS analysis (Supplementary Fig. 7), where the Ce 3d spectra remain unchanged from the pristine state to 1.0 V and 0.01 V, confirming the electrochemical inertness of Ce3+. In contrast, the Ti 2p spectra show a progressive shift of Ti4+ peaks toward lower binding energies, indicating the stepwise reduction of Ti4+ to Ti3+ and subsequently to Ti2+ upon deep lithiation.

Compared with commercial intercalation-type negative electrodes, graphite and LTO, CTO demonstrates improved operating potential and specific capacity (Fig. 2b). Specifically, CTO operates at a safer potential of ~0.8 V (vs. Li+/Li) than that of graphite, which effectively mitigates the risk of lithium dendrite formation. In addition, CTO delivers a higher reversible specific capacity of 220 mAh g−1 at 0.1 C than that of LTO, thereby enabling a higher overall energy density. Moreover, CTO exhibits high-rate charge/discharge performance, maintaining high specific capacities even at rates up to 50 C (Fig. 2c). As the current rate increases from 0.1 C to 50 C, the reversible capacities gradually decrease from 220 to 218, 206, 192, 169, 148, 124, and 81 mAh g−1, corresponding to retention ratios of 99.1, 93.6, 87.3, 76.8, 67.3, 56.4, and 36.8%, respectively. Upon returning to 0.1 C, the reversible capacity recovers to its initial value (Fig. 2d), confirming the fast-charging capability and stability of CTO. When benchmarked against intercalated-type negative electrodes such as graphite, LTO, La0.5Li0.5TiO3, etc.45,52–57, CTO demonstrates both high specific capacities and improved rate performance (Fig. 2f and Supplementary Table 2). In terms of cycling performance, CTO exhibits long-term durability at both low and high rates. As plotted in Fig. 2e, CTO retains a specific capacity of 184 mAh g−1 after 500 cycles at 1 C, corresponding to a retention of 90.6%. Even at higher rates of 10 C and 20 C, the capacity retention remains above 80% after extended cycles (Fig. 2g and Supplementary Figs. 8 and 9). These results indicate the robustness of the CTO lattice and its ability to sustain long-term cycling stability, showing performance comparable to other intercalation-type negative electrodes for LIBs (Supplementary Fig. 10 and Supplementary Table 3).

Many fast-charging negative electrodes reported to date rely on thin electrodes with low mass loadings (1-2 mg cm−2), which limit their practical applicability. To evaluate the practical potential of CTO, we examined thick electrodes with a high mass loading of ~10.2 mg cm−2. As shown in Supplementary Fig. 11, the CTO electrode delivers specific capacities of 193, 173, 160, 146, 117, 89, 63, and 30 mAh g−1 at current rates from 0.1 C to 50 C, demonstrating great rate capability even under high-loading conditions. Moreover, it retains 80% of its initial capacity after 1000 cycles at 20 C, indicating considerable cycling stability. Beyond gravimetric performance, we also assessed the volumetric capacity of CTO. At a low mass loading of 1.5 mg cm−2, CTO exhibits a maximum volumetric capacity of 550 mAh cm−3, and still achieves 394 mAh cm−3 at the high loading of 10.2 mg cm−2 (Supplementary Fig. 12). This performance is comparable to that of previously reported micron-sized or nano-structured intercalation-type negative electrodes (Supplementary Fig. 12c), underscoring the practical advantages of CTO for high-energy-density and fast-charging LIBs.

Low-temperature performance is also a critical metric for practical battery applications, so the electrochemical performance of CTO was evaluated at 0 °C and −15 °C (Supplementary Figs. 13 and 14). At 0 °C, CTO exhibits only a moderate capacity reduction compared to 25 °C, still delivering a reversible capacity of 183 mAh g−1 at 0.1 C. Moreover, it maintains a great rate capability, delivering 170, 161, 149, 130, 111, 89, and 54 mAh g−1 as the rate increases from 0.5 C to 50 C. Long-term cycling tests further demonstrate durable stability, with negligible capacity decay over 1000 cycles at 10 C and over 10000 cycles at 20 C. Even at a lower temperature of −15 °C, where Li+ diffusion kinetics are further hindered, CTO retains considerable cycling stability at 5 C and 10 C. The low-temperature performance of CTO is comparable to conventional commercial negative electrodes such as LTO and graphite58–60, highlighting its promising potential for practical applications.

Furthermore, the practical performance of CTO was evaluated in a full-cell configuration (CTO||LiFePO4). The comparison of charge/discharge curves shows that the operating potential difference between LiFePO4 and CTO reaches 2.76 V (Supplementary Fig. 15), which is favorable for a high output energy density. The full cell delivers a high discharge capacity of 150 mAh g−1 at 0.1 C, with a first-cycle coulombic efficiency of 62.5% (Supplementary Fig. 16a), comparable to that of the half cell. Upon increasing the current rate to 0.5, 1, 2, 5, and 10 C, the full cell maintains specific capacities of 125, 115, 108, 100, and 92 mAh g−1, respectively (Supplementary Fig. 16b). The capacity retention at 10 C reaches 61.3%, demonstrating reliable rate capability. Furthermore, long-term cycling at 5 C confirms electrochemical stability, with 80% capacity retention after 1000 cycles (Supplementary Fig. 16c). Thus, the practical potential of CTO is proven, combining its high specific capacity, fast-charging capability, and cycling stability in a full-cell configuration.

Electrochemical kinetic analysis for Li+ diffusion

To elucidate the underlying mechanisms responsible for the high-rate performance, electrochemical kinetic analysis was performed to investigate Li+ diffusion behavior. As shown in Fig. 3a, the CV curves retain their characteristic shape with progressively enhanced peak currents and small polarization as the scan rate increases from 0.2 mV s−1 to 0.4, 0.7, and 1.1 mV s−1, indicating high electrochemical reversibility and rapid reaction kinetics. Based on the relationship between scan rate and peak current, the b-values were calculated to evaluate the reaction kinetics (Supplementary Note 1). The obtained b-values for the cathodic and anodic peaks are 0.78 and 0.77 (Fig. 3b), respectively, corresponding to Li+ insertion and extraction processes. The b-values between 0.5 and 1.0 suggest kinetic behavior controlled by a combination of both diffusion and surface-capacitive mechanisms. To quantify these contributions, a current-scan rate separation model was employed (Supplementary Note 2). The analysis reveals high pseudocapacitive contribution ratios in the micron-sized CTO, accounting for 72, 76, 82, and 86% of the total capacity at 0.2, 0.4, 0.7, and 1.1 mV s−1, respectively (Fig. 3c and Supplementary Fig. 17). These pseudocapacitive behaviors are comparable to those of many nanoscale electrodes61–63, which facilitate fast charge storage and mitigate structural degradation caused by deep Li+ diffusion.

The Li+ diffusion coefficients in CTO were further determined through a galvanostatic intermittent titration technique (GITT). As depicted in Fig. 3d, alternating current pulses and relaxation steps were performed at a low rate of 0.1 C, aiming to probe Li+ diffusion during lithiation and delithiation under near-equilibrium conditions. As shown in the enlarged view (Fig. 3e), ΔEτ denotes the potential change during a single current pulse, while ΔEs represents the potential difference between two adjacent equilibrium states. Based on Fick’s law of diffusion (Supplementary Note 3), the Li+ diffusion coefficients were calculated as a function of potential. The average diffusion coefficients during lithiation and delithiation are 5.29×10−11 and 9.39×10−11 cm2 s−1, respectively, which are comparable to those of many fast-charging electrodes (Supplementary Table 4). These results confirm the rapid Li+ diffusion within the micron-sized CTO structure. The Li+ diffusion coefficient exhibits a sharp decline near ~1.25 V during both lithiation and delithiation, followed by an increase below this potential (Fig. 3f). This behavior suggests a possible phase transition, accompanied by accelerated Li+ diffusion kinetics in the post-transition phase.

In situ XRD analysis of the sequential lithiation process

To gain insight into the sequential lithiation process in CTO, in situ XRD analysis was conducted in tandem with electrochemical discharge/charge measurements. As depicted in Fig. 4a, b, the evolution of in situ XRD patterns was synchronized with the discharge/charge curves at 0.5 C, providing real-time information on structural changes. In the pristine state, CTO exhibited three distinct diffraction peaks at 25.6°, 32.6°, and 40.2°, corresponding to the (101), (110), and (112) planes of the tetragonal structure (Supplementary Fig. 18). As lithiation proceeded, these diffraction peaks gradually shifted to lower angles, reflecting lattice expansion driven by Li+ insertion. A transition in the diffraction pattern occurred near 1.25 V, where the peaks shifted more rapidly toward lower angles until full lithiation was reached. No new peaks, peak splitting, or disappearance was observed, suggesting that the transformation around 1.25 V involved a topological phase transition. Upon delithiation, the diffraction peaks returned to their original positions, demonstrating the considerable structural reversibility of CTO. XRD Rietveld refinement below 1.25 V corroborated the retention of the tetragonal system, but with a reduction in the unit-cell parameters c that approached the value of parameter a (Supplementary Fig. 19). This structure was thus identified as a pseudo-cubic phase, exhibiting a similar structural symmetry to the cubic phase. Raman spectra collected across the phase transition region (Supplementary Fig. 20) reveal systematic vibrational changes during lithiation. The gradual downshift of main peaks, along with the progressive merging of low-frequency modes, indicates an increased structural symmetry, consistent with the transition toward a pseudo-cubic phase with higher symmetry. The refinement structural parameters displayed abrupt changes in the unit-cell parameters a, c, and lattice volume V at the transition point, confirming a topological phase transition between tetragonal and pseudo-cubic phases (Fig. 4c, d). Despite the noticeable unit-cell volume variation during lithiation, the overall periodicity of the original tetragonal structure was preserved. This topological phase transition underscores the structural resilience of CTO, which is favorable for achieving long-term cycling stability.

Moreover, the unit-cell parameters a and c displayed gradual variations during the early stage of Li+ insertion, followed by a rapid expansion after the transition point at 1.25 V. This indicates that the post-transition phase provides a more favorable structural environment for Li+ diffusion and insertion, in agreement with the calculated Li+ diffusion coefficients. Based on the discharge curve and the above results (Fig. 4e), we propose a possible sequence for the lithiation process. Initially, Li+ ions preferentially insert into the Ce-poor layers, which possess a higher vacancy concentration that facilitates Li+ diffusion. As the Ce-poor layers become partially filled, the CTO lattice undergoes a topological phase transition, generating the pseudo-cubic phase. Ultimately, Li+ ions occupy the residual vacancies in both Ce-poor and Ce-rich layers, resulting in the complete transformation of CTO into the pseudo-cubic structure.

Atomic-scale structural evolution during Li+ insertion

Due to the limitations of in situ XRD analysis in resolving topological structural information, atomic-resolution scanning transmission electron microscopy (STEM) was performed to directly visualize the structural evolution of CTO at various lithiation states. As shown in Fig. 5a–c, HAADF-STEM images were acquired along the [110] zone axis for three representative states: pristine, discharged to 1.25 V, and fully discharged to 0.01 V, corresponding to the initial, phase transition, and fully lithiated states, respectively. In the pristine state, CTO exhibits a well-defined vacancy-ordered superlattice, characterized by alternating Ce-rich and Ce-poor layers with pronounced contrast differences (Fig. 5a). Upon discharge to 1.25 V, a two-phase coexistence becomes evident (Supplementary Fig. 21), marked by different colors (Fig. 5b). As the vacancies in Ce-poor layers are partially occupied by Li+, the contrast difference between Ce-poor and Ce-rich layers diminishes. Consequently, a topological phase emerges, characterized by nearly uniform contrast across all Ce sites, indicative of a pseudo-cubic structure. When fully discharged to 0.01 V, Li+ ions occupy all available vacancies in both Ce-rich and Ce-poor layers, leading to a complete structural transformation of the superlattice into a pseudo-cubic phase (Fig. 5c). The SAED pattern along [100] zone axis after transformation shows additional weak satellite reflections surrounding the primary diffraction spots, which is different from those in the pristine state, proving the formation of a short-range ordered superlattice (Supplementary Fig. 22)64,65. Since the pseudo-cubic structure retains similar periodicity to the original vacancy-ordered superlattice but exhibits a halving of the lattice parameter along the c-axis, this structural transformation is identified as a topological phase transition. This Li+-induced transformation represents a shift from long-range to short-range ordering, as schematically illustrated in Fig. 5d–f. Upon charging back to 3.0 V, the structure reverts to its original vacancy-ordered superlattice (Supplementary Fig. 23), consistent with the in situ XRD analysis. Even after prolonged cycling, the CTO maintains its morphological and structural integrity, with large-scale vacancy-ordered superlattices interspersed with minor pseudo-cubic domains (Supplementary Figs. 24 and 25), demonstrating the high reversibility and stability of the topological phase transition in CTO.

