Identifying cycling current-dependent degradation pathways in high-voltage layered sodium–ion cathodes

Identifying cycling current-dependent degradation pathways in high-voltage layered sodium–ion cathodes | Research Square window.SnipcartSettings = { analytics: { enabled: false } }; (function() { var accessVector = localStorage.getItem('access_vector') || ''; window.dataLayer = window.dataLayer || []; if (accessVector) { window.dataLayer.push({ user: { profile: { profileInfo: { snid: accessVector } } } }); } })(); (function(w,d,s,l,i){w[l]=w[l]||[];w[l].push({'gtm.start':new Date().getTime(),event:'gtm.js'});var f=d.getElementsByTagName(s)[0],j=d.createElement(s),dl=l!='dataLayer'?'&l='+l:'';j.async=true;j.src='https://www.googletagmanager.com/gtm.js?id='+i+dl;f.parentNode.insertBefore(j,f);})(window,document,'script','dataLayer','GTM-K279D39R'); Browse Preprints In Review Journals COVID-19 Preprints AJE Video Bytes Research Tools Research Promotion AJE Professional Editing AJE Rubriq About Preprint Platform In Review Editorial Policies Our Team Advisory Board Help Center Sign In Submit a Preprint Cite Share Download PDF Article Identifying cycling current-dependent degradation pathways in high-voltage layered sodium–ion cathodes Haidong Liu, Yongchun Li, Qiao Zhang, Katherine Mazzio, Yanan Sun, and 8 more This is a preprint; it has not been peer reviewed by a journal. https://doi.org/ 10.21203/rs.3.rs-8109344/v1 This work is licensed under a CC BY 4.0 License Status: Under Review Version 1 posted You are reading this latest preprint version Abstract High-voltage operation boosts the energy density of sodium layered oxides by activating oxygen redox, but the underlying degradation mechanisms remain controversial, hindering rational materials design. Here we demonstrate that the applied cycling current determines which degradation pathway dominates, thereby resolving this long-standing inconsistency. Using synchrotron-based X-ray probes, advanced microscopic characterizations and first-principle calculation on a compositionally controlled set of Sc-, Sc–Ti-, and Sc–Ti–B-doped O3-type oxides, it is revealed that low current preserves oxygen-redox participation and drives oxygen-loss-induced degradation, whereas high current suppresses oxygen redox and induces a transition to a structure-governed evolution pathway along simplified P3-type routes. This compositional set makes the current-dependent degradation behaviour experimentally distinguishable, allowing the oxygen-redox-driven and structure-governed modes to be clearly disentangled. These findings establish a unified current-governed mechanistic framework and provide guidance for designing durable high-voltage sodium-ion cathodes. Physical sciences/Chemistry/Electrochemistry/Batteries Physical sciences/Energy science and technology/Energy storage/Batteries Physical sciences/Materials science/Materials for energy and catalysis/Batteries sodium ion batteries layered oxide cathode cycling current-dependent degradation pathway oxygen redox phase transition synchrotron characterization Figures Figure 1 Figure 2 Figure 3 Figure 4 Figure 5 Figure 6 Figure 7 Introduction Sodium-ion batteries (NIBs) are a promising alternative to lithium-ion batteries (LIBs) for grid-scale energy storage due to the natural abundance and low cost of sodium 1 – 3 . Advancing high-power, durable cathode materials is essential, as the cathode largely determines battery cost and performance 2 . Among various cathodes, O3–type Ni–Mn based layered oxides, characterized by Na⁺ ions in octahedral sites and ABC stacking, can offer high capacity (> 200 mAh g − 1 above 4.1 V vs. Na + /Na) along with high average redox potentials, making them attractive for practical use 4 , 5 . The high capacity is supported by the activation of oxygen redox processes 6 . Nevertheless, despite this advantage, their performance rapidly deteriorates under fast Na de/intercalation reactions (> 3C), manifested by poor rate capability and cycling stability, as widely reported for O3–type Ni–Mn based cathodes 7 – 9 . Such degradation originates from structural collapse, oxygen loss, cation migration, and inhomogeneous strain accumulation, with strongly rate-dependent pathways that ultimately constrain long-term stability and scalability 7 , 10 – 15 . Doping/substitution strategies, as well as the more recent concept of entropy guided optimization, represent rational approaches for designing cathodes with high performance, and have been widely employed to mitigate the above-mentioned issues 9 , 11 , 13 , 16 – 21 . Even trace amounts of dopants can significantly alter the crystal and electronic structures of layered cathodes, thereby influencing phase transition pathways, and ultimately improving electrochemical properties. Suppressing unfavorable phase transitions, such as the P3–O3 transformation in NaNi 0.5 Mn 0.5 O 2 is widely recognized as a key strategy for enhancing the electrochemical performance of O3–type layered oxides 7 . However, while essential, phase–transition suppression alone does not fully determine the long-term stability under high-voltage operation. For example, single-element doping with Li, Mg, Fe, Cu, or Zn has been reported to suppress these unfavorable phase transitions in O3–type NaNi 0.5 Mn 0.5 O 2 by promoting a more reversible structural evolution during low-current cycling 7 , 22 – 25 . Nevertheless, despite the structural improvements achieved through single-element doping, such materials still experience rapid capacity fading, particularly when cycling under high-current densities. By contrast, cathodes incorporating multiple dopants such as NaCu 0.1 Ni 0.2 Co 0.2 Fe 0.2 Mn 0.15 Ti 0.15 O 2 and NaNi 0.35 Mn 0.35 Cu 0.01 Fe 0.1 Ti 0.05 Sn 0.05 O 2 , deliver remarkable stability at high-current densities despite undergoing similar phase transition processes 9 , 13 . However, these materials typically deliver lower specific capacities at low -current densities compared with their single-element doped counterparts. These observations indicate a trade-off in which specific capacity is sacrificed to achieve improved rate stability. Yet, this compromise inevitably limits capacity across both low- and high-current operations, underscoring the intrinsic rate-dependent degradation mechanisms of high-voltage layered oxides. Beyond structural effects, redox activities strongly influence both capacity and cycling stability, with oxygen redox playing a particularly critical role in high-voltage layered oxides. However, oxygen redox is often accompanied by oxygen loss from unstable local environments and resulting structural instabilities, which accelerates degradation and compromises long-term performance 10 . To address this challenge, dopants such as B, S, and Se have been employed to stabilize oxygen redox by enhancing covalent bonding within the lattice 26 – 29 . Nevertheless, most studies have primarily focused on the initial cycles, providing limited direct evidence of oxygen stability during extended operation. Although online/differential electrochemical mass spectrometry (OEMS or DEMS, respectively) has been widely used to probe oxygen redox reversibility and has yielded valuable insights into early-cycle behavior 21 , 26 – 28 , this technique is largely restricted to the first few cycles and may fail to capture additional oxygen loss pathways that emerge during extended cycling. Furthermore, most measurements are conducted at low current densities, which do not adequately reveal the degradation mechanisms that dominate under high-current densities, particularly those associated with structural evolution. Collectively, these findings underscore the presence of rate-dependent degradation pathways which are often underappreciated, and which are critical to advancing the development of stable oxygen–redox cathodes. To tackle both structural and redox instabilities, dopants were chosen based on their distinct yet complementary effects. Scandium (Sc 3+ ), with its relatively large ionic radius and stable 3 + valence, can act as a structural stabilizer that strengthens the transition-metal (TM)–O framework and help suppress cation migration, thereby mitigating phase transitions and slab gliding observed in undoped O3-type oxides. Titanium (Ti 4+ ) is known for its strong Ti–O covalency and electronic stability, and can further enhance the robustness of the framework and help reduce Jahn–Teller distortion by partially replacing Mn 3+ , while moderating irreversible oxygen redox activity at high voltages. In contrast, boron (B 3+ ) serves as an electronic structure modifier that strengthens TM–O–B covalency, stabilizing oxygen redox processes and suppressing oxygen loss during cycling. The combination of these dopants is therefore expected to simultaneously stabilize the cationic and anionic sublattices, providing a model system to disentangle rate-dependent degradation mechanisms in high-voltage O3-type Na layered oxides. In this work, Sc and Ti were introduced to partially substitute Ni and Mn in O3- NaNi 0.5 Mn 0.5 O 2 , while a small amount of B (2 mol%) was further incorporated by partially replacing Ti. This compositional tuning enables a systematic investigation of how these dopants modulate transition-metal (TM) and oxygen redox activities, and how they correlate with the degradation mechanisms under varying current densities. It was found that these cathodes exhibit rate-dependent cycling stability, with Sc–Ti–B co-doping (NaSc 0.1 Ni 0.4 Mn 0.4 Ti 0.08 B 0.02 O 2 ) significantly enhancing both capacity and cycling stability at low currents. The Sc–Ti–B co-doped material delivers 147 mAh g − 1 at 1C (200 mA g – 1 ), retaining 83% of its capacity after 100 cycles within a voltage window of 2.5–4.4 V. However, at a 3C rate, the Sc–Ti co-doped material (without B) (NaSc 0.1 Ni 0.4 Mn 0.4 Ti 0.1 O 2 ) demonstrates superior stability. Using advanced characterization methods, including resonant inelastic X–ray scattering (RIXS), soft/hard X–ray absorption spectroscopy (XAS), transmission electron microscopy (TEM), and density functional theory (DFT) calculations, we identify that B doping greatly improves structural stability and oxygen–redox reversibility at low-current densities, primarily by reducing oxygen loss. These findings indicate that oxygen loss during cycling is a key contributor to capacity fading at low-current densities, beyond structural degradation alone. At high-current densities, by contrast, all O3–type materials eventually undergo a single–phase transition to the P3 phase. For the B-containing composition, although improved oxygen stability enhances structural reversibility at low-current densities, it exacerbates structural changes during rapid de/sodiation, thereby compromising performance at high-current densities. These findings reveal two fundamentally distinct degradation pathways, one governed by oxygen redox stability at low current and the other governed by structural robustness at high current, and they provide a unified mechanistic framework that explains high-voltage behaviour in layered sodium-ion cathodes while establishing the basis for developing materials that remain stable under varying cycling conditions. Results and discussion Compositions of substituted O3 NaNi 0.5 Mn 0.5 O 2 Motivated by our earlier results, O3–NaNi 0.5 Mn 0.5 O 2 (designated as O3–undoped) shows only Ni and oxygen redox activity within the voltage range of 2.5–4.4 V, delivering an initial discharge capacity exceeding 200 mAh g − 1 . However, this material suffers from rapid capacity fading, even with single-element doping such as Mg. 7 To further investigate strategies for stabilizing this high-voltage cathode, various dopants, including Sc, Ti, and B, were incorporated into the parent compound. As shown in Fig. 1 a, the designed compositions are O3–NaSc 0.10 Ni 0.40 Mn 0.50 O 2 (O3–Sc doping), O3–NaSc 0.10 Ni 0.40 Mn 0.40 Ti 0.10 O 2 (O3–ScTi doping), and O3–NaSc 0.10 Ni 0.40 Mn 0.40 Ti 0.08 B 0.02 O 2 (O3–ScTiB doping). All samples exhibit a well-defined layered structure, with only a minor presence of NiO impurity (see Fig. 1 a, 1 b, S1, S2 and Table S1 –S4) 22 , 23 , 30 . Notably, doping significantly reduces the amount of this impurity, particularly in the Sc–Ti–B doped material. Upon Sc doping, the primary (003) reflection is clearly shifted to a lower 2θ angle, indicating an expansion along the c axis. However, with the introduction of Ti, the (003) reflection shifts back toward a higher 2θ angle, and this trend continues with additional B doping. X–ray pair distribution function (PDF) analysis further confirms that the dopants are successfully incorporated into the crystal structure, as evidenced by distinct shifts in the nearest-neighbor transition metal–oxygen (TM–O, TM = Ni, Mn, and Sc) and transition metal–transition metal (TM–TM) distances (see Figure S3). B doping stands out by producing the shortest TM–O bond length and the longest TM–TM distance, which may facilitate the formation of covalent bonds to stabilize O 26 . Moreover, the density of O 2p states near the Fermi level increases significantly, indicating that doping introduces more electronic states, which may affect electronic conductivity and oxygen redox activities (see Fig. 1 d and S4). These observations confirm that each dopant induces distinct modifications in the local structural environment of the layered oxide framework. All samples exhibit a similar morphology, characterized by irregularly shaped particles with sizes up to approximately 3 µm (see Figure S5). High-angle annular dark-field scanning transmission electron microscopy (HAADF-STEM) further confirms that the dopants exert minimal influence on the crystal structure, as evidenced by the similar atomic arrangements and the absence of cation mixing in the bulk regions of both materials (see Fig. 1 e). The accompanying energy-dispersive X-ray spectroscopy (EDS) analysis also verifies the presence of all constituent elements and reveals their uniform spatial distribution. Electrochemical properties The electrochemical properties of all cathode materials were evaluated in half cells within a voltage window of 2.5–4.4 V versus Na + /Na, as shown in Fig. 2 . Each material exhibits distinct charge-discharge capacities and unique voltage curves during the initial cycle (see Fig. 2 a). The specific charge-discharge capacities for O3–undoped, O3–Sc doping, O3–ScTi doping, and O3–ScTiB doping are 233.3/197.0, 173.9/158.3, 191.1/159.1, and 200.2/169.8 mAh g − 1 , respectively. The undoped material (O3–undoped) displays multiple small voltage plateaus between 3.0 and 4.0 V, along with long and flat plateaus below 3.0 V and above 4.0 V. Doping effectively smooths the charge-discharge curves within the 3.0–4.0 V range, though each dopant has a distinct effect on the low-voltage ( 4.0 V) regions. Sc doping preserves both the low- and high-voltage plateaus, although their lengths are shortened. This behavior is consistent with the dQ/dV curves (Figure S6a, b), where the Sc doped sample exhibits oxidation/reduction peaks near 2.7 V and 4.1 V, similar to the undoped material but with reduced intensity. In contrast, the introduction of Ti and B significantly modifies the voltage profile. The low-voltage plateau becomes more pronounced during both charge and discharge, while the high-voltage plateau transforms into a sloping region during charge and nearly disappears during discharge. These changes are also evident in the corresponding dQ/dV curves of the undoped and Sc–Ti–B doped materials (Fig. 2 b). The above observations highlight the distinct roles of each dopant in influencing redox activity, which could contribute to differences in cycling stability among the various cathode materials. All samples were tested at 0.1C, 1C, and 3C (1C = 200 mA g − 1 ) to evaluate the influence of dopants on cycling performance across different charge-discharge rates. As shown in Fig. 2 c, the doped materials exhibit significantly better cycling stability at 0.1C compared to the rapid capacity fading observed in the undoped sample, particularly the Sc–Ti–B doped material. The corresponding charge-discharge curves for all materials are presented in Figure S7. Upon cycling, the low-voltage and high-voltage plateaus vanish quickly for most samples, except for the O3–ScTiB doped material, which retains the low-voltage plateau even after 100 cycles. This suggests a highly reversible electrochemical process enabled by the Sc, Ti, and B co-doping. When cycled at 1C, the Sc–Ti–B co-doped material delivers a high reversible capacity of approximately 147 mAh g⁻¹ and maintains 65% capacity retention after 300 cycles, outperforming both the O3–ScTi sample (60.5%) and the O3–Sc–doped sample (23.2%) (see Fig. 2 d). It is evident that the Sc–Ti–B co-doped sample would exhibit superior capacity retention with further cycling. Unexpectedly, the Sc–Ti co-doped material demonstrates the best cycling stability at a higher rate of 3C, achieving a capacity of about 106 mAh g⁻¹ with 80.1% retention after 500 cycles (see Fig. 2 e and S8). These results indicate that while Sc–Ti–B co-doping significantly enhances cycling stability, its relative advantage becomes less pronounced under higher current densities. As illustrated in Fig. 2 f, the discharge capacities at all cycling currents increase progressively from Sc doping alone to Sc–Ti–B co-doping, with the most prominent improvement observed at 1C. Moreover, B incorporation distinctly improves capacity retention, except under the 3C condition. Furthermore, the rate performance (Figure S9 and S10) reveals that Sc–Ti–B doping enhances the rate capability at lower charge-discharge rates compared to Sc–Ti doping alone. These results indicate that B doping is beneficial for cycling stability at low-current densities but may limit performance at higher-current densities. Redox activity analysis of the first cycle X–ray absorption spectroscopy (XAS) at both the Ni K-edge and L-edge was conducted to investigate the redox behavior and local structural changes of the doped materials at different states of charge: open-circuit voltage (OCV), charged to 4.0 V, charged to 4.4 V, and discharged to 2.5 V (see Fig. 3 and S11–S13). For detailed results on the undoped material, please refer to our recent report 7 . As shown in Fig. 3 a–c and S11, a clear energy shift toward higher values is observed at both the Ni K-edge and L-edge during charging, confirming the oxidation of Ni. At the fully charged state (4.4 V), the O3–ScTi doping sample exhibits the highest edge energy (see inset in Fig. 3 a), followed by the O3–ScTiB doping sample, with the undoped and Sc doped samples showing lower edge energies. These trends suggest that Ti doping enhances Ni redox activity, while B doping slightly suppresses it. Upon discharge to 2.5 V, all doped samples show good reversibility, as evidenced by the near–complete overlap of the spectra with those at the OCV state. In contrast, the extended X–ray absorption fine structure (EXAFS) region at the Ni K-edge reveals a more complex behavior, indicating changes in the local coordination environment of Ni during cycling. As shown in Fig. 3 c, d, all samples exhibit two primary peaks: the first at lower radial distance corresponding to the Ni–O shell, and the second at higher radial distance representing the Ni–Me (B, Sc, Ni, and Mn) shell 7 . During charging, the Ni–O bond length gradually decreases for both the Sc– and Sc–Ti–B doped samples. Notably, the Sc doped sample exhibits a more pronounced contraction, indicating greater distortion in the Ni–O bonds compared to the Sc–Ti–B doped sample. Although the changes in peak intensities resemble those observed in the undoped sample, the overall amplitude variations are smaller in the doped materials, especially in the Sc–Ti–B sample, suggesting improved structural stability. At the discharge state, the Sc–Ti–B doped sample demonstrates superior reversibility, with its EXAFS spectra at OCV and 2.5 V being nearly identical. Figure 3 b presents the EXAFS data for all samples at 4.4 V. The Sc–Ti doped sample exhibits the shortest Ni–O bond length among the materials investigated. At the OCV state, however, its Ni–O radial distance is nearly identical to that of the Sc–Ti–B doped sample, suggesting that B incorporation suppresses the variation in the local Ni environment. Such mitigation of structural distortion is likely to promote improved lattice stability and, consequently, enhanced electrochemical performance. Redox activity of the lattice oxygen (“oxygen redox”) is a complex and highly discussed process, which can be observed particularly in Ni–containing layered oxides operating under high-voltage conditions 31 , 32 . In this study, we employed resonant inelastic X–ray scattering (RIXS) to investigate the influence of various dopants on the oxygen redox behavior. While the origin of oxygen redox signals detected by RIXS remains a topic of ongoing debate, our focus here is not on the mechanism itself but rather on the quantitative variations associated with different dopants 12 , 33 . As shown in Fig. 3 e, distinct differences are observed at the ~ 8 eV energy loss, which is widely considered to represent the spectral fingerprint of oxygen redox 12 , 31 , 33 – 35 . Compared to the O3–undoped sample at the OCV state, all doped samples exhibit similar peaks, albeit with varying intensities. These results confirm that oxygen redox activity is present across all samples to varying extents. To estimate the relative contribution of oxygen redox, the A/B intensity ratios were used and are presented in Fig. 3 h (top). Clearly, all doped materials exhibit lower oxygen redox intensity compared to the undoped sample, likely due to reduced Ni content in the doped compositions. Among the doped materials, the Sc–Ti co-doped sample displays the weakest oxygen redox signals, which aligns with its largest Ni–O bond lengths in the EXAFS (see Fig. 3 b). Interestingly, the 2% B doping (O3–ScTiB doping) markedly enhances oxygen redox activities compared to the Sc–Ti doped counterpart. However, boron is a metalloid, differing fundamentally from metallic dopants such as Li, Mg, Ni, or Zn, which are traditionally associated with triggering oxygen redox by forming the X–O–Na configurations (where X refers to the aforementioned dopants) 36 – 39 . This distinction implies that boron may play a unique and different role in modulating oxygen redox behavior. To further investigate the oxygen redox behavior among these samples during the initial cycle, O K-edge XAS measurements were conducted at various electrochemical states. In the O K-edge spectra (probing the 1s → 2p transition), delocalized states arising from the hybridization between oxygen 2p and Ni/Mn 3d (e g ) orbitals appear within the 527–534 eV range 34 , 40 – 42 . This region is divided into three distinct zones: Region C (527.5–529 eV), Region D (< 530.5 eV), and Region E (< 532.7 eV). In Region C, shoulder features emerge at both the 4.0 V and 4.4 V charge states, corresponding to the formation of oxygen holes previously reported in Na-layered oxides during desodiation 12 , 34 , 37 , 43 . The peak in Region D originates from O 2p states hybridized with unoccupied TM 3d orbitals of spin-down t₂ g and spin-up e g character, reflecting contributions from both exchange-split subbands. The higher energy feature in Region E corresponds to hybridization with spin-down e g states 42 . As shown in Figs. 3 f, 3 g, and S14, S15, all samples exhibit a similar spectral evolution upon charging: the spectra shift toward lower energies, and a new peak appears at 528.6 eV at the 4.0 V charge state. At full charge (4.4 V), the intensity of the 529.4 eV peak in Region D diminishes, indicating that oxygen actively participates in the electrochemical process starting from 4.0 V, with the formation of oxygen holes. Upon discharge, the spectra largely revert to their initial state, suggesting that the oxygen holes are only present at high voltages (> 4.0 V). To quantify the extent of oxygen redox among the samples, the change in the C/D intensity ratio between the 4.0 V and 4.4 V charge states was analyzed and is presented in Fig. 3 h (bottom). Notably, the Δ(C/D) trend closely mirrors the A/B ratio obtained from RIXS measurements, where the intensity decreases from the undoped sample to the Sc and Ti co-doped case, and then increases with B incorporation. Overall, the RIXS and O K-edge XAS data reveal that Sc and Ti co-doping significantly suppresses oxygen redox, whereas B doping reverses this effect and markedly enhances oxygen redox activity. Structural degradation during the first cycle at a low-current density (0.1C) Synchrotron–based operando XRD was employed to study how various dopants influence the structural evolution of O3 type cathodes during the first cycle. For the undoped O3–NaNi 0.5 Mn 0.5 O 2 sample, our previous study revealed complex but reversible structural transformations, following the sequence: O3 (rhombohedral) → O'3 (monoclinic) → P'3 (monoclinic) and P3 (hexagonal) → O'3 (monoclinic) → O3 (rhombohedral), which results in more than 20% lattice contraction at the fully charged state 7 . With doping, all samples exhibit similar trends for changes in the XRD reflections below 4.0 V (see Fig. 4 a, b, and S16–S18). Specifically, the 003 and 006 reflections shift to lower angles and the reflections of 101, 012 and 104 are found to shift toward higher angles, indicating an expansion along the c axis and reduction in the ab plane 7 , 44 . When being charged to 3.0 V, a clear phase transition from O–type to P–type structure is observed, as evidenced by the appearance of 015 P3 at about 5.8° (λ = 0.20733 Å) and 10.1° (λ = 0.35424 Å). Subsequently, all doped materials remain in a single P3 phase up to 4.0 V. Notably, the Sc–Ti and Sc–Ti–B co-doped materials demonstrate a delayed transition to the P3 phase compared to the sample doped with Sc alone, suggesting enhanced structural stability or kinetic hindrance introduced by co-doping. At voltages above 4.0 V, the doped materials exhibit distinct phase transitions. For the Sc doped sample, a similar transformation to that of the undoped counterpart is observed. The main 003 P3 reflection shifts slightly to higher angles, indicating the formation of a monoclinic O'3 phase 8 . This is corroborated by the disappearance of the 012 P3 and 015 P3 reflections. Concurrently, a reflection at 4.8° becomes prominent, indexed as the 100 O'3 reflection and highlighted by the magenta dashed box in Fig. 4 a, further supporting the phase transition. These features closely mirror those seen in the undoped sample. Upon further charging from 4.2 V to 4.4 V, most reflections vanish except for the persistent 100 O'3 reflection, suggesting that the O'3 phase remains dominant in the high-voltage region. Meanwhile, a new reflection appears at 2.4°, and gradually shifts to lower angles upon discharge. This behavior indicates the formation of a new phase, resembling the recently reported OP4 phase in 10% Mg doped P2–Na 0.67 Ni 0.33 Mn 0.67 O 2 31,45,46 . In this study, the newly formed phase is identified as an OP2–type structure 46 . In the Sc and Ti co-doped sample, the 003 P3 reflection significantly weakens at 4.0 V and becomes faint at 4.2 V (see Figures S17 and S19). As highlighted in the 003 reflection in Figure S16b, the diffraction pattern still shows a dominant O'3 phase, along with a broad reflection at 2.0°. This new reflection is similar to the reported P1 phase in NaNi 1/3 Mn 1/3 Co 1/3 O 2 at the desodiated state 47 . These results indicate that both Sc–only doping and Sc–Ti co-doping promote the formation of a P–O intergrowth phase, with the O phase remaining predominant. For the Sc–Ti–B co-doped material, a more continuous structural evolution is observed in the high-voltage region. Upon charging to 4.0 V, the main 003 reflection progressively shifts to higher angles, reaching its maximum 2θ value at the fully charged state. Like the Sc–only doped material, this phase is also an OP2 phase. During the discharge process, the Sc doped sample exhibits only partial reversibility, as indicated by the persistence of mixed phases from 3.7 V down to the fully discharged state. In contrast, the Sc–Ti–B co-doped material shows a highly reversible structural evolution, reflecting improved cycling stability under low-current cycling. These findings suggest that Sc/Ti co-doping effectively suppresses the formation of intermediate phases, while B doping promotes a more continuous and reversible structural transformation in the high-voltage region. To further elucidate the effects of different dopants on structural evolution, the c lattice parameter evolution and phase transition behaviors are illustrated in Fig. 4 c. All samples exhibit lattice expansion up to approximately 4.0 V, followed by contraction at the fully charged state. Among them, the Sc–Ti–B co-doped sample shows the largest expansion ( ~ + 6.37%), even exceeding that of the undoped material. At full charge, all materials exhibit mixed phases except for the Sc–Ti–B doped material (see Figure S20 and S21). The undoped material primarily demonstrates the O3 phase with a significant number of vacancies, along with traces of the O′3 phase. For the Sc– and Sc–Ti doped cases, both are dominated by the O′3 phase, with a small amount of the OP2 phase present in the Sc doped sample and the P1 phase in the Sc–Ti doped sample. In contrast, the Sc–Ti–B co-doped sample displays a single OP2 phase, showing an intermediate level of c lattice parameter changes with the lowest ratio of O phase stacking. This aligns with the observation that a dominant P phase is favorable for better electrochemical performance in the O3–type framework 20 . The phase transition diagram further illustrates this behavior, showing that 3.0 V and 4.0 V act as key boundaries between P–type and O–type phases in O3–type materials. Dopants mainly affect the structure in the high-desodiation state, ultimately influencing structural reversibility during cycling. In summary, Sc and Ti co-doping effectively suppresses the formation of the intermediate O′3 phase, while B doping delays the O3–P3 transition and improves structural reversibility through the stabilization of an OP2 phase at high voltage. HAADF-STEM was employed to investigate the microstructural evolution at different states. As shown in Fig. 5 a, d, O3–undoped and O3–ScTiB doping samples maintain a well-defined layered structure in the bulk, as evidenced by distinct lattice fringes and sharp, regular diffraction spots in the corresponding fast Fourier transform (FFT) patterns. However, in the near-surface region, the O3–undoped sample develops a pronounced cation-mixed layer that is ~ 7 nm thick, whereas doping effectively suppresses the formation of such a disordered region. At the fully charged state, the O3–undoped sample exhibits extensive cracking within the lattice that propagate throughout the crystalline domains, accompanied by significant lattice strain, particularly in the areas near the cracks. In contrast, the Sc–Ti–B co-doped sample preserves a well-ordered layered structure with minimal dislocations and markedly reduced strain in the fully charged state (Fig. 5 b, c, and 5 e, f). These observations are consistent with the smaller contraction of the c lattice parameter in the doped materials. Furthermore, electron energy loss spectroscopy (EELS) at the O K-edge and Ni/Mn L-edges was performed at both the OCV and fully charged states. Both materials exhibit a clear O K-edge signal at the OCV state. However, at 4.4 V, the undoped sample shows a significantly weakened oxygen signal near the surface compared with the Sc–Ti–B co-doped sample. This observation highlights the structural instability of the undoped material in the absence of doping (Figs. 5 g, h, and S22, S23). Taken together, the operando XRD, STEM, and EELS results confirm that lattice contraction-induced strain and oxygen loss are the primary factors driving material degradation in high-voltage cathodes under low-current conditions. Redox degradation during cycling at low current density (0.1C) XAS at the Ni L-edge, Ni K-edge, and O K-edge was performed for both undoped and doped materials to study the redox degradation at 0.1C during cycling. As shown in Figs. 6 a and S24, S25, all samples exhibit noticeable energy shifts in the Ni K-edge XANES spectra over multiple cycles. For the undoped material, the absorption edge gradually shifts to higher energy, indicating progressive oxidation. In contrast, the doped samples show shifts toward lower energies, and the magnitude