The chemical states of the constituent elements in CTO at different lithiation states were probed through EELS. As shown in Supplementary Fig. 26, the Ti L2,3-edge signals progressively shift to lower energies during lithiation, confirming the valence reduction of Ti4+ caused by Li+ insertion. In contrast, the Ce M4,5-edge peaks remain essentially unchanged (Supplementary Fig. 27), suggesting the electrochemical stability of Ce cations. This chemical inertness of Ce helps mitigate lattice distortions associated with the redox activity of the Ti, thereby enhancing structural robustness. To further analyze lattice strain, two-dimensional strain distribution maps were reconstructed for different lithiation states using geometric phase analysis (GPA)66,67. As depicted in Fig. 5g, the color scale in the strain maps and the full width at half maxima (FWHM) of the histograms represent the distribution and magnitude of lattice strain. The pristine superlattice exhibits a homogeneous and low-strain state, whereas the lithiated lattice becomes more distorted with higher strain values. Especially, the lattice strain at 1.25 V shows the highest strain heterogeneity due to the coexistence of two phases with large lattice mismatches, yet the average strain remains as low as +0.4%. The results confirm the low-strain structural characteristic of CTO, consistent with its long-term cycling stability.

To further elucidate the evolution of structural periodicity, fast Fourier transform (FFT) and inverse FFT (IFFT) analyses were performed for pristine and lithiated CTO (Fig. 5h). In the pristine state, the IFFT image shows a long-range homogenous atomic arrangement, with no additional diffraction spots beyond those of CTO in the FFT pattern. In contrast, the FFT pattern of lithiated CTO reveals extra diffraction spots within the specific intervals, confirming the emergence of additional ordering. The corresponding IFFT image further proves the presence of short-range ordered domains. Similar structural modulations are observed along the [100] zone axis (Supplementary Figs. 28 and 29), originating from the topological transition to the pseudo-cubic phase. Pair distribution function (PDF) analysis was performed to compare the pristine and lithiated states. As shown in Supplementary Fig. 30a, the PDF refinements reveal a transition from long-range tetragonal ordering to short-range pseudo-cubic ordering, consistent with the structural models in Supplementary Fig. 30b. In the short-range region (3.1-6.1 Å), the intensities of Ce1-Ti and Ce1-Ce2 correlations increase, and a new Ce-O correlation emerges at ~5.0 Å upon lithiation (Supplementary Fig. 30c), indicating rearrangement of the Ce coordination environment. Moreover, the Ce1-Ce1, Ce-O, and Ce1-Ce2 peaks become noticeably narrower after lithiation (Supplementary Fig. 30d), reflecting enhanced short-range ordering of Ce atoms. These results collectively demonstrate that lithiation triggers a topological transition characterized by strengthened short-range structural ordering around the Ce coordination environment. To elucidate the effect of the transition from long-range to short-range ordering, atomic-resolution iDPC-STEM imaging were performed to derive the electric field distribution and charge density distributions (Supplementary Fig. 31)68,69. As shown in Fig. 5i, the pristine lattice displays alternating vacancy-rich and vacancy-poor layers, with distinct electric field and charge density distributions between these layers, as evidenced by the DPC-STEM and charge density maps. After lithiation, the lattice evolves into a uniform atomic arrangement with nearly identical electric fields and charge densities at all Ce sites. This direct atomic-level visualization of structure, electric field, and charge density provides clear evidence of a pronounced topological phase transition from long-range to short-range ordering induced by Li+ insertion. These findings align well with the above analyses, offering a comprehensive understanding of the structural evolution in CTO during lithiation.

Vacancy-ordered superlattice enabled Li+ storage mechanism

To elucidate the Li+ storage mechanism in CTO, density functional theory (DFT) calculations based on first principles were performed. A vacancy-ordered tetragonal superlattice model (Ce10Ti16O48) was constructed to simulate the lithiation behavior, which closely approximates the experimentally observed framework of CTO. Seven types of Li+ insertion sites (Li1–Li7) were identified (Supplementary Fig. 32). Specifically, Li1 and Li2 correspond to Ce vacancies in the Ce-poor layers, while Li3 resides in the Ce-rich layers. In addition, Li4–Li7 occupy O4 windows between adjacent TiO6 octahedra, consistent with the theoretical capacity and experimental observations (Supplementary Fig. 33). Due to the high concentration of ordered vacancies, all predicted sites exhibit low formation energies (Supplementary Table 5), suggesting the structural openness and favorable kinetics of the tetragonal superlattice for Li+ accommodation. Among them, Li1 presents the lowest formation energy, indicating it as the favorable initial insertion site. Once Li1 is occupied, the Ce-poor and Ce-rich layers become structurally equivalent in both atomic ratios and positions, thereby enhancing the structural symmetry of the tetragonal superlattice and driving a topological phase transition toward a pseudo-cubic structure (Fig. 6a).

To investigate the lithiation behavior in the transformed phase, a partially lithiated pseudo-cubic model (Li5Ce8Ti24O72) was constructed, in which five types of Li+ insertion sites (Li8–Li12) were identified (Supplementary Fig. 34). Li8 and Li9 correspond to unoccupied Ce vacancies in the Ce-poor and Ce-rich layers, respectively, while Li10–Li12 occupy the O4 windows between adjacent TiO6 octahedra. These insertion sites exhibit comparable formation energies (Supplementary Table 6), with Ce-vacancy sites being energetically more favorable, confirming their priority in the lithiation sequence. These results predict a sequential lithiation process from the initial tetragonal phase to a partially lithiated pseudo-cubic phase, and eventually to the fully lithiated state (Fig. 6a), consistent with the experimental observations. Concurrently, the electronic structure undergoes evolution during lithiation, as evidenced by the projected density of states (DOS). The pristine tetragonal phase exhibits a wide band gap of ~1.8 eV, with the conduction band far from the Fermi level (Fig. 6b and Supplementary Data 1–3). Upon Li+ insertion, the band gap gradually narrows, accompanied by an increase in electronic states near the Fermi level. The projected DOS of Ti orbitals (Supplementary Fig. 35) reveals electron transfer from Li to Ti 3d orbitals, resulting in a progressive transition from spin-symmetric to spin-asymmetric Ti 3d occupancy, indicative of the stepwise reduction of Ti4+. Moreover, the topological phase transition induces charge redistribution. Projected charge density maps along the (100) plane (Fig. 6c) reveal that localized charge accumulation appears at former Ce-vacancy sites after phase transition, in agreement with the DPC-STEM observations (Fig. 5i). This charge reorganization enhances both electronic conductivity and Li+ mobility in the pseudo-cubic phase, suggesting a strong correlation between structural evolution and electrochemical properties.

To further quantify Li+ diffusion in both the vacancy-ordered tetragonal superlattice and the pseudo-cubic structure, climbing image nudged elastic band (CI-NEB) calculations were carried out. Benefiting from the open channels, bond valence sum (BVS) energy maps (Fig. 6d) show continuous Li+ diffusion pathways connecting adjacent TiO6 octahedra and Ce vacancies in both structures. Two dominant Li+ diffusion pathways were identified, including interlayer migration between Ce-poor and Ce-rich layers, and intralayer diffusion within the Ce-poor layer. CI-NEB calculations (Fig. 6d, e) show low Li+ migration barriers along the interlayer pathways for both phases, with the pseudo-cubic phase exhibiting a lower energy barrier (0.47 eV) than that of the tetragonal phase (0.62 eV). Similarly, the intralayer diffusion barrier within the Ce-poor layer is reduced in the pseudo-cubic phase (0.50 eV) compared to the tetragonal superlattice (0.64 eV) (Supplementary Figs. 36 and 37), indicating that the topological phase transition accelerates Li+ diffusion kinetics. Overall, the DFT results reaffirm the advantages of the vacancy-ordered superlattice in promoting efficient Li+ storage and diffusion. The lithiation-induced topological transformation to the pseudo-cubic phase further enhances Li+ diffusion kinetics, thereby underpinning the great electrochemical performance of CTO, including high specific capacity, fast-charging capability, and long-term cycling stability.

Methods
Preparation of CTO

Micron-sized CTO was synthesized via a two-step high-temperature solid-state reaction. Initially, high-purity CeO2 (≥99.0%, Aladdin) and TiO2 (≥99.0%, Aladdin) powders with the stoichiometric ratio were weighed and mixed. The mixture was then subjected to high-energy ball milling (SPEX 8000M) for 10 h at 300 rpm using agate milling balls. Then, the resulting mixture was calcined at 1000 °C for 5 h in air to obtain a white intermediate product. In the subsequent step, the intermediate was sintered at 1200 °C for 10 h under a 5% H2/Ar atmosphere, with a heating rate of 10 °C min−1. After natural cooling, brownish-black CTO particles were successfully obtained.

Structural and physical characterizations

Powder XRD data were collected using a high-resolution X-ray diffractometer (D8 Advance, Bruker), with a scan rate of 3° min−1 and a step size of 0.01°. PDF measurements were performed on the same X-ray diffractometer equipped with an Mo-Kα radiation source (λ=0.7107 Å) operating at 50 kV and 30 mA. The crystal structures were constructed and visualized using the VESTA software package70. Electrochemically synchronized in situ XRD measurements were conducted on the same X-ray diffractometer integrated with an in situ battery unit and an external electrochemical control system. Scanning electron microscopy (SEM) images were collected using a field-emission scanning electron microscope (Hitachi S-4800, Japan) operated at an accelerating voltage of 5 kV. Raman spectra were collected using a confocal Raman spectrometer (Renishaw inVia) equipped with a 633 nm excitation laser. XPS spectra were acquired on an X-ray photoelectron spectrometer (Thermo Scientific ESCALAB 250Xi). Samples were transferred to the analysis chamber using an airtight container to avoid surface contamination. The binding energies were calibrated using the C 1s peak at 284.8 eV as a reference. HAADF-STEM imaging was carried out using a spherical aberration-corrected transmission electron microscope (Spectra 300, Thermo Fisher Scientific) operated at 300 kV. Atomic-resolution EDX elemental mapping was acquired by a high-resolution energy detector. EELS data were collected through a Gatan imaging filter detector, with an energy resolution of 0.8 eV for the zero-loss peak and an energy dispersion of 0.2 eV per channel. The iDPC-STEM images were generated by processing the segregated signals from four-quadrant annular detectors, while DPC-STEM images were extracted to visualize the electric-field color vectors. Color-coded atomic contrast heatmaps were reconstructed by a custom Python script, which extracted atomic column positions and intensities from the raw HAADF-STEM images. Two-dimensional lattice strain maps and statistical strain analyses were calculated based on the GPA method to quantify atomic column displacements relative to ideal positions. All STEM samples were prepared via a plasma-focused ion beam (FIB, Helios 5 PFIB DualBeam, Thermo Fisher Scientific), ensuring a specimen thickness below 100 nm. All FIB procedures were conducted following standardized protocols to ensure consistency and reproducibility. For ex situ characterizations, all cells were measured at 25 ± 1 °C and then transferred into an argon-filled glovebox (H2O < 0.1 ppm, O2 < 0.1 ppm, Mikrouna) for disassembly. The electrodes were sealed in airtight containers for subsequent structural and compositional analyses.

Electrochemical measurements

Electrochemical performance was evaluated using CR2016-type coin cells, with microporous polyolefin membranes (25 μm thick, Celgard 2325, Shenzhen Kejing) serving as separators. Electrode sheets were prepared following a standard slurry-casting process. The electrode slurry consisted of active material, carbon black (Super P, Shenzhen Kejing), and polyvinylidene fluoride (PVDF, Shenzhen Kejing) in a mass ratio of 8:1:1, dispersed in N-methyl-2-pyrrolidinone (NMP, ≥99.0%, Sinopharm Chemical Reagent) and stirred for 24 h. The homogeneous slurry was uniformly coated on one side of copper foils (≥99.9%, 15 μm thick, Shenzhen Kejing) using a doctor blade. The coated electrodes were dried under vacuum at 80 °C for 10 h and subsequently punched into 12-mm diameter disks using a slicer. The average mass loadings were ~1.5 mg cm−2 and ~10.2 mg cm−2 for thin and thick electrodes, respectively. For half-cell configurations, commercial lithium metal foil (1.2 mm thick, 14 mm diameter, China Energy Lithium Co., Ltd.) was used as the counter electrode, while LiFePO4 (≥99.0%, Shenzhen Kejing) electrode (active material: carbon black: PVDF=8:1:1) with a mass loading of ~2.0 mg cm−2 was employed as the positive electrode in full cells. The electrolyte (Canrd Technology Co., Ltd.) was composed of 1 M LiPF6 dissolved in a 1: 1: 1 (v/v/v) mixture of dimethyl carbonate (DMC), ethylene carbonate (EC), and diethylene carbonate (DEC). The electrolyte volume used for each cell assembly was ~100 μL.

CV measurements were conducted using an electrochemical workstation (CHI660E, CH Instruments), while galvanostatic charge-discharge performance was assessed with a battery testing system (CT-3008, Neware). The electrochemical window for half-cells was set as 0.01-3.0 V, and a rate of 1 C was defined as a specific current of 250 mAh g−1. GITT measurements were performed at a specific current of 25 mAh g−1 with a current pulse duration of 10 min followed by a relaxation period of 90 min, with data acquired at a rate of one point per second. Electrochemical performance tests of half-cells were carried out in a constant-temperature chamber at controlled temperatures of 25 ± 1 °C, 0 ± 1 °C, and −15 ± 1 °C. For full cells, the mass ratio of active materials in the negative and positive electrodes was optimized with an N/P ratio of ~1.1, and the electrochemical window was set as 0.8–3.2 V. The specific capacity of full cells was calculated based on the mass of the positive electrode. All electrochemical performance tests were conducted using at least three independent cells to ensure data reliability.