of these shifts diminishes with continued cycling, especially for the Sc–Ti–B doped sample. In the EXAFS spectra, a clear shift in the Ni–O shell is observed for the undoped and Sc–Ti doped materials, suggesting an unstable coordination environment. In comparison, the O3–ScTiB doped sample maintains a consistent Ni–O radial distance throughout cycling, indicating a more stable local structure. As illustrated in Fig. 3 f, the shoulder feature associated with O hole contributions diminishes during cycling for the undoped sample. However, this change is minimal in the Sc–Ti–B doped material, with nearly no variation between the 1st and 15th cycles. To better compare the effects of doping on O redox behavior, summaries are provided in Fig. 6 e for O3–undoped, O3–ScTi doping, and O3–ScTiB doping, respectively. Notably, in the undoped sample, Ni oxidation increases while O hole contributions decrease over 30 cycles. This suggests a growing reliance on Ni redox activity and a declining contribution from O redox, consistent with the charge-discharge curves shown in Figure S7. In doped materials, both Ni and O redox activities gradually decrease over cycling. However, the O3–ScTiB doping sample demonstrates superior stability, particularly in maintaining O redox contributions. Furthermore, it retains a stable Ni–O coordination environment, with negligible changes in radial distance during cycling. These findings align with the enhanced cycling stability observed at low current rates. To further clarify the influence of dopants on oxygen redox, DFT calculations were carried out for O3–undoped, O3–Sc-doped, O3–ScTi-doped, and O3–ScTiB-doped samples. As shown in Figs. 6 g and S26, the computational results reveal the same trend as observed experimentally: oxygen redox activities decrease from the undoped case to the Sc–Ti co-doped case, and then increases with B incorporation. This behavior arises from the stronger orbital overlap between O 2p and Ni 3d t 2g bonding orbitals induced by Sc and Ti, which stabilizes electrons in the bonding states and suppresses oxygen redox activities. B doping further enhances the O 2p–Ni 3d t 2g overlap to stabilize oxygen, but at the same time modulates orbital energy levels and electron distribution by increasing the occupation of Ni 3d e g antibonding orbitals. This adjustment leads to an overall enhancement of oxygen redox activities. Collectively, the synergistic contributions of Sc, Ti, and B produce an optimized electronic configuration of oxygen that suppresses side reactions such as oxygen loss to ensure stability, while enabling highly reversible redox with maximal electron transfer during electrochemical cycling. Such electronic-level regulation markedly reinforces oxygen redox activity in Ni–Mn layered oxides, in excellent agreement with experimental results. Building on these investigations at low-current densities, a comparison of the electrochemical performance, structural evolution, and oxygen stability is summarized in the spider chart shown in Fig. 6 h. For the undoped material, structural instability and oxygen loss are the primary contributors to degradation. In contrast, doped materials, whether with single or multiple dopants, exhibit similar structural effects, effectively mitigating lattice strain compared to the undoped O3–NaNi 0.5 Mn 0.5 O 2 . In these cases, oxygen stability becomes the dominant factor governing long-term cycling performance at the low-current densities. It is worth emphasizing, however, that oxygen loss is inevitable in high-voltage cathodes; the role of doping is to slow down this process rather than eliminate it. Structure degradation at high current densities At high cycling currents (> 1C), the capacity contribution from oxygen redox becomes negligible (see rate performance in Figure S9). This is attributed to the intrinsically sluggish kinetics of oxygen redox during cycling (see GITT data in Figure S27), suggesting that oxygen stability plays a much less critical role at high currents compared to low-current cycling. To elucidate the dominant factor under fast de/intercalation process, synchrotron–based operando XRD was employed to track the structural evolution. As shown in Fig. 7 a, b and S28, the phase transition pathway becomes markedly simplified at high-current densities. In the Sc–Ti doped sample, a clear O3–P3 phase transition occurs during the first charge. During subsequent cycles at both 1C and 3C, the material stabilizes in a single P3 phase, showing only a minimal shift in the 2θ angle, corresponding to a small change of ~ + 1.5% in the c lattice parameter. In contrast, the B doped sample exhibits a more complex transition pathway. At 1C, it undergoes reversible O3 → P3 → OP2 transitions, similar to those observed at 0.1C. When cycled at 3C, however, a distinct X phase emerges at the fully charged state, following the sequence O3 → P3 → X. This process becomes irreversible during discharge, as the material only transitions back from X to P3. Notably, the Sc–Ti–B co-doped sample shows much larger variations in the c lattice parameter, with contractions of -7.4% at 1C and − 1.7% at 3C, compared to the Sc–Ti doped sample. This greater structural fluctuation is likely a key factor underlying its inferior capacity retention under high-current cycling conditions. Discussion At low current densities, the oxygen redox reaction can fully participate in reversible charge compensation, which increases its contribution to the overall capacity. For the Sc–Ti–B co-doped sample, the additional boron induces a larger change in the c lattice parameter. This effect is attributed to the stronger B–O bond, which effectively suppresses oxygen loss. 26 This stabilization enhances the reversibility of both structural evolution and the oxygen redox process. As a result, the B doped material delivers higher capacity and improved cycling stability compared to its Sc–Ti doped counterpart. These findings suggest that at low cycling currents, long-term cycling performance is primarily controlled by the stability of oxygen redox, and that suppressing oxygen loss is more decisive than mitigating high voltage phase transitions. However, as the cycling current increases, the behavior changes: regardless of dopant type, the material primarily undergoes a simple P3 phase transition during cycling (see Fig. 7 c). Under these rapid charge-discharge conditions, the contribution of oxygen redox to capacity becomes negligible, while structural changes emerge as the dominant factor governing cycling performance. Consequently, the Sc–Ti doped sample exhibits superior cycling stability at high-current densities. Furthermore, a spider chart summarizing the respective roles of TM redox, oxygen redox, and changes in the c lattice parameter on the capacity and cycling stability of high-voltage layered cathodes corroborates this trend (see Fig. 7 d). It highlights that oxygen redox is more critical for determining capacity and stability at low currents, whereas structural evolution exerts a stronger influence on cycling performance under high-current cycling. In summary, we have synthesized a series of Sc–, Sc–Ti–, and Sc–Ti–B co-doped O3–type Ni-Mn based layered oxides via a solid-state reaction method. Notably, the Sc–Ti–B co-doping significantly enhances the Na storage performance at low cycling currents, achieving a capacity of 147 mAh g − 1 at 1C with a capacity retention of 83.5% after 100 cycles in the voltage range of 2.5 to 4.4 V. Using synchrotron-based techniques such as XRD, XAS, RIXS, together with electron microscopy and DFT calculations, it is revealed that the additional B doping effectively suppresses oxygen loss, thereby improving the reversibility of both structural evolution and the oxygen redox process at low-current densities. However, at high cycling current (3C), the Sc–Ti doped material (without B) demonstrates superior capacity retention. This indicates that structural integrity plays a more critical role in capacity retention than oxygen redox stability under high-current cycling. These findings uncover a rate-dependent degradation mechanism in layered oxides under high-voltage operation, in which oxygen stability is the key to achieving high capacity and cycling stability at low currents, whereas structural stability becomes the primary factor shaping degradation at high currents. This current dependent structural and redox behaviour is particularly relevant for practical sodium ion batteries, because real drive cycles impose widely varying current densities. Cells operated under such conditions may age through mechanisms that differ from those inferred from constant current laboratory tests, and the framework proposed here provides guidance for developing high voltage layered cathodes that remain stable under realistic cycling profiles. Methods Synthesis methods The synthesis procedure for all O3–type cathodes followed the method reported previously 7 . The starting materials included Na 2 O 2 (> 97%, Sigma-Aldrich), Mn 2 O 3 (> 99%, Sigma-Aldrich), NiO (> 99%, Sigma-Aldrich), Sc 2 O 3 (> 99.9%, Sigma-Aldrich), TiO 2 (> 99.9%, Sigma-Aldrich), and B 2 O 3 (> 99.98%, Sigma-Aldrich). An excess of 5 mol% Na 2 O 2 was added to compensate for Na loss during the high temperature reaction, and 0.015 mol of product was synthesized for each material. All raw materials were mixed for 5 h using a planetary ball mill with a reversible procedure, then the mixture was fabricated as a pellet by pressing under 15 MPa pressure. After that, the pellet was calcined at 950°C for 15 h under O 2 atmosphere to obtain the final products. Before cooling down to 200°C, the pellet was transferred to an Ar-filled Glovebox for further processing and measurement. Electrochemical characterization The process of fabricating working electrodes was the same as previously reported 7 . The electrode composition consisted of active material, Super P (SP, MTI Corp.), and polyvinylidene difluoride (PVDF, MTI Corp.) at a weight ratio of 80:10:10. The PVDF binder was dissolved in N-methyl-2-pyrrolidone (NMP, Sigma Aldrich) at a concentration of 5 wt%. The active material and Super P were first mixed in a shaker ball mill for 10 min. The PVDF solution was then added, and the mixture was further milled for 30 min to obtain a homogeneous slurry. This slurry was cast onto carbon-coated Al foil (MTI Corp.) using a doctor blade with a wet thickness of 250 µm. All fabrication steps were conducted in an Ar-filled glovebox. The coated electrodes were dried on a hot plate, punched into 12 mm discs, and subsequently vacuum-dried at 120°C for 12 h in a Büchi oven. The electrolyte consisted of 1 M NaClO 4 dissolved in a mixture of propylene carbonate (PC, Sigma Aldrich) and fluoroethylene carbonate (FEC, Sigma Aldrich) with a volume ratio of 98:2. Glass fiber (Whatman) was used as the separator. Coin-type 2025 cells were assembled in an Ar-filled glovebox, with sodium metal (BASF Corp.) serving as the counter electrode. Electrochemical tests were performed in the voltage range of 2.5–4.4 V (vs. Na + /Na) under constant-current conditions at various current rates using a Neware battery test system (CT-4008-5V10mA-164, Neware). For GITT measurements, each step consisted of charging/discharging at 0.1C for 30 min, followed by a rest period of 10 h. Structural characterization Powder XRD was conducted using a Bruker D8 Twin-Twin diffractometer equipped with a Cu Kα source (λ = 1.5406 Å), operated at 40 kV and 40 mA over the range of 10–70° (2θ) with a scanning step of 0.01° s − 1 . Operando XRD measurements for O3–NaSc 0.1 Ni 0.4 Mn 0.5 O 2 and O3–NaSc 0.1 Ni 0.4 Mn 0.4 Ti 0.1 O 2 , along with ex-situ PDF measurements, were performed at beamline P02.1 of PETRA III at DESY (Deutsches Elektronensynchrotron), Hamburg, Germany, using synchrotron radiation with an energy of ~ 60 keV 49 , 50 . In-house designed operando cells were assembled in an Ar-filled glovebox (H 2 O and O 2 concentrations < 5 ppm, as provided by the beamline). Electrochemical measurements were carried out with a Biologic VMP-3 potentiostat controlled by EC-Lab software (v11.34). Prior to cycling, the cells were rested for 20 min, followed by charge-discharge at 0.1 C within the voltage window of 2.5–4.4 V. Operando XRD data were collected sequentially from up to five cells operating simultaneously in Debye-Scherrer geometry, with an acquisition time of 1 min per cell. Data collection was automated using a motorized XY stage with a 350 mm travel range in both directions, controlled by a Python3 script to synchronize stage movement and data acquisition. Consequently, each operando cell was measured for ~ 1 min approximately every 6 min. Operando XRD measurement for O3–NaSc 0.10 Ni 0.4 Mn 0.4 Ti 0.08 B 0.02 O 2 were carried out at the DanMAX beamline of Max IV, Lund, Sweden. The pouch cells were mounted in a custom-designed holder capable of accommodating up to six cells simultaneously 51 . Data were collected in transmission mode using a Dectris Pilatus3 2M CdTe detector, with an acquisition interval of 2 min per pouch cell. The X-ray wavelength was calibrated to λ = 0.35424 Å using a LaB 6 standard. All Rietveld refinements of the diffraction patterns were performed with Fullprof 52 . Spectroscopic characterization Ex situ Ni and Mn K-edge XAS measurements were performed at room temperature at the KMC-2 end-station of BESSY II (Berlin, Germany) in both transmission and fluorescence modes 53 . Reference spectra for energy calibration were collected before and after each sample. XAS spectra were processed using Athena software by subtracting the pre-edge background and normalizing with a spline fit. The k 2 -weighted EXAFS was Fourier transformed over the k-range of 3–11 Å −1 (for general analysis). RIXS measurements were performed at room temperature under a vacuum of ~ 10 − 8 mbar at the U41–PEAXIS end-station of BESSY II (Berlin, Germany) 54 . The spectrometer was aligned in specular geometry at a 60° scattering angle and optimized to a total energy resolution of 90 meV using carbon tape. The excitation energy for each 1D spectrum was set to the O K-edge (531 eV), with an acquisition time of 20 min. For sample preparation, coin cells were charged/discharged to the target voltage, disassembled, and the electrodes were washed with DMC, dried, and mounted on double-sided Cu tape inside a N 2 -filled glovebox. The samples were then transferred to the measurement chamber using a N 2 -sealed transfer suitcase, ensuring no exposure to air. Data were processed using the beamline software, Adler-4.0. Soft XAS measurements were carried out at the Flexible PhotoElectron Spectroscopy (FlexPES) beamline of the MAX IV synchrotron (Lund, Sweden) under ultra-high vacuum at room temperature 55 . An exit slit of 10 µm was used, providing a photon flux of ~ 10 12 photons s − 1 and an energy resolution of 26 meV. The beam spot at the sample was 1 × 0.5 mm 2 (defocused mode). Data were collected in TFY detection mode, and all spectra were background-subtracted and normalized. SEM, STEM and EELS measurements SEM images were taken with a Phenom Pharos Desktop SEM from Phenom world using an accelerating voltage of 10 kV and a secondary electron detector. The transmission electron microscope (TEM) images were collected on a JEOL JEM-3200FS operating at an accelerating voltage of 300 kV. The electron energy loss spectroscopy (EELS) was performed and atomically resolved high angle annular dark-field scanning transmission electron microscopy (HAADF-STEM) images were carried out on a double Cs-corrected TEM (FEI, Titan Themis G2) operated at 300 kV. TEM samples were prepared by FEI Scios microscope operated at 2–30 kV. DFT calculations All calculations were performed within the framework of Density Functional Theory (DFT) using the Vienna Ab initio Simulation Package (VASP) 56 . The Projector-Augmented Wave (PAW) 57 method was employed to describe the electron-nucleus interaction potential, while the Perdew-Burke-Ernzerhof (PBE) functional 58 was adopted for the exchange-correlation potential. A cutoff energy of 500 eV was set for the plane-wave basis set, and the DFT-D3 correction 59 was incorporated to account for weak intermolecular interactions. Structural optimizations were carried out for four crystal models, namely undoped, Sc doping, ScTi