DFT calculations

DFT calculations were conducted using the Vienna Ab initio Simulation Package (VASP) to investigate the electronic and structural properties of CTO. The generalized gradient approximation (GGA) within the Perdew-Burke-Ernzerhof (PBE) functional was employed to accurately describe the exchange-correlation interactions. To closely approximate the experimentally observed structures, a vacancy-ordered tetragonal superlattice (Ce10Ti16O48) and a lithiation-induced pseudo-cubic phase (Li5Ce8Ti24O72) were reconstructed. The corresponding k-point meshes were set to 4×4×2 and 4×4×4, respectively, to ensure adequate Brillouin zone sampling during structural relaxation. The plane-wave cutoff energy was set as 500 eV, and a force convergence criterion of 0.01 eV Å−1 was applied to ensure robust geometric convergence. To account for the strong electron correlation effects in Ce and Ti atoms, the Hubbard U correction was incorporated into the calculations. The on-site Coulomb interaction parameters were set as 5.0 eV for Ce f-states and 3.0 eV for Ti d-states. Possible Li+ insertion sites were systematically identified by calculating the formation energies of different lithiated configurations. The formation energy for each Li+-inserted configuration was calculated as follows:

Eformation = [ELi_x-pristine − Epristine − xELi] (1)

Where ELi_x-pristine and Epristine are the total energies of the lithiated and pristine structures, respectively, ELi denotes the energy per atom of bulk body-centered cubic lithium, and x is the number of inserted Li+ ions. To further assess Li+ migration kinetics, the CI-NEB method was employed to determine the transition states and migration energy barriers. The CI-NEB calculations were conducted using the same cutoff energy (500 eV) and force convergence criterion (0.01 eV Å−1) as in geometric optimizations.