doping, and ScTiB doping systems, to ensure the models were in a thermodynamically stable state. For the electronic step calculations, the system was considered to have achieved energy convergence when the energy change was less than 10 − 5 eV. During the structural optimization process, the current configuration was regarded as a minimum point on the potential energy surface when the average atomic force was less than 0.05 eV/Å. The Partial Density of States (PDOS) of the optimized systems was calculated, and visualization of the results was conducted using the VESTA 60 . Declarations Acknowledgements H.L. and Y.L. acknowledge the support from the Swedish Energy Agency (P2022-00055 and P2023-00603) and STandUP for Energy. P.A. and K.A.M. thank the Bundesministerium für Bildung und Forschung (BMBF) for funding over project KAROFEST (03XP0498A) and TRANSITION (03XP0533C). P.A. also acknowledges support from the German Research Foundation (DFG, grant no. 501562980). J.W. and W.Z. are grateful for the financial support from the National Natural Science Foundation of China (52371225 and 92472115). The authors acknowledge DESY (Hamburg, Germany), a member of the Helmholtz Association HGF, for the provision of experimental facilities and beamtime (proposals I-20220574, I-20211173) with the kind support of Dr. Martin Aaskov Karlsen and Dr. Volodymyr Baran. This work was performed at DESY/PETRA III beamline P02.1, and Y.L. and Y.S. also acknowledge DESY for support with travel costs. Soft XAS and part of the operando XRD experiments were carried out at the FlexPES and DanMAX beamlines (Proposals no. 20241703 and 20241814), MAX IV Laboratory, Sweden. The authors gratefully acknowledge Dr. Jiefang Zhu, Dr. Dan Li, Mohammad Baghban Shemirani and Hao Wang from Uppsala University, as well as Dr. Frederik Holm Gjørup and Dr. Eleanor Frampton (MAX IV) for their valuable support during the beamtime. Hard XAS and RIXS were carried out at the KMC-2 and U41-PEAXIS beamlines at the BESSY II electron storage ring operated by the Helmholtz-Zentrum Berlin für Materialien und Energie, GmbH, which is also acknowledged for provision of beamtime. We acknowledge Myfab Uppsala for providing facilities and experimental support. Myfab is funded by the Swedish Research Council (2020-00207) as a national research infrastructure. Author Contributions Y.L., P.A. and H.L. conceived the idea, Y.L. synthesized and characterized the materials, performed Rietveld refinement studies, and all electrochemical measurements. Operando XRD measurements were done by Y.L. and K.A.M. PDF measurements were performed by Y.L. and Y.S. Hard XAS experiments were performed by Y.L., K.A.M., Y.S., C.B., and M.Y. with the help of G.S. RIXS measurements were done by D.W. and Y.L. DFT was calculated by Q.Z. and J.W. STEM was performed by W.Z. H. Z., and J.W. The data and manuscript were organized by L.Y. and H.L. The manuscript was written by Y.L., H.L., J.W. and P.A. with input from all authors. All authors contributed to discussion of the results. References Nayak PK, Yang L, Brehm W, Adelhelm P (2018) From Lithium-Ion to Sodium-Ion Batteries: Advantages, Challenges, and Surprises. Angew Chem Int Ed 57:102–120 Vaalma C, Buchholz D, Weil M, Passerini S (2018) A cost and resource analysis of sodium-ion batteries. 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Phys Rev Lett 77:3865–3868 Grimme S et al (2010) A consistent and accurate ab initio parametrization of density functional dispersion correction (DFT-D) for the 94 elements H-Pu. J Chem Phys 132:154104 Momma K (2011) VESTA 3for three-dimensional visualization of crystal, volumetric and morphology data. J Appl Cryst 44:1272–1276 Additional Declarations There is NO Competing Interest. Supplementary Files SupportingInformationforIdentifyingcyclingcurrentdependentdegradationpathwaysHL.docx Supplementary Information Cite Share Download PDF Status: Under Review Version 1 posted You are reading this latest preprint version Research Square lets you share your work early, gain feedback from the community, and start making changes to your manuscript prior to peer review in a journal. As a division of Research Square Company, we’re committed to making research communication faster, fairer, and more useful. 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University","correspondingAuthor":false,"prefix":"","firstName":"Mingman","middleName":"","lastName":"Yuan","suffix":""},{"id":558828658,"identity":"564a7feb-6d15-4627-b70b-b04357f8dca3","order_by":11,"name":"Jun Wang","email":"","orcid":"https://orcid.org/0000-0001-9561-5857","institution":"Southern University of Science and Technology","correspondingAuthor":false,"prefix":"","firstName":"Jun","middleName":"","lastName":"Wang","suffix":""},{"id":558828659,"identity":"54c43406-2f49-4c87-8b1f-6357b3ba3c4e","order_by":12,"name":"Philipp Adelhelm","email":"","orcid":"https://orcid.org/0000-0003-2439-8802","institution":"Humboldt-University Berlin and Helmholtz-Zentrum Berlin (HZB)","correspondingAuthor":false,"prefix":"","firstName":"Philipp","middleName":"","lastName":"Adelhelm","suffix":""}],"badges":[],"createdAt":"2025-11-13 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07:51:21","extension":"png","order_by":15,"title":"","display":"","copyAsset":false,"role":"acdc-reference","size":205562,"visible":true,"origin":"","legend":"","description":"","filename":"Onlinefloatimage5.png","url":"https://assets-eu.researchsquare.com/files/rs-8109344/v1/f11b042a0a51f7be3e627ad3.png"},{"id":98426993,"identity":"6240b091-0c20-4b7c-81f2-a3e75d83be73","added_by":"auto","created_at":"2025-12-17 16:39:12","extension":"png","order_by":16,"title":"","display":"","copyAsset":false,"role":"acdc-reference","size":86030,"visible":true,"origin":"","legend":"","description":"","filename":"Onlinefloatimage6.png","url":"https://assets-eu.researchsquare.com/files/rs-8109344/v1/3961de454ca7c9cfd5132975.png"},{"id":98043630,"identity":"6de76711-d2ed-419a-8322-053d1c2616d0","added_by":"auto","created_at":"2025-12-12 07:51:21","extension":"png","order_by":17,"title":"","display":"","copyAsset":false,"role":"acdc-reference","size":94846,"visible":true,"origin":"","legend":"","description":"","filename":"Onlinefloatimage7.png","url":"https://assets-eu.researchsquare.com/files/rs-8109344/v1/9962ef948c4e938c5ce04b5b.png"},{"id":98043632,"identity":"2100c453-4d8a-410c-b47d-18925f488879","added_by":"auto","created_at":"2025-12-12 07:51:21","extension":"xml","order_by":18,"title":"","display":"","copyAsset":false,"role":"acdc-reference","size":144551,"visible":true,"origin":"","legend":"","description":"","filename":"NCOMMS25919140structuring.xml","url":"https://assets-eu.researchsquare.com/files/rs-8109344/v1/94af96e35f159d9c46e82a11.xml"},{"id":98043631,"identity":"75a072e3-dfd5-4ce0-9091-f9b31ab971bb","added_by":"auto","created_at":"2025-12-12 07:51:21","extension":"html","order_by":19,"title":"","display":"","copyAsset":false,"role":"acdc-reference","size":154907,"visible":true,"origin":"","legend":"","description":"","filename":"earlyproof.html","url":"https://assets-eu.researchsquare.com/files/rs-8109344/v1/ebf9d95028127f5ba3ba2ea2.html"},{"id":98427381,"identity":"e275eb3e-12b1-4326-a2d0-7964e0229e20","added_by":"auto","created_at":"2025-12-17 16:40:16","extension":"png","order_by":1,"title":"Figure 1","display":"","copyAsset":false,"role":"figure","size":1001891,"visible":true,"origin":"","legend":"\u003cp\u003e\u003cstrong\u003eStructural characterizations.\u003c/strong\u003e (a) Schematic illustration of the composition design routes and their corresponding abbreviations. A high-entropy doping state readily emerges when more than four elements are introduced as dopants\u003csup\u003e9,11,18\u003c/sup\u003e. b) XRD patterns of O3–undoped, O3–Sc doping, O3–ScTi doping and O3–ScTiB doping. The insert magnifies the 2θ range of 16°–17.5° to highlight the shift of the 003 reflection. (c) XRD Rietveld refinement pattern of O3–ScTiB doping cathode. (d) O 2p orbital projected density of states (pDOS). (e) HAADF-STEM images with interlayer distances of O3-No doping and O3-ScTiB doping. (f) HAADF EDS maps of O3–ScTiB doping.\u003c/p\u003e","description":"","filename":"floatimage1.png","url":"https://assets-eu.researchsquare.com/files/rs-8109344/v1/029c95e68c92d2a0ee96d30e.png"},{"id":98427771,"identity":"7549ea34-9f8e-4b25-bd37-8df14584747a","added_by":"auto","created_at":"2025-12-17 16:41:09","extension":"png","order_by":2,"title":"Figure 2","display":"","copyAsset":false,"role":"figure","size":567853,"visible":true,"origin":"","legend":"\u003cp\u003e\u003cstrong\u003eElectrochemical properties in 2.5–4.4 V.\u003c/strong\u003e (a) First charge-discharge curves of O3–undoped, O3–Sc doping, O3–ScTi doping, and O3–ScTiB doping. (b) Corresponding dQ/dV curves of O3–undoped and O3–ScTiB doping cathodes. (c–e) Cycling performance at 0.1C, 1C and 3C rates, respectively (1C = 200 mA g\u003csup\u003e–1\u003c/sup\u003e). (f) Maximum discharge capacity and capacity retention after 100 cycles at different cycling currents, with dotted lines indicating trend variations.\u003c/p\u003e","description":"","filename":"floatimage2.png","url":"https://assets-eu.researchsquare.com/files/rs-8109344/v1/5e2e18b2c324d501caefc12d.png"},{"id":98043607,"identity":"98ab8ca9-3511-49be-8712-be6128987c67","added_by":"auto","created_at":"2025-12-12 07:51:20","extension":"png","order_by":3,"title":"Figure 3","display":"","copyAsset":false,"role":"figure","size":653155,"visible":true,"origin":"","legend":"\u003cp\u003e\u003cstrong\u003eXAS and RIXS studies.\u003c/strong\u003e (a) Ni K-edge XANES spectra for O3–Sc doping and O3–ScTiB doping at the various charge-discharge states, along with a comparison at the fully charged state. (b–d) Corresponding Ni K-edge EXAFS spectra. (e) O K-edge RIXS spectra with an excitation energy of 531 eV for O3–undoped, O3–Sc doping, O3–ScTi doping and O3–ScTiB doping at the fully charged state. (f, g) O K-edge XAS at the various charge-discharge states for O3–undoped and O3–ScTiB doping. (h) A comparison based on the O K-edge RIXS and XAS of the reactive oxygen amount versus dopant type.\u003c/p\u003e","description":"","filename":"floatimage3.png","url":"https://assets-eu.researchsquare.com/files/rs-8109344/v1/c2c539113f3815cafa6a312f.png"},{"id":98428382,"identity":"538b6789-da7f-4bab-a7b8-99bcd4aad906","added_by":"auto","created_at":"2025-12-17 16:41:58","extension":"png","order_by":4,"title":"Figure 4","display":"","copyAsset":false,"role":"figure","size":696299,"visible":true,"origin":"","legend":"\u003cp\u003e\u003cstrong\u003eSynchrotron-based \u003c/strong\u003e\u003cem\u003e\u003cstrong\u003eoperando\u003c/strong\u003e\u003c/em\u003e\u003cstrong\u003e X–ray diffraction investigations.\u003c/strong\u003e Contour plots along with phase evolution and corresponding \u003cem\u003ec\u003c/em\u003e lattice parameter for (a) O3–Sc doping (λ = 0.20733 Å), and (b) O3–ScTiB doping (λ = 0.35424 Å) during the first cycle at the 0.1C rate. (c) Schematic illustration of doping effects on the phase transition behavior.\u003c/p\u003e","description":"","filename":"floatimage4.png","url":"https://assets-eu.researchsquare.com/files/rs-8109344/v1/4f9c4e5bc4e7b65713af5a59.png"},{"id":98427274,"identity":"4d834b41-305e-4823-bb10-3a8128ad3678","added_by":"auto","created_at":"2025-12-17 16:40:03","extension":"png","order_by":5,"title":"Figure 5","display":"","copyAsset":false,"role":"figure","size":953750,"visible":true,"origin":"","legend":"\u003cp\u003e\u003cstrong\u003eMicrostructural analysis.\u003c/strong\u003e HAADF-STEM images and the corresponding FFT patterns of (a, b) O3–undoped and (d, e) O3–ScTiB doping at the OCV state and after charging to 4.4 V, respectively. Strain maps obtained by geometrical phase analysis\u003csup\u003e48\u003c/sup\u003e (GPA) for fully charged samples: (c) O3–undoped and (f) O3–ScTiB doping. EELS spectra at the O K-edge for (g) O3–undoped and (h) O3–ScTiB doping at the OCV state and after charging to 4.4 V. The red line marks the onset of signal detection at the oxygen K-edge.\u003c/p\u003e","description":"","filename":"floatimage5.png","url":"https://assets-eu.researchsquare.com/files/rs-8109344/v1/5376ed1528243547c0ebe5b7.png"},{"id":98425874,"identity":"36e55cf4-7f14-441c-9a38-598a1503c785","added_by":"auto","created_at":"2025-12-17 16:35:19","extension":"png","order_by":6,"title":"Figure 6","display":"","copyAsset":false,"role":"figure","size":573017,"visible":true,"origin":"","legend":"\u003cp\u003e\u003cstrong\u003eRedox and local structural degradation during the cycling.\u003c/strong\u003e(a) Ni K-edge XANES of O3–undoped and O3–ScTiB doping at the 1st, 15th and 30th cycle under the fully charged state. (b, c) Corresponding EXAFS spectra at the different cycles. (d, e) O K-edge XAS at the different cycles. (f) changes in oxygen redox activity (C/D ratio) at both OCV and fully charged states across cycling. (g) Schematic illustration of the calculated DOS for all cathodes. (h) Spider chart illustrating the various factors affecting electrochemical performance at low cycling rate.\u003c/p\u003e","description":"","filename":"floatimage6.png","url":"https://assets-eu.researchsquare.com/files/rs-8109344/v1/1216f51c8dab6e9717f8a3fd.png"},{"id":98043617,"identity":"ceda9222-a732-4804-bced-45116b0ea924","added_by":"auto","created_at":"2025-12-12 07:51:20","extension":"png","order_by":7,"title":"Figure 7","display":"","copyAsset":false,"role":"figure","size":522121,"visible":true,"origin":"","legend":"\u003cp\u003e\u003cstrong\u003eStructural evolution upon fast charging.\u003c/strong\u003e Synchrotron–based \u003cem\u003eoperando\u003c/em\u003eXRD at 1C and 3C rates for (a) O3–ScTi doping (λ = 0.20733 Å) and (b) O3–ScTiB doping (λ = 0.35424 Å). (c) A comparison of phase transitions and \u003cem\u003ec\u003c/em\u003e lattice parameters at different charge-discharge rates. X denotes the intergrowth phase formed between the P-type and O-type stacking sequences. (d) Spider chart illustrating the various factors affecting electrochemical performance under low (\u0026lt;1C), moderate (1–3C), and high (\u0026gt; 3C) cycling current conditions.\u003c/p\u003e","description":"","filename":"floatimage7.png","url":"https://assets-eu.researchsquare.com/files/rs-8109344/v1/3d8d4ca6c9a50aaef4740451.png"},{"id":98622178,"identity":"c7c33f89-ed14-4288-9a54-bf08c2a1192e","added_by":"auto","created_at":"2025-12-19 16:47:56","extension":"pdf","order_by":0,"title":"","display":"","copyAsset":false,"role":"manuscript-pdf","size":5907115,"visible":true,"origin":"","legend":"","description":"","filename":"manuscript.pdf","url":"https://assets-eu.researchsquare.com/files/rs-8109344/v1/970a3b04-8a03-4e22-b9c5-b783c4700656.pdf"},{"id":98427328,"identity":"956e3ae3-0f1d-468f-8df9-339e64cb0b65","added_by":"auto","created_at":"2025-12-17 16:40:09","extension":"docx","order_by":1,"title":"","display":"","copyAsset":false,"role":"supplement","size":7132348,"visible":true,"origin":"","legend":"Supplementary Information","description":"","filename":"SupportingInformationforIdentifyingcyclingcurrentdependentdegradationpathwaysHL.docx","url":"https://assets-eu.researchsquare.com/files/rs-8109344/v1/fba279cd54914b7737e99590.docx"}],"financialInterests":"There is \u003cb\u003eNO\u003c/b\u003e Competing Interest.","formattedTitle":"Identifying cycling current-dependent degradation pathways in high-voltage layered sodium–ion cathodes","fulltext":[{"header":"Introduction","content":"\u003cp\u003eSodium-ion batteries (NIBs) are a promising alternative to lithium-ion batteries (LIBs) for grid-scale energy storage due to the natural abundance and low cost of sodium\u003csup\u003e\u003cspan additionalcitationids=\"CR2\" citationid=\"CR1\" class=\"CitationRef\"\u003e1\u003c/span\u003e\u0026ndash;\u003cspan citationid=\"CR3\" class=\"CitationRef\"\u003e3\u003c/span\u003e\u003c/sup\u003e. Advancing high-power, durable cathode materials is essential, as the cathode largely determines battery cost and performance\u003csup\u003e\u003cspan citationid=\"CR2\" class=\"CitationRef\"\u003e2\u003c/span\u003e\u003c/sup\u003e. Among various cathodes, O3\u0026ndash;type Ni\u0026ndash;Mn based layered oxides, characterized by Na⁺ ions in octahedral sites and ABC stacking, can offer high capacity (\u0026gt;\u0026thinsp;200 mAh g\u003csup\u003e\u0026minus;\u0026thinsp;1\u003c/sup\u003e above 4.1 V vs. Na\u003csup\u003e+\u003c/sup\u003e/Na) along with high average redox potentials, making them attractive for practical use\u003csup\u003e\u003cspan citationid=\"CR4\" class=\"CitationRef\"\u003e4\u003c/span\u003e,\u003cspan citationid=\"CR5\" class=\"CitationRef\"\u003e5\u003c/span\u003e\u003c/sup\u003e. The high capacity is supported by the activation of oxygen redox processes\u003csup\u003e\u003cspan citationid=\"CR6\" class=\"CitationRef\"\u003e6\u003c/span\u003e\u003c/sup\u003e. Nevertheless, despite this advantage, their performance rapidly deteriorates under fast Na de/intercalation reactions (\u0026gt;\u0026thinsp;3C), manifested by poor rate capability and cycling stability, as widely reported for O3\u0026ndash;type Ni\u0026ndash;Mn based cathodes\u003csup\u003e\u003cspan additionalcitationids=\"CR8\" citationid=\"CR7\" class=\"CitationRef\"\u003e7\u003c/span\u003e\u0026ndash;\u003cspan citationid=\"CR9\" class=\"CitationRef\"\u003e9\u003c/span\u003e\u003c/sup\u003e. Such degradation originates from structural collapse, oxygen loss, cation migration, and inhomogeneous strain accumulation, with strongly rate-dependent pathways that ultimately constrain long-term stability and scalability\u003csup\u003e\u003cspan citationid=\"CR7\" class=\"CitationRef\"\u003e7\u003c/span\u003e,\u003cspan additionalcitationids=\"CR11 CR12 CR13 CR14\" citationid=\"CR10\" class=\"CitationRef\"\u003e10\u003c/span\u003e\u0026ndash;\u003cspan citationid=\"CR15\" class=\"CitationRef\"\u003e15\u003c/span\u003e\u003c/sup\u003e.