Peer review rounds
[{"round": 1, "reviewer_comments": {"reviewer_1": "This manuscript reports a new anode material for lithium-ion batteries with a perovskite structure. The electrochemical\nperformance, kinetic properties, and energy storage mechanism were investigated in detail. A dynamic\ntransformation between long-range and short-range ordering was claimed to contribute\nto the superior lithium storage performance.\nHowever, the correlation between such structural transformation and the lithium storage properties is not unequivocally\nelucidated. Moreover, as the full cells show low Coulombic efficiency, the practical\napplication potential of the anode is largely reduced. On top of that, the scientific\nfindings in this manuscript are not enough to meet the high standards of Nature\nCommunications. Below are some specific points:\n1. In Figure 5b, the boundary between two phases is not clear.\n2. Apart from electron microscopy, which only shows localized information, more experimental evidence should be provided\nto highlight the properties of such short-range ordering.\n3. What is the driving force behind the transformation between long-range and short-range ordering?\n4. The theoretical calculation part clearly only considered the ion storage in the long-range ordered crystal structure. The\ninfluence of the short-range ordering is overlooked. Moreover, the formation energy of the three states does not show much\ndifference.\nRegarding the electrochemical properties:\n5. Please specify the redox reactions in the CV curves.\n6.The specific capacity exhibits severe fluctuation at high current rates, please provide reasonable explanation for this.\n7. In Figure 3e, it is quite obvious that the potential has not reached an equilibrium state. Therefore, the accuracy of the\ndiffusion coefficient calculation is questionable.", "reviewer_2": "I co-reviewed this manuscript with one of the reviewers who provided the listed reports. This is part of the Nature\nCommunications initiative to facilitate training in peer review and to provide appropriate recognition for Early Career\nResearchers who co-review manuscripts.", "reviewer_3": "In this work, the authors report an excellent electrochemical Li+ storage performance using the vacancy-ordered perovskite\nCe TiO as the anode materials for half and full cells at ambient temperature, compared to other perovskite and\n2/3 3\nintercalation-type anodes. Especially, the moderate operating voltage (~0.8 V) avoids the safety hazard of graphite due to\ngenerating lithium dendrites. Moreover, the authors perform an exhaustive characterization for the crystal structure and\ncharge-discharge process. The results are interesting and valuable. However, there exist some critical problems in the\ncurrent version, especially about the charge-discharge mechanism. Therefore, I can’t recommend its publication on Nature\nCommunications unless the following issues are satisfactorily resolved.\n1.A-vacancy Ln TiO is a type of perovskites that have been widely researched in other fields (Journal of Materials\n2/3 3\nScience Letters, 1999, 18, 1181; Journal of Solid State Chemistry, 2000, 149, 354; J. Am. Ceram. Soc., 2002, 85 1339;\nCeramics International, 2016, 42, 1698), including Ce TiO (Journal of Alloys and Compounds, 2000, 305, 72). The\n2/3 3\nauthors should introduce the characters of Ln TiO and the considerations that Ce TiO is selected in this work.\n2/3 3 2/3 3\n2.The authors emphasize the superiority of CTO with respect to the commercial anodes. Nowadays, a good low-temperature\nperformance has become an urgent demand for commercial secondary batteries. In order to better satisfy the practical\nrequirement, the low-temperature performance of CTO should be examined.\n3.The authors state that the intercalated Li+ occupy the A-site vacancies of CTO in the whole lithiation process. However, I\ncannot fully agree with this view. Ce TiO has the A-vacancy concentration of only 1/3. The resultant theoretical capacity is\n2/3 3\nonly ~47 mAh/g when all A vacancies are filled. Actually, it has been reported that the inserted Li+ may occupy the O4\nwindow positions in perovskite anodes (Chem Mater 2013, 25, 1607). If this case holds true in this work, it should be\nemphatically considered in the related experimental characterization and theoretical analysis.\n4.In the pristine CTO sample, Ce and Ce layers uniformly alternate. And, the authors conclude that the intercalated\n0.889 0.444\nLi+ preferentially occupy the A-vacancies in the Ce layers. So, the distribution of inserted Li+ should be relatively\n0.444\nuniform in all Ce layers. And the lithiation-induced pseudo-cubic phase should homogeneously mixes with the pristine\n0.444\ntetragonal phase, as observed in Fig. 5g (middle 1.25 V). Why does a distinct boundary between the pure tetragonal and\npure pseudo-cubic phases occur, as shown in Fig. 5b and Fig. S15? Moreover, the difference between blue and orange\nregions in Fig. S15 is hardly distinguished, how to determine the phase boundary?\n5.The authors state that the valence configurations of Ti atoms are 3d04s0 in pristine structure, and 3d14s0 in lithiated\nstructures. If so, why are remarkably asymmetric for the spin-up and spin-down PDOS of Ti-d and Ti-p in pristine structure\n(Fig. S22)? And why are symmetric or quasi-symmetric for the spin-resolved PDOS of Ti 3d14s0 under lithiations?\n6.An obviously contradictory result. The DOS results of pristine model are inconsistent in Fig. 6b and Fig. 22. For example,\nthe former shows a wide bandgap about 2 eV, whereas the latter shows the nature of a half metal.\n7.The constructed pristine CTO model is Ce Ti O in Fig. 6a, not the Ce Ti O described in DFT calculations, Methods. If\n3 8 24 4 8 24\nthe constructed supercell model is Ce Ti O , the A-vacancy concentration is up to 5/8, much higher than the experimental\n3 8 24\n1/3 in Ce TiO . In addition, the constructed pristine model is a quasi-cubic supercell, different from the synthesized\n2/3 3\ntetragonal cell (c≈2a). How do authors ensure the reliability of DFT calculations?\n8.In Fig S8, the full cell has a coulombic efficiency only ~30% for the first cycle, much lower than the 64.4% of half cell, why?\n9.Line 313, “the valance reduction” may be “the valence reduction”.", "reviewer_4": "Review:\nIn this manuscript, vacancy-ordered Ce2/3TiO3 (CTO) is reported as a perovskite-type anode with high gravimetric capacity\n(> 200 mAh g⁻¹ at ≈ 0.8 V), ultrafast-rate capability (≈ 80 mA h g-1 at 50 C), and sustained cycling stability (90% capacity\nretention after 10,000 cycles). Accelerating Li+ transport by A-site vacancy ordering would be a new route which, if\nconfirmed under electrode loadings, can be much better than graphite and spinel Li4Ti5O12. The proof of concept is\nachieved utilizing significant measurements, however need to refine some sections as: electrochemical access is restricted\nto thin (≈ 1.5 mg cm-2) electrodes without any volumetric benchmarking; the 10,000 cycle claim rests simply upon a single\nend-point value; and the \"pseudo-cubic\" lithiated phase, as well as potential Ce-valence changes. Consequently, I\nrecommend minor revision with an emphasis of the need for extensive electrode level validation, detailed cycle dataset,\nquantitative structural analysis and direct comparison with the industry benchmarks.\nFirst, the reversible capacity of 220–221 mA h g-1 at 0.1 C surpasses the Ti4+ → Ti3+ theoretical limit for CTO (≈ 142 mA h\ng-1), while EELS shows no Ce valence shift, implying that the extra ≈ 80 mA h g-1 likely arises from over-estimated active\nmass or parasitic side reactions rather than intrinsic redox. Second, the initial Coulombic efficiency is only 64 %, yet the\nsolid-electrolyte interphase is uncharacterised, leaving the source of irreversible capacity loss- and its consequences for\nlong-term stability—unclear. Third, the purported reversible transition from tetragonal to pseudo-cubic symmetry is supported\nonly by in-situ XRD peak shifts and qualitative STEM images; rigorous Rietveld refinement of the lithiated phase and Ce-\nvalence analysis are required to validate this structural evolution. Finally, complementary diagnostics—such as operando or\nex-situ Raman spectroscopy near the 1.25 V transition and post-cycling XPS or FTIR of the electrodes- are needed to\ncorroborate the symmetry change and identify SEI constituents responsible for the low first-cycle efficiency."}, "author_response": "Response Letter to Reviewers\nResponses to the comments from Reviewer 1\nGeneral comments: This manuscript reports a new anode material for lithium-ion\nbatteries with a perovskite structure. The electrochemical performance, kinetic\nproperties, and energy storage mechanism were investigated in detail. A dynamic\ntransformation between long-range and short-range ordering was claimed to contribute\nto the superior lithium storage performance. However, the correlation between such\nstructural transformation and the lithium storage properties is not unequivocally\nelucidated. Moreover, as the full cells show low Coulombic efficiency, the practical\napplication potential of the anode is largely reduced. On top of that, the scientific\nfindings in this manuscript are not enough to meet the high standards of Nature\nCommunications.\nResponse: We sincerely thank the reviewer’s careful evaluation and constructive\nfeedback. We appreciate the positive remarks on our comprehensive investigation of\nthe electrochemical performance, kinetics, and energy storage mechanism of CTO. We\nalso acknowledge the reviewer’s concerns and have addressed the specific points as\nfollows:\n(1) Correlation between structural transformation and Li storage properties:\nIn response, we have significantly reinforced this correlation through both\nexperimental and theoretical evidence. Experimentally, we supplemented pair\ndistribution function (PDF) analysis of pristine and lithiated CTO. The results (now\npresented in Supplementary Fig. 29) further reveal a clear transformation from a long-\nrange ordered tetragonal phase to a short-range ordered pseudo-cubic structure upon\nlithiation. These results are detailed in our response to Comment 2. Theoretically, we\nhave rebuilt our DFT models to reflect both the vacancy-ordered tetragonal phase and\nthe lithiation-induced pseudo-cubic phase. By systematically comparing lithium-\ninsertion sites, electronic structures and Li+ diffusion pathways of the two phases, we\ndemonstrate that the structural transformation not only enhances electronic activity but\nalso facilitates more favorable Li+ transport kinetics. The revised DFT results are\nincluded in revised Fig. 6 and Supplementary Figs. 31–36.\nCollectively, these experimental and computational results establish a coherent\nmechanism whereby lithium insertion modulates lattice symmetry and electronic\nstructure, thereby improving lithium storage and diffusion kinetics.\n(2) Low first-cycle coulombic efficiency of full cells:\nThrough detailed analysis, we identified that this issue primarily stemmed from\nthe suboptimal N/P ratio design in previous full-cell configurations. We sincerely\napologize for this oversight and have optimized the capacity balance (N/P = 1.1-1.2).\nThe optimized full cells show a significantly improved first-cycle coulombic efficiency\nof 62.5%, approaching that of the half cells (64.4%). Furthermore, the rate capability\nand cycling performance of the optimized full cells were greatly enhanced. It delivered\na capacity retention of 61.3% at 10 C and maintained 80% of its initial capacity after\n1000 cycles at 5 C, demonstrating the excellent electrochemical performance of the full\ncells and the practical application potential of CTO. The updated results are shown in\nSupplementary Fig. 15, and the relevant discussion has been added to the revised\nmanuscript as follows:\nLines 255-262, Page 10: At 0.1 C, the full cell delivers a high discharge capacity of\n150 mAh g-1, with a first-cycle coulombic efficiency of 62.5% (Supplementary Fig.\n15a), comparable to that of the half cell. As the rate increases to 0.5, 1, 2, 5, and 10 C,\nthe full cell maintains specific capacities of 125, 115, 108, 100, and 92 mAh g-1,\nrespectively (Supplementary Fig. 15b). Notably, the capacity retention at 10 C reaches\n61.3%, demonstrating excellent rate capability. Furthermore, long-term cycling at 5 C\nconfirms outstanding stability, with 80% capacity retention after 1000 cycles\n(Supplementary Fig. 15c).\nSupplementary Fig. 15 Electrochemical performance of LiFePO ||CTO full cells. a\n62.5%. b Charge/discharge profiles at different current rates from 0.1 C to 10 C. c Rate\ncapability and long-term cycling performance at 5 C.\n(3) Scientific findings and novelty:\nTo the best of our knowledge, this study represents the first report of a vacancy-\nordered perovskite anode CTO that simultaneously achieves a safe voltage (~0.8 V),\nhigh reversible capacity (>200 mAh g-1), ultrafast rate capability (up to 50 C), and\nexceptional cycling stability (10000 cycles). Moreover, we uncover a novel mechanism\nabout the topological phase transition between long-range and short-range structural\nordering. The electrochemical performance of CTO is outstanding compared to many\nstate-of-the-art intercalation-type anodes, and the mechanism has not been explored in\nprior studies. Additionally, we have conducted full-cell validation, detailed theoretical\ncalculations, and advanced atomic-scale characterizations to support these findings. We\nbelieve that this work provides new insights and design paradigm for the development\nof high-performance intercalation-type anode materials.\nWe hope that the revisions and the point-by-point responses below satisfactorily\naddress the reviewer’s concerns. Thank you again for your insightful comments.\nComment 1: In Figure 5b, the boundary between two phases is not clear.\nResponse: We thank the reviewer for raising this important point. The phase boundary\nbetween the tetragonal and pseudo-cubic phases is primarily distinguished by the\ncontrast difference between Ce1 (Ce-rich) and Ce2 (Ce-poor) layers in the HAADF-\nSTEM image. In the tetragonal phase, these alternating layers show a distinct contrast\ndifference, whereas this contrast becomes nearly uniform in the pseudo-cubic phase.\nTo more clearly visualize the phase boundary, we developed a Python-based image\nprocessing algorithm to extract atomic column intensities from the raw HAADF-STEM\nimage and reconstruct a color-coded atomic contrast heatmap. This approach enhances\natomic contrast visualization and enables clear identification of the phase boundary.\nIn response to the reviewer’s comment, we have added the original HAADF-\nSTEM image and the corresponding atomic contrast heatmap to Supplementary Fig.\n20.\nSupplementary Fig. 20 a HAADF-STEM image at 1.25 V. b Reconstructed atomic\ncontrast heatmap derived from the raw HAADF-STEM image. The phase boundary\nbetween the tetragonal and pseudo-cubic phases is marked by white dash lines.\nComment 2: Apart from electron microscopy, which only shows localized information,\nmore experimental evidence should be provided to highlight the properties of such\nshort-range ordering.\nResponse: We appreciate the reviewer’s valuable suggestion. To further substantiate\nthe short-range ordering, we have supplemented pair distribution function (PDF) \nanalysis on both pristine and lithiated CTO samples, as PDF provides statistically\naveraged structural information beyond localized electron microscopy.\nThe PDF refinement results show that the pristine structure is well-fitted by a long-\nrange ordered tetragonal model, while the lithiated state is matched with a short-range\nordered pseudo-cubic structure, consistent with the phase evolution observed in\nHAADF-STEM images. In the short-range region (3.1 - 6.1 Å), the lithiated sample\nexhibits enhanced peak intensities corresponding to Ce1-Ti and Ce1-Ce2 correlations,\nalong with the emergence of a new Ce-O peak at ~5.0 Å, indicating a rearrangement of\nthe Ce coordination environment upon lithiation. Moreover, the Ce1-Ce1, Ce-O, and\nCe1-Ce2 peaks show notably narrower widths after lithiation, suggesting enhanced\nshort-range ordering of Ce atoms in the lithiated phase. Together with the electron\nmicroscopy results, these findings provide robust evidence for the formation of a well-\ndefined short-range ordered structure induced by lithium insertion.