\u003c/p\u003e \u003cp\u003eDoping/substitution strategies, as well as the more recent concept of entropy guided optimization, represent rational approaches for designing cathodes with high performance, and have been widely employed to mitigate the above-mentioned issues\u003csup\u003e\u003cspan citationid=\"CR9\" class=\"CitationRef\"\u003e9\u003c/span\u003e,\u003cspan citationid=\"CR11\" class=\"CitationRef\"\u003e11\u003c/span\u003e,\u003cspan citationid=\"CR13\" class=\"CitationRef\"\u003e13\u003c/span\u003e,\u003cspan additionalcitationids=\"CR17 CR18 CR19 CR20\" citationid=\"CR16\" class=\"CitationRef\"\u003e16\u003c/span\u003e\u0026ndash;\u003cspan citationid=\"CR21\" class=\"CitationRef\"\u003e21\u003c/span\u003e\u003c/sup\u003e. Even trace amounts of dopants can significantly alter the crystal and electronic structures of layered cathodes, thereby influencing phase transition pathways, and ultimately improving electrochemical properties. Suppressing unfavorable phase transitions, such as the P3\u0026ndash;O3 transformation in NaNi\u003csub\u003e0.5\u003c/sub\u003eMn\u003csub\u003e0.5\u003c/sub\u003eO\u003csub\u003e2\u003c/sub\u003e is widely recognized as a key strategy for enhancing the electrochemical performance of O3\u0026ndash;type layered oxides\u003csup\u003e\u003cspan citationid=\"CR7\" class=\"CitationRef\"\u003e7\u003c/span\u003e\u003c/sup\u003e. However, while essential, phase\u0026ndash;transition suppression alone does not fully determine the long-term stability under high-voltage operation. For example, single-element doping with Li, Mg, Fe, Cu, or Zn has been reported to suppress these unfavorable phase transitions in O3\u0026ndash;type NaNi\u003csub\u003e0.5\u003c/sub\u003eMn\u003csub\u003e0.5\u003c/sub\u003eO\u003csub\u003e2\u003c/sub\u003e by promoting a more reversible structural evolution during low-current cycling\u003csup\u003e\u003cspan citationid=\"CR7\" class=\"CitationRef\"\u003e7\u003c/span\u003e,\u003cspan additionalcitationids=\"CR23 CR24\" citationid=\"CR22\" class=\"CitationRef\"\u003e22\u003c/span\u003e\u0026ndash;\u003cspan citationid=\"CR25\" class=\"CitationRef\"\u003e25\u003c/span\u003e\u003c/sup\u003e. Nevertheless, despite the structural improvements achieved through single-element doping, such materials still experience rapid capacity fading, particularly when cycling under high-current densities. By contrast, cathodes incorporating multiple dopants such as NaCu\u003csub\u003e0.1\u003c/sub\u003eNi\u003csub\u003e0.2\u003c/sub\u003eCo\u003csub\u003e0.2\u003c/sub\u003eFe\u003csub\u003e0.2\u003c/sub\u003eMn\u003csub\u003e0.15\u003c/sub\u003eTi\u003csub\u003e0.15\u003c/sub\u003eO\u003csub\u003e2\u003c/sub\u003e and NaNi\u003csub\u003e0.35\u003c/sub\u003eMn\u003csub\u003e0.35\u003c/sub\u003eCu\u003csub\u003e0.01\u003c/sub\u003eFe\u003csub\u003e0.1\u003c/sub\u003eTi\u003csub\u003e0.05\u003c/sub\u003eSn\u003csub\u003e0.05\u003c/sub\u003eO\u003csub\u003e2\u003c/sub\u003e, deliver remarkable stability at high-current densities despite undergoing similar phase transition processes\u003csup\u003e\u003cspan citationid=\"CR9\" class=\"CitationRef\"\u003e9\u003c/span\u003e,\u003cspan citationid=\"CR13\" class=\"CitationRef\"\u003e13\u003c/span\u003e\u003c/sup\u003e. However, these materials typically deliver lower specific capacities at low -current densities compared with their single-element doped counterparts. These observations indicate a trade-off in which specific capacity is sacrificed to achieve improved rate stability. Yet, this compromise inevitably limits capacity across both low- and high-current operations, underscoring the intrinsic rate-dependent degradation mechanisms of high-voltage layered oxides.\u003c/p\u003e \u003cp\u003eBeyond structural effects, redox activities strongly influence both capacity and cycling stability, with oxygen redox playing a particularly critical role in high-voltage layered oxides. However, oxygen redox is often accompanied by oxygen loss from unstable local environments and resulting structural instabilities, which accelerates degradation and compromises long-term performance\u003csup\u003e\u003cspan citationid=\"CR10\" class=\"CitationRef\"\u003e10\u003c/span\u003e\u003c/sup\u003e. To address this challenge, dopants such as B, S, and Se have been employed to stabilize oxygen redox by enhancing covalent bonding within the lattice\u003csup\u003e\u003cspan additionalcitationids=\"CR27 CR28\" citationid=\"CR26\" class=\"CitationRef\"\u003e26\u003c/span\u003e\u0026ndash;\u003cspan citationid=\"CR29\" class=\"CitationRef\"\u003e29\u003c/span\u003e\u003c/sup\u003e. Nevertheless, most studies have primarily focused on the initial cycles, providing limited direct evidence of oxygen stability during extended operation. Although online/differential electrochemical mass spectrometry (OEMS or DEMS, respectively) has been widely used to probe oxygen redox reversibility and has yielded valuable insights into early-cycle behavior\u003csup\u003e\u003cspan citationid=\"CR21\" class=\"CitationRef\"\u003e21\u003c/span\u003e,\u003cspan additionalcitationids=\"CR27\" citationid=\"CR26\" class=\"CitationRef\"\u003e26\u003c/span\u003e\u0026ndash;\u003cspan citationid=\"CR28\" class=\"CitationRef\"\u003e28\u003c/span\u003e\u003c/sup\u003e, this technique is largely restricted to the first few cycles and may fail to capture additional oxygen loss pathways that emerge during extended cycling. Furthermore, most measurements are conducted at low current densities, which do not adequately reveal the degradation mechanisms that dominate under high-current densities, particularly those associated with structural evolution. Collectively, these findings underscore the presence of rate-dependent degradation pathways which are often underappreciated, and which are critical to advancing the development of stable oxygen\u0026ndash;redox cathodes.\u003c/p\u003e \u003cp\u003eTo tackle both structural and redox instabilities, dopants were chosen based on their distinct yet complementary effects. Scandium (Sc\u003csup\u003e3+\u003c/sup\u003e), with its relatively large ionic radius and stable 3\u0026thinsp;+\u0026thinsp;valence, can act as a structural stabilizer that strengthens the transition-metal (TM)\u0026ndash;O framework and help suppress cation migration, thereby mitigating phase transitions and slab gliding observed in undoped O3-type oxides. Titanium (Ti\u003csup\u003e4+\u003c/sup\u003e) is known for its strong Ti\u0026ndash;O covalency and electronic stability, and can further enhance the robustness of the framework and help reduce Jahn\u0026ndash;Teller distortion by partially replacing Mn\u003csup\u003e3+\u003c/sup\u003e, while moderating irreversible oxygen redox activity at high voltages. In contrast, boron (B\u003csup\u003e3+\u003c/sup\u003e) serves as an electronic structure modifier that strengthens TM\u0026ndash;O\u0026ndash;B covalency, stabilizing oxygen redox processes and suppressing oxygen loss during cycling. The combination of these dopants is therefore expected to simultaneously stabilize the cationic and anionic sublattices, providing a model system to disentangle rate-dependent degradation mechanisms in high-voltage O3-type Na layered oxides.\u003c/p\u003e \u003cp\u003eIn this work, Sc and Ti were introduced to partially substitute Ni and Mn in O3- NaNi\u003csub\u003e0.5\u003c/sub\u003eMn\u003csub\u003e0.5\u003c/sub\u003eO\u003csub\u003e2\u003c/sub\u003e, while a small amount of B (2 mol%) was further incorporated by partially replacing Ti. This compositional tuning enables a systematic investigation of how these dopants modulate transition-metal (TM) and oxygen redox activities, and how they correlate with the degradation mechanisms under varying current densities. It was found that these cathodes exhibit rate-dependent cycling stability, with Sc\u0026ndash;Ti\u0026ndash;B co-doping (NaSc\u003csub\u003e0.1\u003c/sub\u003eNi\u003csub\u003e0.4\u003c/sub\u003eMn\u003csub\u003e0.4\u003c/sub\u003eTi\u003csub\u003e0.08\u003c/sub\u003eB\u003csub\u003e0.02\u003c/sub\u003eO\u003csub\u003e2\u003c/sub\u003e) significantly enhancing both capacity and cycling stability at low currents. The Sc\u0026ndash;Ti\u0026ndash;B co-doped material delivers 147 mAh g\u003csup\u003e\u0026minus;\u0026thinsp;1\u003c/sup\u003e at 1C (200 mA g\u003csup\u003e\u0026ndash;\u003cspan citationid=\"CR1\" class=\"CitationRef\"\u003e1\u003c/span\u003e\u003c/sup\u003e), retaining 83% of its capacity after 100 cycles within a voltage window of 2.5\u0026ndash;4.4 V. However, at a 3C rate, the Sc\u0026ndash;Ti co-doped material (without B) (NaSc\u003csub\u003e0.1\u003c/sub\u003eNi\u003csub\u003e0.4\u003c/sub\u003eMn\u003csub\u003e0.4\u003c/sub\u003eTi\u003csub\u003e0.1\u003c/sub\u003eO\u003csub\u003e2\u003c/sub\u003e) demonstrates superior stability. Using advanced characterization methods, including resonant inelastic X\u0026ndash;ray scattering (RIXS), soft/hard X\u0026ndash;ray absorption spectroscopy (XAS), transmission electron microscopy (TEM), and density functional theory (DFT) calculations, we identify that B doping greatly improves structural stability and oxygen\u0026ndash;redox reversibility at low-current densities, primarily by reducing oxygen loss. These findings indicate that oxygen loss during cycling is a key contributor to capacity fading at low-current densities, beyond structural degradation alone. At high-current densities, by contrast, all O3\u0026ndash;type materials eventually undergo a single\u0026ndash;phase transition to the P3 phase. For the B-containing composition, although improved oxygen stability enhances structural reversibility at low-current densities, it exacerbates structural changes during rapid de/sodiation, thereby compromising performance at high-current densities. These findings reveal two fundamentally distinct degradation pathways, one governed by oxygen redox stability at low current and the other governed by structural robustness at high current, and they provide a unified mechanistic framework that explains high-voltage behaviour in layered sodium-ion cathodes while establishing the basis for developing materials that remain stable under varying cycling conditions.\u003c/p\u003e"},{"header":"Results and discussion","content":"\u003cdiv id=\"Sec3\" class=\"Section2\"\u003e \u003ch2\u003eCompositions of substituted O3 NaNi\u003csub\u003e0.5\u003c/sub\u003eMn\u003csub\u003e0.5\u003c/sub\u003eO\u003csub\u003e2\u003c/sub\u003e\u003c/h2\u003e \u003cp\u003eMotivated by our earlier results, O3\u0026ndash;NaNi\u003csub\u003e0.5\u003c/sub\u003eMn\u003csub\u003e0.5\u003c/sub\u003eO\u003csub\u003e2\u003c/sub\u003e (designated as O3\u0026ndash;undoped) shows only Ni and oxygen redox activity within the voltage range of 2.5\u0026ndash;4.4 V, delivering an initial discharge capacity exceeding 200 mAh g\u003csup\u003e\u0026minus;\u0026thinsp;1\u003c/sup\u003e. However, this material suffers from rapid capacity fading, even with single-element doping such as Mg.\u003csup\u003e7\u003c/sup\u003e To further investigate strategies for stabilizing this high-voltage cathode, various dopants, including Sc, Ti, and B, were incorporated into the parent compound. As shown in Fig.\u0026nbsp;\u003cspan refid=\"Fig1\" class=\"InternalRef\"\u003e1\u003c/span\u003ea, the designed compositions are O3\u0026ndash;NaSc\u003csub\u003e0.10\u003c/sub\u003eNi\u003csub\u003e0.40\u003c/sub\u003eMn\u003csub\u003e0.50\u003c/sub\u003eO\u003csub\u003e2\u003c/sub\u003e (O3\u0026ndash;Sc doping), O3\u0026ndash;NaSc\u003csub\u003e0.10\u003c/sub\u003eNi\u003csub\u003e0.40\u003c/sub\u003eMn\u003csub\u003e0.40\u003c/sub\u003eTi\u003csub\u003e0.10\u003c/sub\u003eO\u003csub\u003e2\u003c/sub\u003e (O3\u0026ndash;ScTi doping), and O3\u0026ndash;NaSc\u003csub\u003e0.10\u003c/sub\u003eNi\u003csub\u003e0.40\u003c/sub\u003eMn\u003csub\u003e0.40\u003c/sub\u003eTi\u003csub\u003e0.08\u003c/sub\u003eB\u003csub\u003e0.02\u003c/sub\u003eO\u003csub\u003e2\u003c/sub\u003e (O3\u0026ndash;ScTiB doping). All samples exhibit a well-defined layered structure, with only a minor presence of NiO impurity (see Fig.\u0026nbsp;\u003cspan refid=\"Fig1\" class=\"InternalRef\"\u003e1\u003c/span\u003ea, \u003cspan refid=\"Fig1\" class=\"InternalRef\"\u003e1\u003c/span\u003eb, S1, S2 and Table \u003cspan refid=\"MOESM1\" class=\"InternalRef\"\u003eS1\u003c/span\u003e\u0026ndash;S4)\u003csup\u003e\u003cspan citationid=\"CR22\" class=\"CitationRef\"\u003e22\u003c/span\u003e,\u003cspan citationid=\"CR23\" class=\"CitationRef\"\u003e23\u003c/span\u003e,\u003cspan citationid=\"CR30\" class=\"CitationRef\"\u003e30\u003c/span\u003e\u003c/sup\u003e. Notably, doping significantly reduces the amount of this impurity, particularly in the Sc\u0026ndash;Ti\u0026ndash;B doped material. Upon Sc doping, the primary (003) reflection is clearly shifted to a lower 2θ angle, indicating an expansion along the \u003cem\u003ec\u003c/em\u003e axis. However, with the introduction of Ti, the (003) reflection shifts back toward a higher 2θ angle, and this trend continues with additional B doping. X\u0026ndash;ray pair distribution function (PDF) analysis further confirms that the dopants are successfully incorporated into the crystal structure, as evidenced by distinct shifts in the nearest-neighbor transition metal\u0026ndash;oxygen (TM\u0026ndash;O, TM\u0026thinsp;=\u0026thinsp;Ni, Mn, and Sc) and transition metal\u0026ndash;transition metal (TM\u0026ndash;TM) distances (see Figure S3). B doping stands out by producing the shortest TM\u0026ndash;O bond length and the longest TM\u0026ndash;TM distance, which may facilitate the formation of covalent bonds to stabilize O\u003csup\u003e26\u003c/sup\u003e. Moreover, the density of O 2p states near the Fermi level increases significantly, indicating that doping introduces more electronic states, which may affect electronic conductivity and oxygen redox activities (see Fig.\u0026nbsp;\u003cspan refid=\"Fig1\" class=\"InternalRef\"\u003e1\u003c/span\u003ed and S4). These observations confirm that each dopant induces distinct modifications in the local structural environment of the layered oxide framework. All samples exhibit a similar morphology, characterized by irregularly shaped particles with sizes up to approximately 3 \u0026micro;m (see Figure S5). High-angle annular dark-field scanning transmission electron microscopy (HAADF-STEM) further confirms that the dopants exert minimal influence on the crystal structure, as evidenced by the similar atomic arrangements and the absence of cation mixing in the bulk regions of both materials (see Fig.\u0026nbsp;\u003cspan refid=\"Fig1\" class=\"InternalRef\"\u003e1\u003c/span\u003ee). The accompanying energy-dispersive X-ray spectroscopy (EDS) analysis also verifies the presence of all constituent elements and reveals their uniform spatial distribution.