\nThe PDF results have been included in Supplementary Fig. 29, and the\ncorresponding discussion has been added to the revised manuscript as follows:\nLines 402-414, Page 15: Pair distribution function (PDF) analysis was performed to\ncompare the pristine and lithiated states. As shown in Supplementary Fig. 29a, the\nPDF refinements reveal a transition from long-range tetragonal ordering to short-range\npseudo-cubic ordering, consistent with the structural models in Supplementary Fig.\n29b. In the short-range region (3.1 - 6.1 Å), the intensities of Ce1-Ti and Ce1-Ce2\ncorrelations increase, and a new Ce-O correlation emerges at ~5.0 Å upon lithiation\n(Supplementary Fig. 29c), indicating a rearrangement of the Ce coordination\nenvironment. Moreover, the Ce1-Ce1, Ce-O, and Ce1-Ce2 peaks become noticeably\nnarrower after lithiation (Supplementary Fig. 29d), reflecting enhanced short-range\nordering of Ce atoms. These results collectively demonstrate that lithiation triggers a\ntopological transition characterized by enhanced short-range structural ordering around\nCe coordination environment.\nSupplementary Fig. 29 PDF analysis revealing lithiation-induced transformation from\nlong-range tetragonal ordering to short-range pseudo-cubic ordering. a PDF\nrefinements of pristine and lithiated CTO. b Structural models of the long-range ordered\ntetragonal and short-range ordered pseudo-cubic phases. c Comparison of PDF peaks\nin the short-range region (3.1 - 6.1 Å), where the three dominant peaks are assigned to\nCe1-Ti1 (~3.4 Å), Ce1-Ce1 (~3.8 Å), and Ce1-Ce2 (~5.5 Å) correlations, additional\nminor peaks are attributed to Ti-O and Ce-O distances. d Comparison of Ce1-Ce1, Ce-\nO, and Ce1-Ce2 peaks.\nComment 3: What is the driving force behind the transformation between long-range\nand short-range ordering?\nResponse: We appreciate the reviewer’s insightful question. Although direct\ninvestigations on the driving force for structural transformation in CTO are limited,\nanalogous system La Li TiO (LLTO) with similar vacancy-ordered perovskite\n2/3-x x 3\nframeworks have been extensively studied and provide valuable insight. In LLTO, it\nhas been well established that structural symmetry and ordering are strongly modulated\nby Li content and A-site vacancy concentration. Specifically, prior studies (e.g., Solid\nState Ionics, 2000, 134, 219-228; Ionics, 2005, 11, 333-342; Angew. Chem. Int. Ed.,\n2000, 39, 619-621; Chem. Mater., 2005, 17, 2404–2412; Chem. Mater., 2024, 36,\n1197–1213) have shown that low Li content with high La vacancy concentration\nstabilizes a long-range ordered orthorhombic structure, while increasing Li content and\nreducing La vacancy concentration drive a transition to higher-symmetry tetragonal and\npseudo-cubic structures.\nBy analogy, our DFT calculations reveal that Li+ ions preferentially occupy A-site\nvacancies within the Ce-poor layers during early lithiation. This occupancy reduces the\natomic ratio difference between Ce-rich and Ce-poor layers, thereby leading to a\nsymmetry change from tetragonal to pseudo-cubic, along with a transition from long-\nrange to short-range ordering. The inserted Li⁺ compensates for the A-site vacancies in\nthe tetragonal superlattice and disrupts the long-range periodicity of Ce vacancies,\nresulting in a short-range ordered structure. We also calculated the Gibbs free energy\nchange (ΔG) of the transition from tetragonal to pseudo-cubic structure, which shows\na negative ΔG of -1.23 eV (as shown in Fig. R1), indicating that the transformation is\nthermodynamically favorable.\nSimilar lithiation-induced transition from tetragonal to pseudo-cubic phase has\nalso been reported in other perovskite oxides, such as La Li TiO (Nat. Commun.,\n0.5 0.5 3\n2020, 11, 3490) and Li Pr TiO (Chem. Eng. J., 2024, 479, 147765), where the\n0.38 0.54 3\ntransformation was likewise attributed to Li+ occupation of A-site vacancies and the\nresulting increase in lattice symmetry.\nTherefore, we think that the driving force behind the transition from long-range to\nshort-range ordering in CTO arises from lithium insertion-induced reduction of A-site\nvacancy concentration and the associated enhancement of lattice symmetry.\nFig. R1: Calculated Gibbs free energy change (ΔG) for the structural transformation\nfrom tetragonal to pseudo-cubic phase.\nComment 4: The theoretical calculation part clearly only considered the ion storage in\nthe long-range ordered crystal structure. The influence of the short-range ordering is\noverlooked. Moreover, the formation energy of the three states does not show much\ndifference.\nResponse: We thank the reviewer for this insightful comment. In the revised\nmanuscript, we have substantially updated the theoretical calculations to more\naccurately reflect the lithiation behavior in both long-range and short-range ordered\nphases.\nSpecifically, we reconstructed two representative structural models: (1) a vacancy-\nordered tetragonal superlattice Ce Ti O , corresponding to the pristine long-range\n10 16 48\nordered phase, and (2) a partial lithiation-induced pseudo-cubic structure LiCe Ti O ,\n5 8 24\nrepresenting the short-range ordered phase. We systematically calculated the Li+\ninsertion energies at various crystallographic sites in both models (Supplementary\nFigs. 31 and 33). The revised calculations also show that the formation energies for\ndifferent insertion sites within the same phase are comparable (Supplementary Tables\n3 and 4), suggesting that multiple sites are energetically favorable for lithium\naccommodation. This result indicates that the open framework and abundant vacancies\nin both phases provide numerous low-barrier sites for Li+ insertion, thereby enabling\nhigh lithium storage capacity and fast diffusion kinetics. These findings align well with\nthe experimentally observed high specific capacity and excellent rate capability of CTO.\nWe have updated Fig. 6 and the corresponding discussion (Lines 428-485, Pages\n16-18) in the revised manuscript.\nSupplementary Fig. 31 Optimized structure models depicting potential Li+ storage\nsites in the tetragonal phase.\nSupplementary Fig. 33 Optimized structure models depicting potential Li+ storage\nsites in the pseudo-cubic phase.\nSupplementary Table 3 Formation energies of different Li+ storage sites in tetragonal\nphase.\nLi+ storage sites Free energy (eV) Formation energy (eV)\nPristine -605.49906117 0.00000000\nLi -611.85515242 -4.45219713\ncubic phase.\nLi+ storage sites Free energy (eV) Formation energy (eV)\nPristine -308.48216030 0.00000000\nLi -311.98629707 -1.60024265\nResponse: We sincerely thank the reviewer for this valuable suggestion. In response,\nwe have carefully analyzed the CV curves in Fig. 2a and now provide a detailed\nassignment of the observed redox features.\nAs shown in Fig. 2a, three distinct redox couples are identified at approximately\n1.45 V/1.24 V, 0.81 V/0.66 V, and 0.30 V/0.10 V. The redox couples at 1.45 V/1.24 V\nand 0.81 V/0.66 V correspond to a two-step Ti4+/Ti3+ redox process, consistent with\nprevious reports on Ti-based anodes (e.g., Nano Energy, 2024, 119, 109065; Chem. Eng.\nJ., 2024, 479, 147765; Electrochem. Commun., 2013, 32, 5-8; Adv. Energy Mater., 2023,\n13, 2302015). The redox couple at 0.30 V/0.10 V is attributed to the Ti3+/Ti2+ redox\ntransition, which usually occurs under deep lithiation below 0.5 V, as reported in other\nTi-based anodes (e.g., Energy Environ. Sci., 2017, 10, 1456-1464; Adv. Sci., 2025, 12,\n2410543; Nano Energy, 2022, 94, 106972; ACS Appl. Mater. Interfaces, 2024, 16,\n898–906).\nThis redox assignment is further supported by the evolution of the Ti L -edge\n2,3\nEELS spectra (Supplementary Fig. 25), which exhibit progressive shifts of the Ti L\n2\nand L peaks to lower energies upon discharge, indicating the stepwise reduction of Ti4+\n3\ntoward lower valence states. Additionally, the supplemented ex situ XPS measurements\n(Supplementary Fig. 6) confirm this evolution, showing a sequential reduction of Ti4+\nto Ti3+ and finally to Ti2+ during lithiation.\nWe have added the relevant discussion to the revised manuscript as follows:\nLines 188-196, Pages 7-8: In addition, the CV curves exhibit three distinct pairs of\nredox peaks. The pairs located at 1.45 V/1.24 V and 0.81 V/0.66 V correspond to the\ntwo-step Ti4+/Ti3+ redox process46, while the pair at 0.30 V/0.10 V is attributed to the\nTi3+/Ti2+ transition during deep lithiation51. These assignments are further supported by\nex situ XPS analysis (Supplementary Fig. 6). It reveals that the Ce 3d spectra remain\nunchanged from the pristine state to 1.0 V and 0.01 V, confirming the electrochemical\ninactivity of Ce. In contrast, the Ti 2p spectra show a progressive shift of Ti4+ signal\npeaks toward lower binding energies, indicating the step reduction of Ti4+ to Ti3+ and\nsubsequently to Ti2+ upon deep lithiation.\nReferences:\n46. Liu H., et al. Cation-deficient perovskite Li Nd TiO as a high-performance\n0.35 0.55 3\nanode for lithium-ion batteries. Nano Energy 119, 109065 (2024).\n51. Jang M., et al. Two Steps Li Ion Storage Mechanism in Ruddlesden–Popper\nLi La Ti O . Adv. Sci., 12, 2410543 (2025).\n2 2 3 10\nSupplementary Fig. 6 XPS spectra of CTO at the pristine state, after discharging to\n1.0 V, and after discharging to 0.01 V. a Evolution of Ce 3d high-resolution spectra. b\nEvolution of Ti 2p high-resolution spectra.\nComment 6: The specific capacity exhibits severe fluctuation at high current rates,\nplease provide reasonable explanation for this.\nResponse: We thank the reviewer for raising this important point. To the best of our\nknowledge, similar capacity fluctuation phenomena during high-rate cycling have been\nobserved in many other fast-charging anode materials (e.g., Adv. Mater., 2020, 32,\n1905295; Adv. Mater., 2022, 34, 2107262; Adv. Energy Mater., 2023, 13, 2302107; Adv.\nEnergy Mater., 2022, 12, 2102972). According to the explanations in prior studies (Adv.\nMater., 2024, 36, 2412266; Nano Energy, 2024, 119, 109065; J. Energy Chem., 2022,\n74, 34-44), the initial capacity increase is typically attributed to gradual electrolyte\ninfiltration and the progressive activation of electrochemically accessible regions\nwithin the bulk electrode. Subsequent capacity decline is commonly associated with\nSEI growth, electrode passivation, or structural degradation.\nIn our case, multiple capacity fluctuations are observed during extended high-rate\ncycling. We speculate that it originates from the delayed activation of interior active\nregions. Specifically, due to the micron-scale size of CTO particles and the large current\ndensities applied, lithium diffusion during the early cycles is insufficient to fully access\nthe particle interior. The rapid lithiation/delithiation at high rates limits lithium\npenetration depth, restricting full utilization of accessible sites. As cycling proceeds,\nthe accumulated mechanical stress and repeated volume changes induce particle\nfragmentation or local SEI disruption, thereby exposing previously isolated regions.\nThis reactivation process allows further utilization of internal intercalation sites,\nleading to repeated capacity recovery during prolonged cycling. Despite these capacity\nfluctuations, the CTO electrode maintains high capacity retention (> 80%) after\nthousands of cycles, demonstrating its excellent long-term cycling stability and\npractical potential for high-rate applications.\nComment 7: In Figure 3e, it is quite obvious that the potential has not reached an\nequilibrium state. Therefore, the accuracy of the diffusion coefficient calculation is\nquestionable.\nResponse: We thank the reviewer for this valuable comment. To address this concern,\nwe repeated the GITT measurements using an extended relaxation time of 90 minutes\n(instead of the original 20 minutes). This adjustment ensures that the potential\nsufficiently approaches a near-equilibrium state during each titration step, thereby\nimproving the accuracy of the calculated diffusion coefficients.\nThe updated results confirm that CTO exhibits fast Li+ diffusion kinetics, with\naverage diffusion coefficients of 5.29×10−11 cm2 s-1 during lithiation and 9.39×\n10−11 cm2 s-1 during delithiation. These values are close to the previously calculated\nresults, reaffirming the intrinsically rapid Li+ transport properties of CTO.\nAccordingly, we have updated Fig. 3d-f with the revised GITT profiles and\ndiffusion coefficient plots. The updated average diffusion coefficients have also been\nincorporated into the revised manuscript and Supplementary Table 2.\nLines 294-297, Page 11: The average diffusion coefficients during lithiation and\ndelithiation were found to be 5.29×10−11 and 9.39×10−11 cm2 s−1, respectively,\nwhich exceed those of many fast-charging electrode materials (Supplementary Table\n2).\nFig. 3 Kinetic analysis for Li⁺ diffusion within CTO electrode. d GITT curves during\nlithiation and delithiation at a rate of 0.1 C. e Enlarged view of GITT curves for the\ndetermination of ∆𝐸 and ∆𝐸 . f Calculated Li⁺ diffusion coefficients during both\n𝜏 𝑠\nlithiation and delithiation processes.\nResponses to the comments from Reviewer 2\nReviewer 2\nGeneral comments: I co-reviewed this manuscript with one of the reviewers who\nprovided the listed reports. This is part of the Nature Communications initiative to\nfacilitate training in peer review and to provide appropriate recognition for Early Career\nResearchers who co-review manuscripts.\nResponse: We sincerely thank you for your constructive feedback and valuable\nsuggestions, as well as the time and effort you invested in the thorough evaluation of\nour manuscript.\nResponses to the comments from Reviewer 3\nReviewer 3\nGeneral comments: In this work, the authors report an excellent electrochemical Li+\nstorage performance using the vacancy-ordered perovskite Ce TiO as the anode\n2/3 3\nmaterials for half and full cells at ambient temperature, compared to other perovskite\nand intercalation-type anodes. Especially, the moderate operating voltage (~0.8 V)\navoids the safety hazard of graphite due to generating lithium dendrites. Moreover, the\nauthors perform an exhaustive characterization for the crystal structure and charge-\ndischarge process. The results are interesting and valuable. However, there exist some\ncritical problems in the current version, especially about the charge-discharge\nmechanism. Therefore, I can’t recommend its publication on Nature Communications\nunless the following issues are satisfactorily resolved.\nResponse: We sincerely thank the reviewer for the positive evaluation of our work. We\ngreatly appreciate the reviewer’s recognition of the study’s novelty, the electrochemical\nadvantages of CTO, and the comprehensive characterizations. At the same time, we\nfully acknowledge the reviewer’s concerns regarding the charge-discharge mechanism\nand agree that these issues are critical for validating our conclusions.\nIn response, we have accordingly revised the manuscript to address all the\ncomments. We hope that the point-by-point responses and newly added data\nsatisfactorily address the reviewer’s concerns. The detailed responses to each comment\nare provided below.\nComment 1: A-vacancy Ln TiO is a type of perovskites that have been widely\n2/3 3\nresearched in other fields (Journal of Materials Science Letters, 1999, 18, 1181; Journal\nof Solid State Chemistry, 2000, 149, 354; J. Am. Ceram. Soc., 2002, 85 1339; Ceramics\nInternational, 2016, 42, 1698), including Ce TiO (Journal of Alloys and Compounds,\n2/3 3\n2000, 305, 72). The authors should introduce the characters of Ln TiO and the\n2/3 3\nconsiderations that Ce TiO is selected in this work.