\u003c/p\u003e \u003cp\u003e \u003c/p\u003e \u003c/div\u003e\n\u003ch3\u003eElectrochemical properties\u003c/h3\u003e\n\u003cp\u003eThe electrochemical properties of all cathode materials were evaluated in half cells within a voltage window of 2.5\u0026ndash;4.4 V versus Na\u003csup\u003e+\u003c/sup\u003e/Na, as shown in Fig.\u0026nbsp;\u003cspan refid=\"Fig2\" class=\"InternalRef\"\u003e2\u003c/span\u003e. Each material exhibits distinct charge-discharge capacities and unique voltage curves during the initial cycle (see Fig.\u0026nbsp;\u003cspan refid=\"Fig2\" class=\"InternalRef\"\u003e2\u003c/span\u003ea). The specific charge-discharge capacities for O3\u0026ndash;undoped, O3\u0026ndash;Sc doping, O3\u0026ndash;ScTi doping, and O3\u0026ndash;ScTiB doping are 233.3/197.0, 173.9/158.3, 191.1/159.1, and 200.2/169.8 mAh g\u003csup\u003e\u0026minus;\u0026thinsp;1\u003c/sup\u003e, respectively. The undoped material (O3\u0026ndash;undoped) displays multiple small voltage plateaus between 3.0 and 4.0 V, along with long and flat plateaus below 3.0 V and above 4.0 V. Doping effectively smooths the charge-discharge curves within the 3.0\u0026ndash;4.0 V range, though each dopant has a distinct effect on the low-voltage (\u0026lt;\u0026thinsp;3.0 V) and high-voltage (\u0026gt;\u0026thinsp;4.0 V) regions. Sc doping preserves both the low- and high-voltage plateaus, although their lengths are shortened. This behavior is consistent with the dQ/dV curves (Figure S6a, b), where the Sc doped sample exhibits oxidation/reduction peaks near 2.7 V and 4.1 V, similar to the undoped material but with reduced intensity. In contrast, the introduction of Ti and B significantly modifies the voltage profile. The low-voltage plateau becomes more pronounced during both charge and discharge, while the high-voltage plateau transforms into a sloping region during charge and nearly disappears during discharge. These changes are also evident in the corresponding dQ/dV curves of the undoped and Sc\u0026ndash;Ti\u0026ndash;B doped materials (Fig.\u0026nbsp;\u003cspan refid=\"Fig2\" class=\"InternalRef\"\u003e2\u003c/span\u003eb). The above observations highlight the distinct roles of each dopant in influencing redox activity, which could contribute to differences in cycling stability among the various cathode materials.\u003c/p\u003e \u003cp\u003eAll samples were tested at 0.1C, 1C, and 3C (1C\u0026thinsp;=\u0026thinsp;200 mA g\u003csup\u003e\u0026minus;\u0026thinsp;1\u003c/sup\u003e) to evaluate the influence of dopants on cycling performance across different charge-discharge rates. As shown in Fig.\u0026nbsp;\u003cspan refid=\"Fig2\" class=\"InternalRef\"\u003e2\u003c/span\u003ec, the doped materials exhibit significantly better cycling stability at 0.1C compared to the rapid capacity fading observed in the undoped sample, particularly the Sc\u0026ndash;Ti\u0026ndash;B doped material. The corresponding charge-discharge curves for all materials are presented in Figure S7. Upon cycling, the low-voltage and high-voltage plateaus vanish quickly for most samples, except for the O3\u0026ndash;ScTiB doped material, which retains the low-voltage plateau even after 100 cycles. This suggests a highly reversible electrochemical process enabled by the Sc, Ti, and B co-doping. When cycled at 1C, the Sc\u0026ndash;Ti\u0026ndash;B co-doped material delivers a high reversible capacity of approximately 147 mAh g⁻\u0026sup1; and maintains 65% capacity retention after 300 cycles, outperforming both the O3\u0026ndash;ScTi sample (60.5%) and the O3\u0026ndash;Sc\u0026ndash;doped sample (23.2%) (see Fig.\u0026nbsp;\u003cspan refid=\"Fig2\" class=\"InternalRef\"\u003e2\u003c/span\u003ed). It is evident that the Sc\u0026ndash;Ti\u0026ndash;B co-doped sample would exhibit superior capacity retention with further cycling. Unexpectedly, the Sc\u0026ndash;Ti co-doped material demonstrates the best cycling stability at a higher rate of 3C, achieving a capacity of about 106 mAh g⁻\u0026sup1; with 80.1% retention after 500 cycles (see Fig.\u0026nbsp;\u003cspan refid=\"Fig2\" class=\"InternalRef\"\u003e2\u003c/span\u003ee and S8). These results indicate that while Sc\u0026ndash;Ti\u0026ndash;B co-doping significantly enhances cycling stability, its relative advantage becomes less pronounced under higher current densities. As illustrated in Fig.\u0026nbsp;\u003cspan refid=\"Fig2\" class=\"InternalRef\"\u003e2\u003c/span\u003ef, the discharge capacities at all cycling currents increase progressively from Sc doping alone to Sc\u0026ndash;Ti\u0026ndash;B co-doping, with the most prominent improvement observed at 1C. Moreover, B incorporation distinctly improves capacity retention, except under the 3C condition. Furthermore, the rate performance (Figure S9 and S10) reveals that Sc\u0026ndash;Ti\u0026ndash;B doping enhances the rate capability at lower charge-discharge rates compared to Sc\u0026ndash;Ti doping alone. These results indicate that B doping is beneficial for cycling stability at low-current densities but may limit performance at higher-current densities.\u003c/p\u003e \u003cp\u003e \u003c/p\u003e\n\u003ch3\u003eRedox activity analysis of the first cycle\u003c/h3\u003e\n\u003cp\u003eX\u0026ndash;ray absorption spectroscopy (XAS) at both the Ni K-edge and L-edge was conducted to investigate the redox behavior and local structural changes of the doped materials at different states of charge: open-circuit voltage (OCV), charged to 4.0 V, charged to 4.4 V, and discharged to 2.5 V (see Fig.\u0026nbsp;\u003cspan refid=\"Fig3\" class=\"InternalRef\"\u003e3\u003c/span\u003e and S11\u0026ndash;S13). For detailed results on the undoped material, please refer to our recent report\u003csup\u003e\u003cspan citationid=\"CR7\" class=\"CitationRef\"\u003e7\u003c/span\u003e\u003c/sup\u003e. As shown in Fig.\u0026nbsp;\u003cspan refid=\"Fig3\" class=\"InternalRef\"\u003e3\u003c/span\u003ea\u0026ndash;c and S11, a clear energy shift toward higher values is observed at both the Ni K-edge and L-edge during charging, confirming the oxidation of Ni. At the fully charged state (4.4 V), the O3\u0026ndash;ScTi doping sample exhibits the highest edge energy (see inset in Fig.\u0026nbsp;\u003cspan refid=\"Fig3\" class=\"InternalRef\"\u003e3\u003c/span\u003ea), followed by the O3\u0026ndash;ScTiB doping sample, with the undoped and Sc doped samples showing lower edge energies. These trends suggest that Ti doping enhances Ni redox activity, while B doping slightly suppresses it. Upon discharge to 2.5 V, all doped samples show good reversibility, as evidenced by the near\u0026ndash;complete overlap of the spectra with those at the OCV state. In contrast, the extended X\u0026ndash;ray absorption fine structure (EXAFS) region at the Ni K-edge reveals a more complex behavior, indicating changes in the local coordination environment of Ni during cycling. As shown in Fig.\u0026nbsp;\u003cspan refid=\"Fig3\" class=\"InternalRef\"\u003e3\u003c/span\u003ec, d, all samples exhibit two primary peaks: the first at lower radial distance corresponding to the Ni\u0026ndash;O shell, and the second at higher radial distance representing the Ni\u0026ndash;Me (B, Sc, Ni, and Mn) shell\u003csup\u003e\u003cspan citationid=\"CR7\" class=\"CitationRef\"\u003e7\u003c/span\u003e\u003c/sup\u003e. During charging, the Ni\u0026ndash;O bond length gradually decreases for both the Sc\u0026ndash; and Sc\u0026ndash;Ti\u0026ndash;B doped samples. Notably, the Sc doped sample exhibits a more pronounced contraction, indicating greater distortion in the Ni\u0026ndash;O bonds compared to the Sc\u0026ndash;Ti\u0026ndash;B doped sample. Although the changes in peak intensities resemble those observed in the undoped sample, the overall amplitude variations are smaller in the doped materials, especially in the Sc\u0026ndash;Ti\u0026ndash;B sample, suggesting improved structural stability. At the discharge state, the Sc\u0026ndash;Ti\u0026ndash;B doped sample demonstrates superior reversibility, with its EXAFS spectra at OCV and 2.5 V being nearly identical. Figure\u0026nbsp;\u003cspan refid=\"Fig3\" class=\"InternalRef\"\u003e3\u003c/span\u003eb presents the EXAFS data for all samples at 4.4 V. The Sc\u0026ndash;Ti doped sample exhibits the shortest Ni\u0026ndash;O bond length among the materials investigated. At the OCV state, however, its Ni\u0026ndash;O radial distance is nearly identical to that of the Sc\u0026ndash;Ti\u0026ndash;B doped sample, suggesting that B incorporation suppresses the variation in the local Ni environment. Such mitigation of structural distortion is likely to promote improved lattice stability and, consequently, enhanced electrochemical performance.\u003c/p\u003e \u003cp\u003eRedox activity of the lattice oxygen (\u0026ldquo;oxygen redox\u0026rdquo;) is a complex and highly discussed process, which can be observed particularly in Ni\u0026ndash;containing layered oxides operating under high-voltage conditions\u003csup\u003e\u003cspan citationid=\"CR31\" class=\"CitationRef\"\u003e31\u003c/span\u003e,\u003cspan citationid=\"CR32\" class=\"CitationRef\"\u003e32\u003c/span\u003e\u003c/sup\u003e. In this study, we employed resonant inelastic X\u0026ndash;ray scattering (RIXS) to investigate the influence of various dopants on the oxygen redox behavior. While the origin of oxygen redox signals detected by RIXS remains a topic of ongoing debate, our focus here is not on the mechanism itself but rather on the quantitative variations associated with different dopants\u003csup\u003e\u003cspan citationid=\"CR12\" class=\"CitationRef\"\u003e12\u003c/span\u003e,\u003cspan citationid=\"CR33\" class=\"CitationRef\"\u003e33\u003c/span\u003e\u003c/sup\u003e. As shown in Fig.\u0026nbsp;\u003cspan refid=\"Fig3\" class=\"InternalRef\"\u003e3\u003c/span\u003ee, distinct differences are observed at the ~\u0026thinsp;8 eV energy loss, which is widely considered to represent the spectral fingerprint of oxygen redox\u003csup\u003e\u003cspan citationid=\"CR12\" class=\"CitationRef\"\u003e12\u003c/span\u003e,\u003cspan citationid=\"CR31\" class=\"CitationRef\"\u003e31\u003c/span\u003e,\u003cspan additionalcitationids=\"CR34\" citationid=\"CR33\" class=\"CitationRef\"\u003e33\u003c/span\u003e\u0026ndash;\u003cspan citationid=\"CR35\" class=\"CitationRef\"\u003e35\u003c/span\u003e\u003c/sup\u003e. Compared to the O3\u0026ndash;undoped sample at the OCV state, all doped samples exhibit similar peaks, albeit with varying intensities. These results confirm that oxygen redox activity is present across all samples to varying extents.\u003c/p\u003e \u003cp\u003eTo estimate the relative contribution of oxygen redox, the A/B intensity ratios were used and are presented in Fig.\u0026nbsp;\u003cspan refid=\"Fig3\" class=\"InternalRef\"\u003e3\u003c/span\u003eh (top). Clearly, all doped materials exhibit lower oxygen redox intensity compared to the undoped sample, likely due to reduced Ni content in the doped compositions. Among the doped materials, the Sc\u0026ndash;Ti co-doped sample displays the weakest oxygen redox signals, which aligns with its largest Ni\u0026ndash;O bond lengths in the EXAFS (see Fig.\u0026nbsp;\u003cspan refid=\"Fig3\" class=\"InternalRef\"\u003e3\u003c/span\u003eb). Interestingly, the 2% B doping (O3\u0026ndash;ScTiB doping) markedly enhances oxygen redox activities compared to the Sc\u0026ndash;Ti doped counterpart. However, boron is a metalloid, differing fundamentally from metallic dopants such as Li, Mg, Ni, or Zn, which are traditionally associated with triggering oxygen redox by forming the X\u0026ndash;O\u0026ndash;Na configurations (where X refers to the aforementioned dopants)\u003csup\u003e\u003cspan additionalcitationids=\"CR37 CR38\" citationid=\"CR36\" class=\"CitationRef\"\u003e36\u003c/span\u003e\u0026ndash;\u003cspan citationid=\"CR39\" class=\"CitationRef\"\u003e39\u003c/span\u003e\u003c/sup\u003e. This distinction implies that boron may play a unique and different role in modulating oxygen redox behavior.\u003c/p\u003e \u003cp\u003eTo further investigate the oxygen redox behavior among these samples during the initial cycle, O K-edge XAS measurements were conducted at various electrochemical states. In the O K-edge spectra (probing the 1s \u0026rarr; 2p transition), delocalized states arising from the hybridization between oxygen 2p and Ni/Mn 3d (e\u003csub\u003eg\u003c/sub\u003e) orbitals appear within the 527\u0026ndash;534 eV range\u003csup\u003e\u003cspan citationid=\"CR34\" class=\"CitationRef\"\u003e34\u003c/span\u003e,\u003cspan additionalcitationids=\"CR41\" citationid=\"CR40\" class=\"CitationRef\"\u003e40\u003c/span\u003e\u0026ndash;\u003cspan citationid=\"CR42\" class=\"CitationRef\"\u003e42\u003c/span\u003e\u003c/sup\u003e. This region is divided into three distinct zones: Region C (527.5\u0026ndash;529 eV), Region D (\u0026lt;\u0026thinsp;530.5 eV), and Region E (\u0026lt;\u0026thinsp;532.7 eV). In Region C, shoulder features emerge at both the 4.0 V and 4.4 V charge states, corresponding to the formation of oxygen holes previously reported in Na-layered oxides during desodiation\u003csup\u003e\u003cspan citationid=\"CR12\" class=\"CitationRef\"\u003e12\u003c/span\u003e,\u003cspan citationid=\"CR34\" class=\"CitationRef\"\u003e34\u003c/span\u003e,\u003cspan citationid=\"CR37\" class=\"CitationRef\"\u003e37\u003c/span\u003e,\u003cspan citationid=\"CR43\" class=\"CitationRef\"\u003e43\u003c/span\u003e\u003c/sup\u003e. The peak in Region D originates from O 2p states hybridized with unoccupied TM 3d orbitals of spin-down t₂\u003csub\u003eg\u003c/sub\u003e and spin-up e\u003csub\u003eg\u003c/sub\u003e character, reflecting contributions from both exchange-split subbands. The higher energy feature in Region E corresponds to hybridization with spin-down e\u003csub\u003eg\u003c/sub\u003e states\u003csup\u003e\u003cspan citationid=\"CR42\" class=\"CitationRef\"\u003e42\u003c/span\u003e\u003c/sup\u003e. As shown in Figs.\u0026nbsp;\u003cspan refid=\"Fig3\" class=\"InternalRef\"\u003e3\u003c/span\u003ef, \u003cspan refid=\"Fig3\" class=\"InternalRef\"\u003e3\u003c/span\u003eg, and S14, S15, all samples exhibit a similar spectral evolution upon charging: the spectra shift toward lower energies, and a new peak appears at 528.6 eV at the 4.0 V charge state. At full charge (4.4 V), the intensity of the 529.4 eV peak in Region D diminishes, indicating that oxygen actively participates in the electrochemical process starting from 4.0 V, with the formation of oxygen holes. Upon discharge, the spectra largely revert to their initial state, suggesting that the oxygen holes are only present at high voltages (\u0026gt;\u0026thinsp;4.0 V). To quantify the extent of oxygen redox among the samples, the change in the C/D intensity ratio between the 4.0 V and 4.4 V charge states was analyzed and is presented in Fig.\u0026nbsp;\u003cspan refid=\"Fig3\" class=\"InternalRef\"\u003e3\u003c/span\u003eh (bottom). Notably, the Δ(C/D) trend closely mirrors the A/B ratio obtained from RIXS measurements, where the intensity decreases from the undoped sample to the Sc and Ti co-doped case, and then increases with B incorporation. Overall, the RIXS and O K-edge XAS data reveal that Sc and Ti co-doping significantly suppresses oxygen redox, whereas B doping reverses this effect and markedly enhances oxygen redox activity.\u003c/p\u003e \u003cp\u003e \u003c/p\u003e\n\u003ch3\u003eStructural degradation during the first cycle at a low-current density (0.1C)\u003c/h3\u003e\n\u003cp\u003eSynchrotron\u0026ndash;based \u003cem\u003eoperando\u003c/em\u003e XRD was employed to study how various dopants influence the structural evolution of O3 type cathodes during the first cycle. For the undoped O3\u0026ndash;NaNi\u003csub\u003e0.5\u003c/sub\u003eMn\u003csub\u003e0.5\u003c/sub\u003eO\u003csub\u003e2\u003c/sub\u003e sample, our previous study revealed complex but reversible structural transformations, following the sequence: O3 (rhombohedral) \u0026rarr; O'3 (monoclinic) \u0026rarr; P'3 (monoclinic) and P3 (hexagonal) \u0026rarr; O'3 (monoclinic) \u0026rarr; O3 (rhombohedral), which results in more than 20% lattice contraction at the fully charged state\u003csup\u003e\u003cspan citationid=\"CR7\" class=\"CitationRef\"\u003e7\u003c/span\u003e\u003c/sup\u003e. With doping, all samples exhibit similar trends for changes in the XRD reflections below 4.0 V (see Fig.