\n2/3 3\nResponse: We sincerely thank the reviewer for this constructive suggestion.\nA-site deficient perovskites with the general formula Ln TiO (Ln = La, Ce, Pr,\n2/.3 3\nNd) adopt an ordered double-perovskite structure characterized by one-third A-site\nvacancies periodically arranged along the c-axis. This vacancy ordering stabilizes the\nTiO octahedral framework and preserves long-range periodicity. Such structural\n6\nfeatures facilitate ion transport and accommodate reversible lattice changes during\ncycling. The intrinsic structural superiority renders A-site deficient perovskites\npromising for intercalation-type anode materials, which require both efficient ion\nmobility and robust structural stability. Although the electrical transport, dielectric, and\nmagnetic properties of Ln TiO have been widely studied, their potential for lithium-\n2/.3 3\nion storage and the underlying microscopic mechanisms remain largely unexplored.\nTo the best of our knowledge, among Ln TiO family, the electrochemical\n2/.3 3\nproperties and associated lithium-storage mechanisms of Ce TiO have never been\n2/3 3\npreviously reported. In contrast, derivative anodes based on other Ln TiO , such as\n2/.3 3\nLi La TiO (Electrochem. Commun., 2013, 32, 5-8), La Li TiO (Nat.\n0.27 0.54 2.945 0.5 0.5 3\nCommun., 2020, 11, 3490), Li Nd TiO (Nano Energy, 2024, 119, 109065), and\n0.35 0.55 3\nLi Pr TiO (Chem. Eng. J., 2024, 479, 147765), have been studied, but these anodes\n0.38 0.54 3\ngenerally exhibit limited rate capability and insufficient long-term cycling stability\ncompared to Ce TiO reported in this study. Furthermore, previous studies on\n2/3 3\nLn TiO anodes have largely focused on electrochemical property, with limited\n2/.3 3\ninsight into the underlying structural mechanisms, leaving the structure-property\nrelationship in A-site deficient perovskites insufficiently understood. Therefore, the\nselection of Ce TiO in this study aims not only to achieve a breakthrough in\n2/3 3\nelectrochemical performance but also to elucidate a deeper understanding of structural\nmechanisms, thereby offering a new design paradigm for high-performance\nintercalation-type anode materials.\nThe relevant description and references have been added to the introduction\nsection of the revised manuscript as follows:\nLines 86-98, Page 4: In this context, A-site deficient perovskites with the general\nformula Ln TiO (Ln = La, Ce, Pr, Nd) represent an ideal structural prototype for\n2/3 3\nintercalation-type anodes, intrinsically combining ordered cation vacancies with a\nrobust TiO octahedral framework39, 40, 41, 42, 43. This vacancy ordering not only\n6\nmodulates the Ti-O-Ti connectivity to promote Li+ transport but also mitigates volume\nchanges during cycling to enhance structural stability, which are critical for high-\nperformance anode materials. Several derivatives of this family, including\nLi La TiO , La Li TiO , Li Nd TiO , and Li Pr TiO , have been\n0.27 0.54 2.945 0.5 0.5 3 0.35 0.55 3 0.38 0.54 3\nexplored as the anodes for LIBs44, 45, 46, 47. However, they usually suffer from limited\nrate capability and insufficient long-term cycling stability, while their underlying\nlithium-storage mechanisms at the microscopic level remain poorly understood.\nTherefore, exploring novel A-site deficient perovskite is essential to overcome current\nperformance limitations and to establish universal structure-property relationships for\nhigh-performance intercalation-type anodes.\nReferences:\n39. Jung W. H. Structure, thermopower and electrical transport properties of La TiO .\n2/3 3−δ\nJ. Mater. Sci. Lett. 18, 1181-1183 (1999).\n40. Yoshii K. Synthesis and Magnetic Properties of Ln TiO (Ln=Pr and Nd). J. Solid\n2/3 3\nState Chem. 149, 354-359 (2000).\n41. Yoshioka H. Structure and Electrical Properties of A-site-Deficient Perovskite\nCompounds in the La TiO -La NbO System. J. Am. Ceram. Soc. 85, 1339-1342\n2/3 3 1/3 3\n(2002).\n42. Bugrov A. N., et al. Soft chemistry synthesis and dielectric properties of A-site\ndeficient perovskite-type compound La TiO . Ceram. Int. 42, 1698-1704 (2016).\n2/3 3−δ\n43. Yoshii K. Structural and magnetic studies of the lanthanide deficient perovskite\nCe TiO . J. Alloys Compd. 305, 72-75 (2000).\n2/3 3\n44. Hua C., Fang X., Wang Z., Chen L. Lithium storage in perovskite lithium\nlanthanum titanate. Electrochem. Commun. 32, 5-8 (2013).\n45. Zhang L., et al. Lithium lanthanum titanate perovskite as an anode for lithium ion\nbatteries. Nat. Commun. 11, 3490 (2020).\n46. Liu H., et al. Cation-deficient perovskite Li Nd TiO as a high-performance\n0.35 0.55 3\nanode for lithium-ion batteries. Nano Energy 119, 109065 (2024).\n47. Liu H., et al. A-site deficient perovskite lithium praseodymium titanate as a high-\nrate anode for lithium-ion batteries. Chem. Eng. J. 479, 147765 (2024).\nComment 2: The authors emphasize the superiority of CTO with respect to the\ncommercial anodes. Nowadays, a good low-temperature performance has become an\nurgent demand for commercial secondary batteries. In order to better satisfy the\npractical requirement, the low-temperature performance of CTO should be examined.\nResponse: We appreciate the reviewer’s insightful suggestion.\nIn response, we have performed additional electrochemical measurements at 0 ℃\nand -15 ℃ to evaluate the low-temperature performance of the CTO electrode. The\nresults show that, although a capacity reduction is observed at low temperatures, CTO\nretains great rate capability and stable cycling performance. The low-temperature\nperformance of CTO is better than that of conventional commercial anodes such as\nLi Ti O and graphite, further highlighting its promising potential for practical\n4 5 12\napplications.\nThe results have been added to Supplementary Figs. 12 and 13, and the relevant\ndiscussion has been included in the revised manuscript as follows:\nLines 237-248, Page 9: Low-temperature performance is also a critical metric for\npractical battery applications, so the electrochemical performance of CTO was\nevaluated at 0 and -15 ℃, (Supplementary Figs. 12 and 13). At 0 ℃, although a\nmoderate reduction in capacity compared to room temperature, CTO still delivers a\nreversible capacity of 183 mAh g-1 at 0.1 C. More notably, it maintains excellent rate\ncapability, delivering 170, 161, 149, 130, 111, 89, and 54 mAh g-1 as the rate increases\nfrom 0.5 C to 50 C. Long-term cycling tests further demonstrate outstanding stability,\nwith negligible capacity decay over 1000 cycles at 10 C and over 10000 cycles at 20 C.\nEven at lower temperature of -15 ℃, where Li+ diffusion kinetics are more severely\nlimited, it can maintain considerable cycling stability at 5 C and 10 C. The low-\ntemperature performance of CTO is better than conventional commercial anodes such\nas Li Ti O and graphite58, 59, 60, underscoring its promising potential for practical\n4 5 12\napplications in extreme environments.\nReferences:\n58. Allen J. L., Jow T. R., Wolfenstine J. Low temperature performance of nanophase\nLi Ti O . J. Power Sources 159, 1340-1345 (2006).\n4 5 12\n59. Ho C.-K., Li C.-Y. V., Deng Z., Chan K.-Y., Yung H., Yang C. Hierarchical\nmacropore-mesoporous shell carbon dispersed with Li Ti O for excellent high rate\n4 5 12\nsub-freezing Li-ion battery performance. Carbon 145, 614-621 (2019).\n60. Hou R., et al. An Unexpected Low-Temperature Battery Formation Technology\nEnabling Fast-Charging Graphite Anodes. Adv. Funct. Mater. 35, 2500481 (2025).\nSupplementary Fig. 12 Low-temperature electrochemical performance of CTO\nelectrode at 0 ℃ a Charge-discharge profiles at various current rates from 0.1 C to 50\nC. b Rate capability and corresponding coulombic efficiency at 0.1, 0.5, 1, 2, 5, 10, 20,\nand 50 C. c Cycling performance at 10 C. d Cycling performance at 20 C.\nSupplementary Fig. 13 Low-temperature electrochemical performance of CTO\nelectrode at -15 ℃ a Charge-discharge profiles at various current rates from 0.1 C to\nand 20 C. c Cycling performance at 5 C. d Cycling performance at 10 C\nComment 3: The authors state that the intercalated Li+ occupy the A-site vacancies of\nCTO in the whole lithiation process. However, I cannot fully agree with this view.\nCe TiO has the A-vacancy concentration of only 1/3. The resultant theoretical\n2/3 3\ncapacity is only ~47 mAh/g when all A vacancies are filled. Actually, it has been\nreported that the inserted Li+ may occupy the O4 window positions in perovskite anodes\n(Chem Mater 2013, 25, 1607). If this case holds true in this work, it should be\nemphatically considered in the related experimental characterization and theoretical\nanalysis.\nResponse: We sincerely thank the reviewer for this valuable comment. We fully agree\nthat the relatively low A-site vacancy concentration (1/3 per formula unit) in Ce TiO\n2/3 3\ncannot account for the high observed capacity, and that additional lithium insertion sites\nbeyond the A-site vacancies must be considered.\nTo address this concern, we have thoroughly revised the theoretical calculations.\nIn the updated DFT calculations, we systematically identified all potential lithium\ninsertion sites in both the vacancy-ordered tetragonal Ce Ti O and the pseudo-cubic\n10 16 48\nLiCe Ti O structures (as shown in Supplementary Figs. 31 and 33). In addition to\n5 8 24\nA-site vacancies, a series of energetically favorable interstitial sites were identified in\nthe O4 windows between adjacent TiO octahedra (i.e., sites Li - Li and Li - Li ),\n6 4 7 10 12\nconsistent with the prior literature (e.g., Chem. Mater., 2013, 25, 1607). The calculated\nformation energies for these O4 window sites are comparable to those of A-site\nvacancies (as shown in Supplementary Tables 3 and 4), indicating there are\nthermodynamically accessible and contribute significantly to the overall lithiation\nprocess.\nTo further confirm the theoretical predictions, we supplemented atomic-resolution\nintegrated differential phase contrast scanning transmission electron microscopy\n(iDPC-STEM) on lithiated CTO, which enables simultaneous imaging of light and\nheavy elements with sub-angstrom resolution. The iDPC-STEM image reveals\nadditional contrast at the O4 window sites, confirming lithium occupation at these\ninterstitial positions (Supplementary Fig. 32).\nOverall, both theoretical and experimental results demonstrate that Li+ insertion in\nCTO occurs not only at A-site vacancies but also extensively at interstitial O4 window\nsites between the TiO octahedra, thereby providing a high lithium storage capacity.\nsites in the tetragonal phase.\nSupplementary Fig. 32 iDPC-STEM image and corresponding structural model of\nlithiated CTO along the [110] zone axis, highlighting Li+ insertion into O4 windows\nbetween adjacent TiO octahedra. The inserted Li+ positions are indicated by red dashed\n6\ncircles.\nSupplementary Fig. 33 Optimized structure models depicting potential Li+ storage\nsites in the pseudo-cubic phase.\nComment 4: In the pristine CTO sample, Ce0.889 and Ce0.444 layers uniformly\nalternate. And, the authors conclude that the intercalated Li+ preferentially occupy the\nA-vacancies in the Ce0.444 layers. So, the distribution of inserted Li+ should be\nrelatively uniform in all Ce0.444 layers. And the lithiation-induced pseudo-cubic phase\nshould homogeneously mixes with the pristine tetragonal phase, as observed in Fig. 5g\n(middle 1.25 V). Why does a distinct boundary between the pure tetragonal and\npure pseudo-cubic phases occur, as shown in Fig. 5b and Fig. S15? Moreover, the\ndifference between blue and orange regions in Fig. S15 is hardly distinguished, how to\ndetermine the phase boundary?\nResponse: We appreciate the reviewer’s thoughtful question regarding the phase\nboundary and the interpretation of Fig. S15. Our response is organized into two parts\nto address the reviewer’s concerns in detail:\n(1) Regarding the distinct phase boundary between tetragonal and pseudo-cubic\nphases:\nIn principle, Li+ ions preferentially intercalate into the Ce-poor (Ce ) layers,\n0.444\nwhich could ideally result in a homogeneous mixture of the lithiation-induced pseudo-\ncubic phase and the pristine tetragonal phase. However, under practical experimental\nconditions, lithium intercalation is rarely spatially uniform due to kinetic factors such\nas diffusion limitations, particle surface effects, and local structural variations. Given\nmicron-scale size of CTO particles, lithiation typically initiates at the particle surface\nand propagates inward, leading to a spatially resolved phase-front propagation process.\nThe pseudo-cubic phase nucleates at the particle surface and gradually extends into the\nbulk as lithiation proceeds. Additionally, the sample preparation by focus ion beam (FIB)\nslicing exposes the cross-section of particles, making the internal phase boundaries\naccessible for observation. As a result, a distinct phase boundary between the tetragonal\nand pseudo-cubic phases was observed at the intermediate state via STEM imaging.\nAnalogous phase-front progression and two-phase coexistence behaviors have been\nwidely observed in other electrodes (e.g., LiNi Co Al O , LiFePO , Fe O , TiO ,\n0.8 0.15 0.05 2 4 3 4 2\netc.), as revealed by in situ or ex situ characterizations (e.g., ACS Energy Lett., 2020, 5,\n2098–2105; Science, 2016, 353, 566-571; ACS Appl. Energy Mater., 2020, 3, 8009–\n8016; Nature Commun., 2016, 7, 11441; Adv. Mater., 2017, 29, 1700236).\n(2) Regarding the visibility and identification of the phase boundary:\nWe apologize for the confusion caused by the original HAADF-STEM image in\nFig. S15. Although the phase regions were colored-labeled according to structural\nassignments, the contrast difference may appear subtle in grayscale rendering. To\naddress this, we reconstructed a color-coded atomic contrast heatmap from the raw\nHAADF-STEM image (now shown in Supplementary Fig. 24). This method maps\natomic column intensities to color gradients, which clearly visualizes the periodic Z-\ncontrast difference between Ce-rich and Ce-poor layers in the tetragonal phase. In\ncontrast, this contrast variation is significantly diminished in the pseudo-cubic region,\nenabling clear identification of the phase boundary.\nSupplementary Fig. 24 a HAADF-STEM image along the [100] zone axis of CTO\nafter long-term cycling. The boundary between tetragonal and pseudo-cubic phases is\nindicated by white dash lines. b Reconstructed atomic contrast heatmap derived from\nthe raw HAADF-STEM image, highlighting the coexistence of tetragonal and pseudo-\ncubic phases. c Line intensity profile along the blue line in the tetragonal region,\nshowing pronounced periodic modulation of Ce atomic column intensities. d Line\nintensity profile along the orange line in the pseudo-cubic phase, where the contrast\nvariation of Ce atomic columns is largely diminished.\nComment 5: The authors state that the valence configurations of Ti atoms are 3d04s0\nin pristine structure, and 3d14s0 in lithiated structures. If so, why are remarkably\nasymmetric for the spin-up and spin-down PDOS of Ti-d and Ti-p in pristine structure\n(Fig. S22)? And why are symmetric or quasi-symmetric for the spin-resolved PDOS of\nTi 3d14s0 under lithiations?\nResponse: We thank the reviewer for pointing out this important issue. We sincerely\napologize for the confusion caused by the previously presented PDOS results, which\narose from the inappropriate structural models in the initial calculations.\nTo address this issue, we reconstructed two representative structural models that\nmore accurately reflect the experimentally observed phases: (1) a tetragonal vacancy-\nordered Ce Ti O superlattice representing the pristine state, and (2) a pseudo-cubic\n10 16 48\nLiCe Ti O structure corresponding to the topologically transformed phase upon\n5 8 24\npartial lithiation. Based on these refined models, we recalculated the spin-resolved\nPDOS of Ti d orbitals at representative lithiation states. In the pristine state, the PDOS\nexhibits symmetric spin-up and spin-down distributions, consistent with the Ti4+ (3d0)\nelectronic configuration. Upon partial lithiation, the pseudo-cubic phase shows a slight\nspin asymmetry of Ti d states, indicating partial electron occupation of Ti 3d orbitals.