\u0026nbsp;\u003cspan refid=\"Fig4\" class=\"InternalRef\"\u003e4\u003c/span\u003ea, b, and S16\u0026ndash;S18). Specifically, the 003 and 006 reflections shift to lower angles and the reflections of 101, 012 and 104 are found to shift toward higher angles, indicating an expansion along the \u003cem\u003ec\u003c/em\u003e axis and reduction in the \u003cem\u003eab\u003c/em\u003e plane\u003csup\u003e\u003cspan citationid=\"CR7\" class=\"CitationRef\"\u003e7\u003c/span\u003e,\u003cspan citationid=\"CR44\" class=\"CitationRef\"\u003e44\u003c/span\u003e\u003c/sup\u003e. When being charged to 3.0 V, a clear phase transition from O\u0026ndash;type to P\u0026ndash;type structure is observed, as evidenced by the appearance of 015\u003csub\u003eP3\u003c/sub\u003e at about 5.8\u0026deg; (λ\u0026thinsp;=\u0026thinsp;0.20733 \u0026Aring;) and 10.1\u0026deg; (λ\u0026thinsp;=\u0026thinsp;0.35424 \u0026Aring;). Subsequently, all doped materials remain in a single P3 phase up to 4.0 V. Notably, the Sc\u0026ndash;Ti and Sc\u0026ndash;Ti\u0026ndash;B co-doped materials demonstrate a delayed transition to the P3 phase compared to the sample doped with Sc alone, suggesting enhanced structural stability or kinetic hindrance introduced by co-doping.\u003c/p\u003e \u003cp\u003eAt voltages above 4.0 V, the doped materials exhibit distinct phase transitions. For the Sc doped sample, a similar transformation to that of the undoped counterpart is observed. The main 003\u003csub\u003eP3\u003c/sub\u003e reflection shifts slightly to higher angles, indicating the formation of a monoclinic O'3 phase\u003csup\u003e\u003cspan citationid=\"CR8\" class=\"CitationRef\"\u003e8\u003c/span\u003e\u003c/sup\u003e. This is corroborated by the disappearance of the 012\u003csub\u003eP3\u003c/sub\u003e and 015\u003csub\u003eP3\u003c/sub\u003e reflections. Concurrently, a reflection at 4.8\u0026deg; becomes prominent, indexed as the 100\u003csub\u003eO'3\u003c/sub\u003e reflection and highlighted by the magenta dashed box in Fig.\u0026nbsp;\u003cspan refid=\"Fig4\" class=\"InternalRef\"\u003e4\u003c/span\u003ea, further supporting the phase transition. These features closely mirror those seen in the undoped sample. Upon further charging from 4.2 V to 4.4 V, most reflections vanish except for the persistent 100\u003csub\u003eO'3\u003c/sub\u003e reflection, suggesting that the O'3 phase remains dominant in the high-voltage region. Meanwhile, a new reflection appears at 2.4\u0026deg;, and gradually shifts to lower angles upon discharge. This behavior indicates the formation of a new phase, resembling the recently reported OP4 phase in 10% Mg doped P2\u0026ndash;Na\u003csub\u003e0.67\u003c/sub\u003eNi\u003csub\u003e0.33\u003c/sub\u003eMn\u003csub\u003e0.67\u003c/sub\u003eO\u003csub\u003e2\u003c/sub\u003e\u003csup\u003e31,45,46\u003c/sup\u003e. In this study, the newly formed phase is identified as an OP2\u0026ndash;type structure\u003csup\u003e\u003cspan citationid=\"CR46\" class=\"CitationRef\"\u003e46\u003c/span\u003e\u003c/sup\u003e. In the Sc and Ti co-doped sample, the 003\u003csub\u003eP3\u003c/sub\u003e reflection significantly weakens at 4.0 V and becomes faint at 4.2 V (see Figures S17 and S19). As highlighted in the 003 reflection in Figure S16b, the diffraction pattern still shows a dominant O'3 phase, along with a broad reflection at 2.0\u0026deg;. This new reflection is similar to the reported P1 phase in NaNi\u003csub\u003e1/3\u003c/sub\u003eMn\u003csub\u003e1/3\u003c/sub\u003eCo\u003csub\u003e1/3\u003c/sub\u003eO\u003csub\u003e2\u003c/sub\u003e at the desodiated state\u003csup\u003e\u003cspan citationid=\"CR47\" class=\"CitationRef\"\u003e47\u003c/span\u003e\u003c/sup\u003e. These results indicate that both Sc\u0026ndash;only doping and Sc\u0026ndash;Ti co-doping promote the formation of a P\u0026ndash;O intergrowth phase, with the O phase remaining predominant. For the Sc\u0026ndash;Ti\u0026ndash;B co-doped material, a more continuous structural evolution is observed in the high-voltage region. Upon charging to 4.0 V, the main 003 reflection progressively shifts to higher angles, reaching its maximum 2θ value at the fully charged state. Like the Sc\u0026ndash;only doped material, this phase is also an OP2 phase. During the discharge process, the Sc doped sample exhibits only partial reversibility, as indicated by the persistence of mixed phases from 3.7 V down to the fully discharged state. In contrast, the Sc\u0026ndash;Ti\u0026ndash;B co-doped material shows a highly reversible structural evolution, reflecting improved cycling stability under low-current cycling. These findings suggest that Sc/Ti co-doping effectively suppresses the formation of intermediate phases, while B doping promotes a more continuous and reversible structural transformation in the high-voltage region.\u003c/p\u003e \u003cp\u003e \u003c/p\u003e \u003cp\u003eTo further elucidate the effects of different dopants on structural evolution, the \u003cem\u003ec\u003c/em\u003e lattice parameter evolution and phase transition behaviors are illustrated in Fig.\u0026nbsp;\u003cspan refid=\"Fig4\" class=\"InternalRef\"\u003e4\u003c/span\u003ec. All samples exhibit lattice expansion up to approximately 4.0 V, followed by contraction at the fully charged state. Among them, the Sc\u0026ndash;Ti\u0026ndash;B co-doped sample shows the largest expansion (\u0026thinsp;~\u0026thinsp;+\u0026thinsp;6.37%), even exceeding that of the undoped material. At full charge, all materials exhibit mixed phases except for the Sc\u0026ndash;Ti\u0026ndash;B doped material (see Figure S20 and S21). The undoped material primarily demonstrates the O3 phase with a significant number of vacancies, along with traces of the O\u0026prime;3 phase. For the Sc\u0026ndash; and Sc\u0026ndash;Ti doped cases, both are dominated by the O\u0026prime;3 phase, with a small amount of the OP2 phase present in the Sc doped sample and the P1 phase in the Sc\u0026ndash;Ti doped sample. In contrast, the Sc\u0026ndash;Ti\u0026ndash;B co-doped sample displays a single OP2 phase, showing an intermediate level of \u003cem\u003ec\u003c/em\u003e lattice parameter changes with the lowest ratio of O phase stacking. This aligns with the observation that a dominant P phase is favorable for better electrochemical performance in the O3\u0026ndash;type framework\u003csup\u003e\u003cspan citationid=\"CR20\" class=\"CitationRef\"\u003e20\u003c/span\u003e\u003c/sup\u003e. The phase transition diagram further illustrates this behavior, showing that 3.0 V and 4.0 V act as key boundaries between P\u0026ndash;type and O\u0026ndash;type phases in O3\u0026ndash;type materials. Dopants mainly affect the structure in the high-desodiation state, ultimately influencing structural reversibility during cycling. In summary, Sc and Ti co-doping effectively suppresses the formation of the intermediate O\u0026prime;3 phase, while B doping delays the O3\u0026ndash;P3 transition and improves structural reversibility through the stabilization of an OP2 phase at high voltage.\u003c/p\u003e \u003cp\u003eHAADF-STEM was employed to investigate the microstructural evolution at different states. As shown in Fig.\u0026nbsp;\u003cspan refid=\"Fig5\" class=\"InternalRef\"\u003e5\u003c/span\u003ea, d, O3\u0026ndash;undoped and O3\u0026ndash;ScTiB doping samples maintain a well-defined layered structure in the bulk, as evidenced by distinct lattice fringes and sharp, regular diffraction spots in the corresponding fast Fourier transform (FFT) patterns. However, in the near-surface region, the O3\u0026ndash;undoped sample develops a pronounced cation-mixed layer that is ~\u0026thinsp;7 nm thick, whereas doping effectively suppresses the formation of such a disordered region. At the fully charged state, the O3\u0026ndash;undoped sample exhibits extensive cracking within the lattice that propagate throughout the crystalline domains, accompanied by significant lattice strain, particularly in the areas near the cracks. In contrast, the Sc\u0026ndash;Ti\u0026ndash;B co-doped sample preserves a well-ordered layered structure with minimal dislocations and markedly reduced strain in the fully charged state (Fig.\u0026nbsp;\u003cspan refid=\"Fig5\" class=\"InternalRef\"\u003e5\u003c/span\u003eb, c, and \u003cspan refid=\"Fig5\" class=\"InternalRef\"\u003e5\u003c/span\u003ee, f). These observations are consistent with the smaller contraction of the \u003cem\u003ec\u003c/em\u003e lattice parameter in the doped materials. Furthermore, electron energy loss spectroscopy (EELS) at the O K-edge and Ni/Mn L-edges was performed at both the OCV and fully charged states. Both materials exhibit a clear O K-edge signal at the OCV state. However, at 4.4 V, the undoped sample shows a significantly weakened oxygen signal near the surface compared with the Sc\u0026ndash;Ti\u0026ndash;B co-doped sample. This observation highlights the structural instability of the undoped material in the absence of doping (Figs.\u0026nbsp;\u003cspan refid=\"Fig5\" class=\"InternalRef\"\u003e5\u003c/span\u003eg, h, and S22, S23). Taken together, the \u003cem\u003eoperando\u003c/em\u003e XRD, STEM, and EELS results confirm that lattice contraction-induced strain and oxygen loss are the primary factors driving material degradation in high-voltage cathodes under low-current conditions.\u003c/p\u003e \u003cp\u003e \u003c/p\u003e\n\u003ch3\u003eRedox degradation during cycling at low current density (0.1C)\u003c/h3\u003e\n\u003cp\u003eXAS at the Ni L-edge, Ni K-edge, and O K-edge was performed for both undoped and doped materials to study the redox degradation at 0.1C during cycling. As shown in Figs.\u0026nbsp;\u003cspan refid=\"Fig6\" class=\"InternalRef\"\u003e6\u003c/span\u003ea and S24, S25, all samples exhibit noticeable energy shifts in the Ni K-edge XANES spectra over multiple cycles. For the undoped material, the absorption edge gradually shifts to higher energy, indicating progressive oxidation. In contrast, the doped samples show shifts toward lower energies, and the magnitude of these shifts diminishes with continued cycling, especially for the Sc\u0026ndash;Ti\u0026ndash;B doped sample. In the EXAFS spectra, a clear shift in the Ni\u0026ndash;O shell is observed for the undoped and Sc\u0026ndash;Ti doped materials, suggesting an unstable coordination environment. In comparison, the O3\u0026ndash;ScTiB doped sample maintains a consistent Ni\u0026ndash;O radial distance throughout cycling, indicating a more stable local structure. As illustrated in Fig.\u0026nbsp;\u003cspan refid=\"Fig3\" class=\"InternalRef\"\u003e3\u003c/span\u003ef, the shoulder feature associated with O hole contributions diminishes during cycling for the undoped sample. However, this change is minimal in the Sc\u0026ndash;Ti\u0026ndash;B doped material, with nearly no variation between the 1st and 15th cycles. To better compare the effects of doping on O redox behavior, summaries are provided in Fig.\u0026nbsp;\u003cspan refid=\"Fig6\" class=\"InternalRef\"\u003e6\u003c/span\u003ee for O3\u0026ndash;undoped, O3\u0026ndash;ScTi doping, and O3\u0026ndash;ScTiB doping, respectively. Notably, in the undoped sample, Ni oxidation increases while O hole contributions decrease over 30 cycles. This suggests a growing reliance on Ni redox activity and a declining contribution from O redox, consistent with the charge-discharge curves shown in Figure S7. In doped materials, both Ni and O redox activities gradually decrease over cycling. However, the O3\u0026ndash;ScTiB doping sample demonstrates superior stability, particularly in maintaining O redox contributions. Furthermore, it retains a stable Ni\u0026ndash;O coordination environment, with negligible changes in radial distance during cycling. These findings align with the enhanced cycling stability observed at low current rates.\u003c/p\u003e \u003cp\u003e \u003c/p\u003e \u003cp\u003eTo further clarify the influence of dopants on oxygen redox, DFT calculations were carried out for O3\u0026ndash;undoped, O3\u0026ndash;Sc-doped, O3\u0026ndash;ScTi-doped, and O3\u0026ndash;ScTiB-doped samples. As shown in Figs.\u0026nbsp;\u003cspan refid=\"Fig6\" class=\"InternalRef\"\u003e6\u003c/span\u003eg and S26, the computational results reveal the same trend as observed experimentally: oxygen redox activities decrease from the undoped case to the Sc\u0026ndash;Ti co-doped case, and then increases with B incorporation. This behavior arises from the stronger orbital overlap between O 2p and Ni 3d t\u003csub\u003e2g\u003c/sub\u003e bonding orbitals induced by Sc and Ti, which stabilizes electrons in the bonding states and suppresses oxygen redox activities. B doping further enhances the O 2p\u0026ndash;Ni 3d t\u003csub\u003e2g\u003c/sub\u003e overlap to stabilize oxygen, but at the same time modulates orbital energy levels and electron distribution by increasing the occupation of Ni 3d e\u003csub\u003eg\u003c/sub\u003e antibonding orbitals. This adjustment leads to an overall enhancement of oxygen redox activities. Collectively, the synergistic contributions of Sc, Ti, and B produce an optimized electronic configuration of oxygen that suppresses side reactions such as oxygen loss to ensure stability, while enabling highly reversible redox with maximal electron transfer during electrochemical cycling. Such electronic-level regulation markedly reinforces oxygen redox activity in Ni\u0026ndash;Mn layered oxides, in excellent agreement with experimental results.\u003c/p\u003e \u003cp\u003eBuilding on these investigations at low-current densities, a comparison of the electrochemical performance, structural evolution, and oxygen stability is summarized in the spider chart shown in Fig.\u0026nbsp;\u003cspan refid=\"Fig6\" class=\"InternalRef\"\u003e6\u003c/span\u003eh. For the undoped material, structural instability and oxygen loss are the primary contributors to degradation. In contrast, doped materials, whether with single or multiple dopants, exhibit similar structural effects, effectively mitigating lattice strain compared to the undoped O3\u0026ndash;NaNi\u003csub\u003e0.5\u003c/sub\u003eMn\u003csub\u003e0.5\u003c/sub\u003eO\u003csub\u003e2\u003c/sub\u003e. In these cases, oxygen stability becomes the dominant factor governing long-term cycling performance at the low-current densities. It is worth emphasizing, however, that oxygen loss is inevitable in high-voltage cathodes; the role of doping is to slow down this process rather than eliminate it.\u003c/p\u003e \u003cdiv id=\"Sec8\" class=\"Section2\"\u003e \u003ch2\u003eStructure degradation at high current densities\u003c/h2\u003e \u003cp\u003eAt high cycling currents (\u0026gt;\u0026thinsp;1C), the capacity contribution from oxygen redox becomes negligible (see rate performance in Figure S9). This is attributed to the intrinsically sluggish kinetics of oxygen redox during cycling (see GITT data in Figure S27), suggesting that oxygen stability plays a much less critical role at high currents compared to low-current cycling. To elucidate the dominant factor under fast de/intercalation process, synchrotron\u0026ndash;based \u003cem\u003eoperando\u003c/em\u003e XRD was employed to track the structural evolution. As shown in Fig.\u0026nbsp;\u003cspan refid=\"Fig7\" class=\"InternalRef\"\u003e7\u003c/span\u003ea, b and S28, the phase transition pathway becomes markedly simplified at high-current densities. In the Sc\u0026ndash;Ti doped sample, a clear O3\u0026ndash;P3 phase transition occurs during the first charge. During subsequent cycles at both 1C and 3C, the material stabilizes in a single P3 phase, showing only a minimal shift in the 2θ angle, corresponding to a small change of ~\u0026thinsp;+\u0026thinsp;1.5% in the \u003cem\u003ec\u003c/em\u003e lattice parameter. In contrast, the B doped sample exhibits a more complex transition pathway. At 1C, it undergoes reversible O3 \u0026rarr; P3 \u0026rarr; OP2 transitions, similar to those observed at 0.1C. When cycled at 3C, however, a distinct X phase emerges at the fully charged state, following the sequence O3 \u0026rarr; P3 \u0026rarr; X. This process becomes irreversible during discharge, as the material only transitions back from X to P3. Notably, the Sc\u0026ndash;Ti\u0026ndash;B co-doped sample shows much larger variations in the \u003cem\u003ec\u003c/em\u003e lattice parameter, with contractions of -7.4% at 1C and \u0026minus;\u0026thinsp;1.7% at 3C, compared to the Sc\u0026ndash;Ti doped sample. This greater structural fluctuation is likely a key factor underlying its inferior capacity retention under high-current cycling conditions.