\nAfter full lithiation, the spin asymmetry becomes more pronounced, suggesting further\nelectron population into Ti 3d orbitals. These results confirm a multistep Ti4+ reduction\nprocess during lithiation.\nThe updated PDOS results are now in Supplementary Fig. 34, and the relevant\ndiscussion has been updated in the revised manuscript as follows:\nLines 458-461, Page 17: The Projected DOS of Ti d orbitals (Supplementary Fig. 34)\nreveal electron transfer from lithium to Ti 3d orbitals, resulting in a progressive\ntransition from spin-symmetric to spin-asymmetric Ti 3d occupancy, indicative of the\nstepwise reduction of Ti4+.\nSupplementary Fig. 34 Projected DOS of Ti d orbitals, illustrating electronic structure\nevolution upon lithiation. a Pristine tetragonal phase. b Partial lithiation-induced\npseudo-cubic phase. c Fully lithiated CTO.\nComment 6: The DOS results of pristine model are inconsistent in Fig. 6b and Fig. 22.\nFor example, the former shows a wide bandgap about 2 eV, whereas the latter shows\nthe nature of a half metal.\nResponse: We thank the reviewer for highlighting this issue and apologize for the\nconfusion caused by the previous calculation results.\nAs noted in our response to Comment 5, we have revised the DFT calculations by\nreconstructing more accurate structural models: a tetragonal Ce Ti O for the pristine\n10 16 48\nstate and a pseudo-cubic LiCe Ti O for the topologically transformed phase. The\n5 8 24\nrecalculated projected DOS reveals that the pristine state exhibits a wide bandgap of\n~1.8 eV. Upon lithiation, the fermi level gradually shifts toward the conduction band,\nand electronic states emerge near the fermi level, indicating a transition toward half-\nmetallic behavior. These trends are consistent with the evolution of the spin-resolved\nTi d PDOS (now presented in Supplementary Fig. 34).\nAccordingly, we have revised Fig. 6b and updated the corresponding discussion\nin the revised manuscript as follows:\nLines 454-461, Page 17: Concurrently, the electronic structure undergoes significant\nmodifications during lithiation, as evidenced by the projected density of states (DOS).\nThe pristine tetragonal phase exhibits a wide bandgap of ~1.8 eV, with the conduction\nband far from the fermi level (Fig. 6b). Upon Li+ insertion, the bandgap gradually\nnarrows and the electronic states near the fermi level increase. The projected DOS of\nTi d orbitals (Supplementary Fig. 34) reveal electron transfer from lithium to Ti 3d\norbitals, resulting in a progressive transition from spin-symmetric to spin-asymmetric\nTi 3d occupancy, indicative of the stepwise reduction of Ti4+.\nFig. 6 b Projected DOS comparison for three representative states: pristine tetragonal\nphase, intermediate pseudo-cubic phase, and fully lithiated phase.\nComment 7: The constructed pristine CTO model is Ce Ti O in Fig. 6a, not the\n3 8 24\nCe Ti O described in DFT calculations, Methods. If the constructed supercell model\n4 8 24\nis Ce Ti O , the A-vacancy concentration is up to 5/8, much higher than the\n3 8 24\nexperimental 1/3 in Ce TiO . In addition, the constructed pristine model is a quasi-\n2/3 3\ncubic supercell, different from the synthesized tetragonal cell (c≈2a). How do authors\nensure the reliability of DFT calculations?\nResponse: We thank the reviewer for this insightful comment and sincerely apologize\nfor the inconsistency between the previously constructed DFT model and the\nexperimental structure.\nTo address this concern, we have reconstructed a more accurate structural model\nfor pristine Ce TiO , adopting a vacancy-ordered tetragonal superlattice Ce Ti O .\n2/3 3 10 16 48\nThis model features an A-site vacancy concentration of 3/8, which closely approximates\nthe experimental value of 1/3 while maintains a balance between computational\ncomplexity and accuracy. This model also retains the experimentally observed\ntetragonal symmetry with c ≈ 2a, ensuring consistency with the actual crystallographic\nphase. To reflect the lithiation-induced structural evolution, we further constructed a\npseudo-cubic LiCe Ti O model representing the topologically transformed phase\n5 8 24\nafter partial lithiation. These two models (Ce Ti O and LiCe Ti O ) were\n10 16 48 5 8 24\nsystematically employed to calculate lithium insertion energetics, electronic structure,\nand Li+ diffusion barriers. These modifications significantly improve the reliability and\naccuracy of our DFT calculations, ensuring closer alignment with the experimental\nresults.\nWe have thoroughly updated Fig. 6 and the corresponding discussions (Lines 428–\n485, Pages 16–18) in the revised manuscript, and have also updated the relevant figures\n(Supplementary Figs. 31-36) and data tables (Supplementary Tables 3 and 4) in\nSupplementary Information.\nFig. 6 DFT calculations for Li+ intercalation dynamics and phase evolution. a\nGeometric structures illustrating Li+ insertion sites and associated structural evolution\nduring lithiation. b Projected DOS comparison for three representative states: pristine\ntetragonal phase, intermediate pseudo-cubic phase, and fully lithiated phase. c Charge\ndensity distribution maps projected along the (100) plane for the tetragonal and pseudo-\ncubic phases. d BVS energy maps and predicted Li+ diffusion pathways in tetragonal\nand pseudo-cubic structures. e Comparison of Li+ diffusion energy barriers in tetragonal\nand pseudo-cubic structures.\nComment 8: In Fig S8, the full cell has a coulombic efficiency only ~30% for the first\ncycle, much lower than the 64.4% of half cell, why?\nResponse: We sincerely thank the reviewer for pointing out this important issue and\napologize for the oversight regarding the low first-cycle coulombic efficiency of the\nfull cell.\nAfter careful evaluation, we think that the low first-cycle coulombic efficiency did\nnot originate from the intrinsic properties of the electrodes, but was primarily caused\nby suboptimal capacity balancing in the previous full cell configuration. Specifically,\nthe previous full cell suffered from an inappropriate capacity ratio between the cathode\nand anode, leading to excessive irreversible lithium consumption during the first cycle.\nThis capacity mismatch hindered efficient utilization of the active lithium source,\nresulting in a low first-cycle coulombic efficiency.\nTo address this issue, we reassembled the full cell with optimized electrode mass\nloadings to ensure a more appropriate N/P ratio of 1.1-1.2. As a result, the first-cycle\ncoulombic efficiency was significantly improved to 62.5%, which was close to that of\nthe half cell (64.4%). Moreover, the optimized full cell exhibited enhanced rate\ncapability and cycling performance. It delivered a capacity retention of 61.3% at 10 C\nand maintained 80% capacity retention after 1000 cycles at 5 C, demonstrating the\nexcellent electrochemical performance of the full cell.\nThe updated electrochemical results of the full cell have been incorporated into\nSupplementary Fig. 15, and the corresponding discussion has been added to the\nrevised manuscript as follows:\nLines 255-262, Page 10: At 0.1 C, the full cell delivers a high discharge capacity of\n150 mAh g-1, with a first-cycle coulombic efficiency of 62.5% (Supplementary Fig.\n15a), comparable to that of the half cell. As the rate increases to 0.5, 1, 2, 5, and 10 C,\nthe full cell maintains specific capacities of 125, 115, 108, 100, and 92 mAh g-1,\nrespectively (Supplementary Fig. 15b). Notably, the capacity retention at 10 C reaches\n61.3%, demonstrating excellent rate capability. Furthermore, long-term cycling at 5 C\nconfirms outstanding stability, with 80% capacity retention after 1000 cycles\n(Supplementary Fig. 15c).\nSupplementary Fig. 15 Electrochemical performance of LiFePO ||CTO full cells. a\n62.5%. b Charge/discharge profiles at different current rates from 0.1 C to 10 C. c Rate\ncapability and long-term cycling performance at 5 C.\nComment 9: Line 313, “the valance reduction” may be “the valence reduction”.\nResponse: We thank the reviewer for pointing out this typographical error. The term\n“valance” has been corrected to “valence” in the revised manuscript (Line 379, Page\n14).\nResponses to the comments from Reviewer 4\nReviewer 4\nGeneral comments: In this manuscript, vacancy-ordered Ce TiO (CTO) is reported\n2/3 3\nas a perovskite-type anode with high gravimetric capacity (> 200 mAh g⁻¹ at ≈ 0.8 V),\nultrafast-rate capability (≈ 80 mA h g-1 at 50 C), and sustained cycling stability (90%\ncapacity retention after 10,000 cycles). Accelerating Li+ transport by A-site vacancy\nordering would be a new route which, if confirmed under electrode loadings, can be\nmuch better than graphite and spinel Li Ti O . The proof of concept is achieved\n4 5 12\nutilizing significant measurements, however need to refine some sections as:\nelectrochemical access is restricted to thin (≈ 1.5 mg cm-2) electrodes without any\nvolumetric benchmarking; the 10,000 cycle claim rests simply upon a single end-point\nvalue; and the \"pseudo-cubic\" lithiated phase, as well as potential Ce-valence changes.\nConsequently, I recommend minor revision with an emphasis of the need for extensive\nelectrode level validation, detailed cycle dataset, quantitative structural analysis and\ndirect comparison with the industry benchmarks.\nResponse: We sincerely thank the reviewer for the constructive and insightful\ncomments. We appreciate the reviewer’s recognition of the novelty and potential of the\nvacancy-ordered CTO as a high-performance anode. In response to the reviewer’s\ncomments, we have accordingly revised the manuscript and provided additional\nexperimental results as outlined below:\n(1) Electrode-level validation and industry benchmarking:\nTo address this comment, we conducted additional electrochemical measurements\nusing thick CTO electrodes with a high mass loading of ~10.2 mg cm-2. The results\n(now presented in Supplementary Fig. 10) demonstrate that the thick CTO electrode\nretains excellent rate capability (e.g., 67.7% capacity retention at 20 C) and cycling\nstability (80% capacity retention after 1000 cycles at 20 C). We also evaluated the\nvolumetric capacity of both thin (~1.5 mg cm-2) and thick electrodes, which achieve\nmaximum volumetric capacities of 550 mAh cm-3 and 394 mAh cm-3, respectively. For\ncomparison, we benchmarked CTO against representative micro-sized and nano-\nstructured intercalation-type anodes (Supplementary Fig. 11), confirming its superior\nperformance over conventional anode materials. These results highlight the practical\napplicability of CTO. The corresponding discussion has been added to the revised\nmanuscript as follows:\nLines 222-236, Page 9: Most fast-charging anode materials reported to data rely on\nthin electrodes with low mass loadings (1-2 mg cm-2), which limit their practical\napplicability. To further demonstrate the superior electrochemical performance of CTO\nunder application-relevant conditions, we evaluated thick electrodes with a high mass\nloading of 10.2 mg cm-2. As shown in Supplementary Fig. 10, the CTO electrode\ndelivers specific capacities of 193, 173, 160, 146, 117, 89, 63, and 30 mAh g-1 at current\nrates from 0.1 C to 50 C, demonstrating excellent rate capability even at a high loading.\nMoreover, it retains 80% of the initial capacity after 1000 cycles at 20 C, maintaining\nconsiderable cycling stability. Beyond gravimetric performance, we also assessed the\nvolumetric capacity of CTO. At a low mass loading of 1.5 mg cm-2, CTO exhibits a\nmaximum volumetric capacity of 550 mAh cm-3, and still achieves 394 mAh cm-3 at\nthe high loading of 10.2 mg cm-2 (Supplementary Fig. 11). This performance\nsignificantly outperforms that of previously reported micro-sized or nano-structured\nintercalation-type anodes (Supplementary Fig. 11c), further highlighting the practical\nadvantages of CTO for high-energy-density and fast-charging LIBs.\nSupplementary Fig. 10 Electrochemical performance of CTO electrode with a high\nmass loading of 10.2 mg cm-2. a Charge-discharge profiles at various current rates from\n0.1 C to 50 C. b Rate capability and corresponding coulombic efficiency at 0.1, 0.5, 1,\n2, 5, 10, 20, and 50 C. c Cycling performance at a high rate of 20 C.\nSupplementary Fig. 11 Volumetric performance of CTO electrode with low (1.5 mg\ncm-2) and high (10.2 mg cm-2) mass loadings. a Cross-sectional SEM images. b Areal\ncurrent density versus volumetric capacity. c Comparison of volumetric capacities at 1\nC and 10 C with previously reported intercalation-type anodes11, 20, 21, 22, 23, 24, 25, 26, 27, 28,\n29, 30, 31.\n(2) Clarification of the 10000-cycle dataset:\nIn response, we have provided representative charge-discharge profiles recorded\nthroughout the 10000-cycle test at 20 C (now presented in Supplementary Fig. 7),\nclearly demonstrating exceptional cycling stability with minimal capacity decay\nbetween the initial and final cycles.\nSupplementary Fig. 7 Detailed discharge/charge curves during prolonged high-rate\ncycling at 20 C.\n(3) Structural validation of the pseudo-cubic phase:\nTo further substantiate the lithiation-induced transition from tetragonal to pseudo-\ncubic phase, we have supplemented additional experimental evidence and analysis.\nSpecifically, pair distribution function (PDF) refinements of pristine and lithiated\nsamples further confirm a transition from a long-range ordered tetragonal structure to a\nshort-range ordered pseudo-cubic phase upon lithiation (Supplementary Fig. 29).\nMoreover, ex situ Raman spectroscopy collected near the transition volage of 1.25 V\nreveals progressive shifts in the dominant vibrational modes and merging of low-\nfrequency modes as lithiation proceeds, consistent with an increase in structural\nsymmetry during phase transition (Supplementary Fig. 19). The details are provided\nin our responses to Comments 3 and 4 below.\n(4) Ce valence state evolution:\nThe Ce M -edge EELS spectra (Supplementary Fig. 26) confirm that the\n4,5\nvalence state of Ce remains unchanged during lithiation. To further validate this, we\nconducted ex situ Ce 3d XPS measurements at different states (Supplementary Fig. 6),\nwhich consistently reveal the electrochemical inactivity of Ce throughout the lithiation\nprocess. These complementary results further indicate that the redox activity during\nlithiation originates from Ti rather than Ce, as further elaborated in our response to\nComment 1 below.\nComment 1: First, the reversible capacity of 220–221 mA h g-1 at 0.1 C surpasses the\nTi4+ → Ti3+ theoretical limit for CTO (≈ 142 mA h g-1), while EELS shows no Ce\nvalence shift, implying that the extra ≈ 80 mA h g-1 likely arises from over-estimated\nactive mass or parasitic side reactions rather than intrinsic redox.\nResponse: We thank the reviewer for this insightful comment. The concern regarding\nthe reversible capacity exceeding the theoretical Ti4+ → Ti3+ redox limit (~ 142 mAh g-\n1) is valid under the assumption of a single-electron redox. However, multiple lines of\nevidence support the occurrence of a two-electron Ti4+ → Ti2+ redox process during\nlithiation, corresponding to a theoretical capacity of 284 mAh g-1, which sufficiently\naccounts for the measured capacity of 221 mAh g-1.\nSpecifically:\n1. The CV curves (Fig. 2a) exhibit three distinct redox couples at 1.45 V/1.24 V, 0.81\nV/0.66 V, and 0.30 V/0.10 V. The redox couples at 1.45 V/1.24 V and 0.81 V/0.66\nV are assigned to the stepwise Ti4+/Ti3+ redox process, consistent with previous\nreports on Ti-based anodes (e.g., Nano Energy, 2024, 119, 109065; Chem. Eng. J.,\n2024, 479, 147765; Electrochem. Commun., 2013, 32, 5-8; Adv. Energy Mater.,\n2023, 13, 2302015). The redox couple at lower potential 0.30 V/0.10 V is attributed\nto a subsequent Ti3+/Ti2+ redox transition under deep lithiation, which is commonly\nobserved below 0.5 V in other Ti-based anodes (e.g., Energy Environ. Sci., 2017,\n10, 1456-1464; Adv. Sci., 2025, 2410543; Nano Energy, 2022, 94, 106972; ACS\nAppl. Mater. Interfaces, 2024, 16, 898–906).\n2. Ti L -edge EELS spectra (Supplementary Fig. 25) collected at the pristine state,\n2,3\n1.25 V, and 0.01 V show progressive shifts of the L and L peaks toward lower\n3 2\nenergies, indicating the stepwise reduction of Ti from +4 to lower oxidation states.\nIn contrast, Ce M -edge EELS spectra remain unchanged throughout lithiation,\n4,5\nconfirming that Ce does not participate in redox reactions.\n3. To further corroborate our interpretation, we have supplemented ex situ XPS\nmeasurements on CTO (Supplementary Fig. 6). The Ti 2p XPS peaks exhibit a\ncontinuous shift toward lower binding energies during lithiation, consistent with the\nTi4+ → Ti3+→ Ti2+ redox sequence. Meanwhile, the Ce 3d spectra remain\nunchanged, further confirming the electrochemical inactivity of Ce during lithiation.\nOverall, these results consistently demonstrate the occurrence of a Ti4+ → Ti2+\ntwo-electron redox process in CTO and rule out the Ce redox, thereby accounting for\nthe observed high reversible capacity.\nSupplementary Fig. 6 XPS spectra of CTO at the pristine state, after discharging to\n1.0 V, and after discharging to 0.01 V. a Evolution of Ce 3d high-resolution spectra. b\nEvolution of Ti 2p high-resolution spectra.\nComment 2: Second, the initial Coulombic efficiency is only 64 %, yet the solid-\nelectrolyte interphase is uncharacterised, leaving the source of irreversible capacity\nloss- and its consequences for long-term stability—unclear.\nResponse: We thank the reviewer for highlighting this important point regarding solid\nelectrolyte interphase (SEI) characterization. In response, we conducted additional SEI\ncharacterizations on CTO electrodes at different stages of the initial discharge (pristine,\ndischarging to 0.5 V, and 0.01 V) using transmission electron microscopy (TEM) and\nX-ray photoelectron spectroscopy (XPS).\nAs shown in Supplementary Fig. 4, TEM images reveal a progressive thickening\nof the surface passivation layer during initial lithiation, indicating the formation and\ngrowth of SEI. Corresponding Li 1s XPS spectra show the emergence and evolution of\ncharacteristic SEI components, including LiF, Li CO , and ROCO Li. As lithiation\n2 3 2\nproceeds, the relative intensities of LiF and Li CO signals increase while those of\n2 3\norganic species decrease, suggesting an enrichment of inorganic SEI components.\nThe irreversible lithium consumption associated with SEI formation accounts for the initial\ncoulombic efficiency of ~64%. Meanwhile, the formation of a robust, LiF/Li CO -rich\n2 3\ninorganic SEI layer is favorable for improving mechanical and interfacial stability of\nthe electrodes, thereby contributing to the long-term cycling stability.\nThese results have been added to Supplementary Fig. 4, and the corresponding\ndiscussion has been included in the revised manuscript as follows:\nLines 171-181, Page 7: The initial capacity loss is primarily attributed to the formation\nof the solid electrolyte interphase (SEI). As shown in Supplementary Fig. 4, TEM\nimages reveal progressive thickening of the SEI layer on CTO particle surfaces during\nthe first discharge. High-resolution Li 1s X-ray photoelectron spectroscopy (XPS)\nspectra indicate that the SEI consists of both organic (ROCO Li) and inorganic (Li CO\n2 2 3\nand LiF) components, which are from different interfacial reactions. ROCO Li and\n2 3\nside reactions with conductive carbon, while LiF arises from the decomposition of\nLiPF salt. As lithiation proceeds, the relative intensities of Li CO and LiF signals\n6 2 3\nincrease while those of organic species decrease, suggesting a gradual enrichment of\ninorganic SEI components at low potentials, which leads to irreversible lithium\nconsumption.\nSupplementary Fig. 4 Evolution of SEI layer on CTO during initial lithiation. a-c TEM\nimages of CTO particle surfaces at the pristine state (a), after discharging to 0.5 V (b),\nand 0.01 V (c), showing the formation and growth of the SEI layer. d-f Corresponding\nhigh-resolution Li 1s XPS spectra, revealing the chemical evolution of Li-containing\nspecies during lithiation.\nComment 3: Third, the purported reversible transition from tetragonal to pseudo-cubic\nsymmetry is supported only by in-situ XRD peak shifts and qualitative STEM images;\nrigorous Rietveld refinement of the lithiated phase and Ce-valence analysis are required\nto validate this structural evolution.\nResponse: We sincerely thank the reviewer for this constructive suggestion. To further\nsubstantiate the structural transition and clarify the Ce valence evolution, we have\nconducted additional structural and spectroscopic analysis, as detailed below:\n(1) Structural validation via pair distribution function (PDF) analysis:\nTo provide more rigorous structural information beyond XRD and STEM, we\nconducted PDF refinements on both pristine and lithiated CTO samples. The results\nconfirm that the pristine and lithiated states correspond to a long-range ordered\ntetragonal phase and a short-range ordered pseudo-cubic phase, respectively.\nAdditionally, in the short-range region (3.1 - 6.1 Å), the lithiated sample exhibits a new\nCe-O correlation peak at ~5.0 Å and enhanced intensities of Ce1-Ti and Ce1-Ce2\ncorrelations. These changes indicate a rearrangement of the local Ce coordination\nenvironment upon lithiation. Moreover, the Ce1-Ce1, Ce-O, and Ce1-Ce2 peaks in the\nlithiated state exhibit narrower widths compared to those in the pristine state, suggesting\nimproved short-range ordering of Ce atoms. The PDF results provide additional\nevidence for a lithiation-induced transition in both structural symmetry and local\nordering.\nThe PDF results have been added to Supplementary Fig. 29, and the\ncorresponding discussion has been included in the revised manuscript as follows:\nLines 402-414, Page 15: Pair distribution function (PDF) analysis was performed to\ncompare the pristine and lithiated states. As shown in Supplementary Fig. 29a, the\nPDF refinements reveal a transition from long-range tetragonal ordering to short-range\npseudo-cubic ordering, consistent with the structural models in Supplementary Fig.\n29b. In the short-range region (3.1 - 6.1 Å), the intensities of Ce1-Ti and Ce1-Ce2\ncorrelations increase, and a new Ce-O correlation emerges at ~5.0 Å upon lithiation\n(Supplementary Fig. 29c), indicating a rearrangement of the Ce coordination\nenvironment. Moreover, the Ce1-Ce1, Ce-O, and Ce1-Ce2 peaks become noticeably\nnarrower after lithiation (Supplementary Fig. 29d), reflecting enhanced short-range\nordering of Ce atoms. These results collectively demonstrate that lithiation triggers a\ntopological transition characterized by enhanced short-range structural ordering around\nCe coordination environment.\nSupplementary Fig. 29 PDF analysis revealing lithiation-induced transformation from\nlong-range tetragonal ordering to short-range pseudo-cubic ordering. a PDF\nrefinements of pristine and lithiated CTO. b Structural models of the long-range ordered\ntetragonal and short-range ordered pseudo-cubic phases. c Comparison of PDF peaks\nin the short-range region (3.1 - 6.1 Å), where the three dominant peaks are assigned to\nCe1-Ti1 (~3.4 Å), Ce1-Ce1 (~3.8 Å), and Ce1-Ce2 (~5.5 Å) correlations, additional\nminor peaks are attributed to Ti-O and Ce-O distances. d Comparison of Ce1-Ce1, Ce-\nO, and Ce1-Ce2 peaks.\n(2) Ce valence evolution confirmed by ex situ XPS:\nAs detailed in our response to Comment 1, we have conducted ex situ Ce 3d XPS\nmeasurements on CTO at the pristine state, a partially lithiated state (1.0 V), and a fully\nlithiated state (0.01 V). The Ce 3d XPS spectra remain essentially unchanged across\nthese states, indicating that Ce maintains a stable +3 oxidation state throughout the\nlithiation process. This result is consistent with the Ce M -edge EELS spectra\n4,5\npresented in Supplementary Fig. 26 and confirms the electrochemical inactivity of Ce.\nThese findings reveal that the structural transition is induced by lithium insertion and\nassociated topological rearrangement, rather than changes in Ce valence. The ex situ Ce\n3d XPS spectra have been included in Supplementary Fig. 6.\nComment 4: Finally, complementary diagnostics--such as operando or ex-situ Raman\nspectroscopy near the 1.25 V transition and post-cycling XPS or FTIR of the electrodes-\nare needed to corroborate the symmetry change and identify SEI constituents\nresponsible for the low first-cycle efficiency.\nResponse: We thank the reviewer for this valuable suggestion. In response, we have\nperformed additional characterizations to corroborate the lithiation-induced symmetry\nchange near the 1.25 V and to identify the SEI constituents after cycling.\n(1) Ex situ Raman spectroscopy near the 1.25 V:\nTo capture the symmetry changes across the phase transition region, we have\nconducted ex situ Raman spectroscopy on CTO electrodes at a series of voltages (1.35\nV, 1.30 V, 1.25 V, 1.20 V, and 1.10 V). As the voltage decreases, the Raman spectra\nexhibit evident changes, including a progressive downshift of the main peaks,\ndisappearance of shoulder modes at ~124 and ~200 cm-1, and increased intensities of\nvibrational modes at ~242 and ~664 cm-1. These spectral changes, especially the\ngradual merging of low-frequency modes in the 100–200 cm-1 region, indicate an\nincrease in structural symmetry, consistent with the transition from a tetragonal phase\nto a higher-symmetry pseudo-cubic phase.\nThe Ramn spectra have been added to Supplementary Fig. 19, and the relevant\ndiscussion has been included in the revised manuscript as follows:\nLines 320-325, Page 12: Raman spectra collected across the phase transition region\n(Supplementary Fig. 19) reveal systematic vibrational changes during lithiation. The\ngradual downshift of main peaks, along with the progressive merging of low-frequency\nmodes, indicates an increased structural symmetry, consistent with the transition toward\na pseudo-cubic phase with higher symmetry.\nSupplementary Fig. 19 Ex situ Raman spectra of CTO across the phase transition\nregion. Raman spectra were collected at a series of discharge voltages (1.35 V, 1.30 V,\n1.25 V, 1.20 V, and 1.10 V) to monitor structural evolution during lithiation. As the\ndischarge proceeds, the main peaks at ~172 and ~354 cm-1 gradually shift to lower\nwavenumbers, while the shoulder peaks at ~124 and ~200 cm-1 progressively weaken\nand merge. Meanwhile, the intensities of modes at ~242 and 664 cm-1 increase below\n1.30 V. These spectral changes, particularly the progressive merging of low-frequency\nmodes in the 100-200 cm-1 region, indicate a gradual increase in structural symmetry,\nconsistent with the transition from a tetragonal to a higher-symmetry pseudo-cubic\nphase.\n(2) Post-cycling XPS analysis of SEI constituents:\nAs detailed in our response to Comment 2, we have already provided TEM images\nand Li 1s XPS spectra of CTO electrodes at different states of the initial discharge,\nwhich elucidates the origin of the low first-cycle coulombic efficiency. To further\naddress the reviewer’s concern, we have now included additional TEM and Li 1s XPS\nanalysis on the CTO electrode after cycling. The post-cycling Li 1s XPS spectrum\nreveals that the SEI is predominately composed of inorganic components such as LiF\nand Li CO , along with a minor organic component ROCO Li, which are nearly\n2 3 2\nconsistent with the SEI formed during the initial discharge. TEM analysis shows a slight\nincrease in SEI thickness after cycling, indicating gradual but stable SEI evolution.\nThese results confirm that the SEI layer remains compositionally stable and\nmechanically robust after cycling, which enables the interfacial stability and long-term\ndurability of the electrode.\nThe post-cycling TEM and XPS results have been added to Supplementary Fig.\n5, and the relevant discussion has been included in the revised manuscript as follows:\nLines 182-186, Page 7: After cycling, the SEI maintains a similar chemical\ncomposition but becomes slightly thicker, with further accumulation of inorganic\ncomponents such as LiF and Li CO (Supplementary Fig. 5). This robust inorganic-\n2 3\nrich SEI enhances the mechanical and chemical stability of the electrode interface,\nwhich is critical for mitigating continuous electrolyte decomposition and enabling long-\nterm cycling durability.\nSupplementary Fig. 5 a TEM image of the CTO electrode after cycling. b\nCorresponding high-resolution Li 1s XPS spectrum."}, {"round": 2, "reviewer_comments": {"reviewer_1": "The authors have addressed the questions in the revised manuscript. It is thus recommended to be accepted.", "reviewer_2": "I co-reviewed this manuscript with one of the reviewers who provided the listed reports. This is part of the Nature Communications initiative to facilitate training in peer review and to provide appropriate recognition for Early Career Researchers who co-review manuscripts.", "reviewer_3": "The authors have provided thorough and detailed responses and revisions to my comments, which appear to have addressed all of my concerns. One minor suggestion is that the solid circles representing capacity and Coulombic efficiency in all cycling plots should be clearly labeled. Otherwise, it would be difficult for beginners in battery research to distinguish which corresponds to capacity and which to Coulombic efficiency.Overall, I recommend the publication of this paper on Nature Communications.", "reviewer_4": "I have reviewed the manuscript and I find that the authors have complied with all comments and criticisms outlined in my earlier review of the manuscript. The revised manuscript is therefore acceptable as it is."}, "author_response": "Response Letter to Reviewers\nResponses to the comments from Reviewer 1\nComments: The authors have addressed the questions in the revised manuscript. It is\nthus recommended to be accepted.\nResponse: We sincerely thank the reviewer for the positive evaluation and\nrecommendation. We are delighted that our revisions have satisfactorily addressed the\nconcerns raised.\nResponses to the comments from Reviewer 2\nComments: I co-reviewed this manuscript with one of the reviewers who provided the\nlisted reports. This is part of the Nature Communications initiative to facilitate training\nin peer review and to provide appropriate recognition for Early Career Researchers who\nco-review manuscripts.\nResponse: We sincerely thank the reviewer for the valuable comment and constructive\nfeedback.\nResponses to the comments from Reviewer 3\nComments: The authors have provided thorough and detailed responses and revisions\nto my comments, which appear to have addressed all of my concerns. One minor\nsuggestion is that the solid circles representing capacity and Coulombic efficiency in\nall cycling plots should be clearly labeled. Otherwise, it would be difficult for beginners\nin battery research to distinguish which corresponds to capacity and which to\nCoulombic efficiency. Overall, I recommend the publication of this paper on Nature\nCommunications.\nResponse: We sincerely thank the reviewer for the positive evaluation and\nrecommendation for publication. We also appreciate the constructive suggestion\nregarding the labeling of solid circles in the cycling plots. Following the advice, we\nhave carefully revised all relevant figures to clearly distinguish capacity and Coulombic\nefficiency, thereby improving clarity for readers.\nResponses to the comments from Reviewer 4\nComments: I have reviewed the manuscript and I find that the authors have complied\nwith all comments and criticisms outlined in my earlier review of the manuscript. The\nrevised manuscript is therefore acceptable as it is.\nResponse: We sincerely thank the reviewer for the positive comment. We are grateful\nfor the constructive feedback provided throughout the review process."}]

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