\u003c/p\u003e \u003cp\u003e \u003c/p\u003e \u003c/div\u003e"},{"header":"Discussion","content":"\u003cp\u003eAt low current densities, the oxygen redox reaction can fully participate in reversible charge compensation, which increases its contribution to the overall capacity. For the Sc\u0026ndash;Ti\u0026ndash;B co-doped sample, the additional boron induces a larger change in the \u003cem\u003ec\u003c/em\u003e lattice parameter. This effect is attributed to the stronger B\u0026ndash;O bond, which effectively suppresses oxygen loss.\u003csup\u003e\u003cspan citationid=\"CR26\" class=\"CitationRef\"\u003e26\u003c/span\u003e\u003c/sup\u003e This stabilization enhances the reversibility of both structural evolution and the oxygen redox process. As a result, the B doped material delivers higher capacity and improved cycling stability compared to its Sc\u0026ndash;Ti doped counterpart. These findings suggest that at low cycling currents, long-term cycling performance is primarily controlled by the stability of oxygen redox, and that suppressing oxygen loss is more decisive than mitigating high voltage phase transitions. However, as the cycling current increases, the behavior changes: regardless of dopant type, the material primarily undergoes a simple P3 phase transition during cycling (see Fig.\u0026nbsp;\u003cspan refid=\"Fig7\" class=\"InternalRef\"\u003e7\u003c/span\u003ec). Under these rapid charge-discharge conditions, the contribution of oxygen redox to capacity becomes negligible, while structural changes emerge as the dominant factor governing cycling performance. Consequently, the Sc\u0026ndash;Ti doped sample exhibits superior cycling stability at high-current densities. Furthermore, a spider chart summarizing the respective roles of TM redox, oxygen redox, and changes in the \u003cem\u003ec\u003c/em\u003e lattice parameter on the capacity and cycling stability of high-voltage layered cathodes corroborates this trend (see Fig.\u0026nbsp;\u003cspan refid=\"Fig7\" class=\"InternalRef\"\u003e7\u003c/span\u003ed). It highlights that oxygen redox is more critical for determining capacity and stability at low currents, whereas structural evolution exerts a stronger influence on cycling performance under high-current cycling.\u003c/p\u003e \u003cp\u003eIn summary, we have synthesized a series of Sc\u0026ndash;, Sc\u0026ndash;Ti\u0026ndash;, and Sc\u0026ndash;Ti\u0026ndash;B co-doped O3\u0026ndash;type Ni-Mn based layered oxides via a solid-state reaction method. Notably, the Sc\u0026ndash;Ti\u0026ndash;B co-doping significantly enhances the Na storage performance at low cycling currents, achieving a capacity of 147 mAh g\u003csup\u003e\u0026minus;\u0026thinsp;1\u003c/sup\u003e at 1C with a capacity retention of 83.5% after 100 cycles in the voltage range of 2.5 to 4.4 V. Using synchrotron-based techniques such as XRD, XAS, RIXS, together with electron microscopy and DFT calculations, it is revealed that the additional B doping effectively suppresses oxygen loss, thereby improving the reversibility of both structural evolution and the oxygen redox process at low-current densities. However, at high cycling current (3C), the Sc\u0026ndash;Ti doped material (without B) demonstrates superior capacity retention. This indicates that structural integrity plays a more critical role in capacity retention than oxygen redox stability under high-current cycling. These findings uncover a rate-dependent degradation mechanism in layered oxides under high-voltage operation, in which oxygen stability is the key to achieving high capacity and cycling stability at low currents, whereas structural stability becomes the primary factor shaping degradation at high currents. This current dependent structural and redox behaviour is particularly relevant for practical sodium ion batteries, because real drive cycles impose widely varying current densities. Cells operated under such conditions may age through mechanisms that differ from those inferred from constant current laboratory tests, and the framework proposed here provides guidance for developing high voltage layered cathodes that remain stable under realistic cycling profiles.\u003c/p\u003e"},{"header":"Methods","content":"\u003cdiv id=\"Sec11\" class=\"Section2\"\u003e \u003ch2\u003eSynthesis methods\u003c/h2\u003e \u003cp\u003eThe synthesis procedure for all O3\u0026ndash;type cathodes followed the method reported previously\u003csup\u003e\u003cspan citationid=\"CR7\" class=\"CitationRef\"\u003e7\u003c/span\u003e\u003c/sup\u003e. The starting materials included Na\u003csub\u003e2\u003c/sub\u003eO\u003csub\u003e2\u003c/sub\u003e (\u0026gt;\u0026thinsp;97%, Sigma-Aldrich), Mn\u003csub\u003e2\u003c/sub\u003eO\u003csub\u003e3\u003c/sub\u003e (\u0026gt;\u0026thinsp;99%, Sigma-Aldrich), NiO (\u0026gt;\u0026thinsp;99%, Sigma-Aldrich), Sc\u003csub\u003e2\u003c/sub\u003eO\u003csub\u003e3\u003c/sub\u003e (\u0026gt;\u0026thinsp;99.9%, Sigma-Aldrich), TiO\u003csub\u003e2\u003c/sub\u003e (\u0026gt;\u0026thinsp;99.9%, Sigma-Aldrich), and B\u003csub\u003e2\u003c/sub\u003eO\u003csub\u003e3\u003c/sub\u003e (\u0026gt;\u0026thinsp;99.98%, Sigma-Aldrich). An excess of 5 mol% Na\u003csub\u003e2\u003c/sub\u003eO\u003csub\u003e2\u003c/sub\u003e was added to compensate for Na loss during the high temperature reaction, and 0.015 mol of product was synthesized for each material. All raw materials were mixed for 5 h using a planetary ball mill with a reversible procedure, then the mixture was fabricated as a pellet by pressing under 15 MPa pressure. After that, the pellet was calcined at 950\u0026deg;C for 15 h under O\u003csub\u003e2\u003c/sub\u003e atmosphere to obtain the final products. Before cooling down to 200\u0026deg;C, the pellet was transferred to an Ar-filled Glovebox for further processing and measurement.\u003c/p\u003e \u003c/div\u003e \u003cdiv id=\"Sec12\" class=\"Section2\"\u003e \u003ch2\u003eElectrochemical characterization\u003c/h2\u003e \u003cp\u003eThe process of fabricating working electrodes was the same as previously reported\u003csup\u003e\u003cspan citationid=\"CR7\" class=\"CitationRef\"\u003e7\u003c/span\u003e\u003c/sup\u003e. The electrode composition consisted of active material, Super P (SP, MTI Corp.), and polyvinylidene difluoride (PVDF, MTI Corp.) at a weight ratio of 80:10:10. The PVDF binder was dissolved in N-methyl-2-pyrrolidone (NMP, Sigma Aldrich) at a concentration of 5 wt%. The active material and Super P were first mixed in a shaker ball mill for 10 min. The PVDF solution was then added, and the mixture was further milled for 30 min to obtain a homogeneous slurry. This slurry was cast onto carbon-coated Al foil (MTI Corp.) using a doctor blade with a wet thickness of 250 \u0026micro;m. All fabrication steps were conducted in an Ar-filled glovebox. The coated electrodes were dried on a hot plate, punched into 12 mm discs, and subsequently vacuum-dried at 120\u0026deg;C for 12 h in a B\u0026uuml;chi oven. The electrolyte consisted of 1 M NaClO\u003csub\u003e4\u003c/sub\u003e dissolved in a mixture of propylene carbonate (PC, Sigma Aldrich) and fluoroethylene carbonate (FEC, Sigma Aldrich) with a volume ratio of 98:2. Glass fiber (Whatman) was used as the separator. Coin-type 2025 cells were assembled in an Ar-filled glovebox, with sodium metal (BASF Corp.) serving as the counter electrode. Electrochemical tests were performed in the voltage range of 2.5\u0026ndash;4.4 V (vs. Na\u003csup\u003e+\u003c/sup\u003e/Na) under constant-current conditions at various current rates using a Neware battery test system (CT-4008-5V10mA-164, Neware). For GITT measurements, each step consisted of charging/discharging at 0.1C for 30 min, followed by a rest period of 10 h.\u003c/p\u003e \u003c/div\u003e \u003cdiv id=\"Sec13\" class=\"Section2\"\u003e \u003ch2\u003eStructural characterization\u003c/h2\u003e \u003cp\u003ePowder XRD was conducted using a Bruker D8 Twin-Twin diffractometer equipped with a Cu Kα source (λ\u0026thinsp;=\u0026thinsp;1.5406 \u0026Aring;), operated at 40 kV and 40 mA over the range of 10\u0026ndash;70\u0026deg; (2θ) with a scanning step of 0.01\u0026deg; s\u003csup\u003e\u0026minus;\u0026thinsp;1\u003c/sup\u003e. \u003cem\u003eOperando\u003c/em\u003e XRD measurements for O3\u0026ndash;NaSc\u003csub\u003e0.1\u003c/sub\u003eNi\u003csub\u003e0.4\u003c/sub\u003eMn\u003csub\u003e0.5\u003c/sub\u003eO\u003csub\u003e2\u003c/sub\u003e and O3\u0026ndash;NaSc\u003csub\u003e0.1\u003c/sub\u003eNi\u003csub\u003e0.4\u003c/sub\u003eMn\u003csub\u003e0.4\u003c/sub\u003eTi\u003csub\u003e0.1\u003c/sub\u003eO\u003csub\u003e2\u003c/sub\u003e, along with \u003cem\u003eex-situ\u003c/em\u003e PDF measurements, were performed at beamline P02.1 of PETRA III at DESY (Deutsches Elektronensynchrotron), Hamburg, Germany, using synchrotron radiation with an energy of ~\u0026thinsp;60 keV\u003csup\u003e\u003cspan citationid=\"CR49\" class=\"CitationRef\"\u003e49\u003c/span\u003e,\u003cspan citationid=\"CR50\" class=\"CitationRef\"\u003e50\u003c/span\u003e\u003c/sup\u003e. In-house designed \u003cem\u003eoperando\u003c/em\u003e cells were assembled in an Ar-filled glovebox (H\u003csub\u003e2\u003c/sub\u003eO and O\u003csub\u003e2\u003c/sub\u003e concentrations\u0026thinsp;\u0026lt;\u0026thinsp;5 ppm, as provided by the beamline). Electrochemical measurements were carried out with a Biologic VMP-3 potentiostat controlled by EC-Lab software (v11.34). Prior to cycling, the cells were rested for 20 min, followed by charge-discharge at 0.1 C within the voltage window of 2.5\u0026ndash;4.4 V. \u003cem\u003eOperando\u003c/em\u003e XRD data were collected sequentially from up to five cells operating simultaneously in Debye-Scherrer geometry, with an acquisition time of 1 min per cell. Data collection was automated using a motorized XY stage with a 350 mm travel range in both directions, controlled by a Python3 script to synchronize stage movement and data acquisition. Consequently, each \u003cem\u003eoperando\u003c/em\u003e cell was measured for ~\u0026thinsp;1 min approximately every 6 min. \u003cem\u003eOperando\u003c/em\u003e XRD measurement for O3\u0026ndash;NaSc\u003csub\u003e0.10\u003c/sub\u003eNi\u003csub\u003e0.4\u003c/sub\u003eMn\u003csub\u003e0.4\u003c/sub\u003eTi\u003csub\u003e0.08\u003c/sub\u003eB\u003csub\u003e0.02\u003c/sub\u003eO\u003csub\u003e2\u003c/sub\u003e were carried out at the DanMAX beamline of Max IV, Lund, Sweden. The pouch cells were mounted in a custom-designed holder capable of accommodating up to six cells simultaneously\u003csup\u003e\u003cspan citationid=\"CR51\" class=\"CitationRef\"\u003e51\u003c/span\u003e\u003c/sup\u003e. Data were collected in transmission mode using a Dectris Pilatus3 2M CdTe detector, with an acquisition interval of 2 min per pouch cell. The X-ray wavelength was calibrated to λ\u0026thinsp;=\u0026thinsp;0.35424 \u0026Aring; using a LaB\u003csub\u003e6\u003c/sub\u003e standard. All Rietveld refinements of the diffraction patterns were performed with Fullprof\u003csup\u003e\u003cspan citationid=\"CR52\" class=\"CitationRef\"\u003e52\u003c/span\u003e\u003c/sup\u003e.\u003c/p\u003e \u003c/div\u003e \u003cdiv id=\"Sec14\" class=\"Section2\"\u003e \u003ch2\u003eSpectroscopic characterization\u003c/h2\u003e \u003cp\u003e \u003cem\u003eEx situ\u003c/em\u003e Ni and Mn K-edge XAS measurements were performed at room temperature at the KMC-2 end-station of BESSY II (Berlin, Germany) in both transmission and fluorescence modes\u003csup\u003e\u003cspan citationid=\"CR53\" class=\"CitationRef\"\u003e53\u003c/span\u003e\u003c/sup\u003e. Reference spectra for energy calibration were collected before and after each sample. XAS spectra were processed using \u003cem\u003eAthena\u003c/em\u003e software by subtracting the pre-edge background and normalizing with a spline fit. The k\u003csup\u003e2\u003c/sup\u003e-weighted EXAFS was Fourier transformed over the k-range of 3\u0026ndash;11 \u0026Aring;\u003csup\u003e\u0026minus;1\u003c/sup\u003e (for general analysis). RIXS measurements were performed at room temperature under a vacuum of ~\u0026thinsp;10\u003csup\u003e\u0026minus;\u0026thinsp;8\u003c/sup\u003e mbar at the U41\u0026ndash;PEAXIS end-station of BESSY II (Berlin, Germany)\u003csup\u003e\u003cspan citationid=\"CR54\" class=\"CitationRef\"\u003e54\u003c/span\u003e\u003c/sup\u003e. The spectrometer was aligned in specular geometry at a 60\u0026deg; scattering angle and optimized to a total energy resolution of 90 meV using carbon tape. The excitation energy for each 1D spectrum was set to the O K-edge (531 eV), with an acquisition time of 20 min. For sample preparation, coin cells were charged/discharged to the target voltage, disassembled, and the electrodes were washed with DMC, dried, and mounted on double-sided Cu tape inside a N\u003csub\u003e2\u003c/sub\u003e-filled glovebox. The samples were then transferred to the measurement chamber using a N\u003csub\u003e2\u003c/sub\u003e-sealed transfer suitcase, ensuring no exposure to air. Data were processed using the beamline software, Adler-4.0. Soft XAS measurements were carried out at the Flexible PhotoElectron Spectroscopy (FlexPES) beamline of the MAX IV synchrotron (Lund, Sweden) under ultra-high vacuum at room temperature\u003csup\u003e\u003cspan citationid=\"CR55\" class=\"CitationRef\"\u003e55\u003c/span\u003e\u003c/sup\u003e. An exit slit of 10 \u0026micro;m was used, providing a photon flux of ~\u0026thinsp;10\u003csup\u003e12\u003c/sup\u003e photons s\u003csup\u003e\u0026minus;\u0026thinsp;1\u003c/sup\u003e and an energy resolution of 26 meV. The beam spot at the sample was 1 \u0026times; 0.5 mm\u003csup\u003e2\u003c/sup\u003e (defocused mode). Data were collected in TFY detection mode, and all spectra were background-subtracted and normalized.\u003c/p\u003e \u003c/div\u003e \u003cdiv id=\"Sec15\" class=\"Section2\"\u003e \u003ch2\u003eSEM, STEM and EELS measurements\u003c/h2\u003e \u003cp\u003eSEM images were taken with a Phenom Pharos Desktop SEM from Phenom world using an accelerating voltage of 10 kV and a secondary electron detector. The transmission electron microscope (TEM) images were collected on a JEOL JEM-3200FS operating at an accelerating voltage of 300 kV. The electron energy loss spectroscopy (EELS) was performed and atomically resolved high angle annular dark-field scanning transmission electron microscopy (HAADF-STEM) images were carried out on a double Cs-corrected TEM (FEI, Titan Themis G2) operated at 300 kV. TEM samples were prepared by FEI Scios microscope operated at 2\u0026ndash;30 kV.\u003c/p\u003e \u003c/div\u003e \u003cdiv id=\"Sec16\" class=\"Section2\"\u003e \u003ch2\u003eDFT calculations\u003c/h2\u003e \u003cp\u003eAll calculations were performed within the framework of Density Functional Theory (DFT) using the Vienna Ab initio Simulation Package (VASP)\u003csup\u003e\u003cspan citationid=\"CR56\" class=\"CitationRef\"\u003e56\u003c/span\u003e\u003c/sup\u003e. The Projector-Augmented Wave (PAW)\u003csup\u003e\u003cspan citationid=\"CR57\" class=\"CitationRef\"\u003e57\u003c/span\u003e\u003c/sup\u003e method was employed to describe the electron-nucleus interaction potential, while the Perdew-Burke-Ernzerhof (PBE) functional\u003csup\u003e\u003cspan citationid=\"CR58\" class=\"CitationRef\"\u003e58\u003c/span\u003e\u003c/sup\u003e was adopted for the exchange-correlation potential. A cutoff energy of 500 eV was set for the plane-wave basis set, and the DFT-D3 correction\u003csup\u003e\u003cspan citationid=\"CR59\" class=\"CitationRef\"\u003e59\u003c/span\u003e\u003c/sup\u003e was incorporated to account for weak intermolecular interactions. Structural optimizations were carried out for four crystal models, namely undoped, Sc doping, ScTi doping, and ScTiB doping systems, to ensure the models were in a thermodynamically stable state. For the electronic step calculations, the system was considered to have achieved energy convergence when the energy change was less than 10\u003csup\u003e\u0026minus;\u0026thinsp;5\u003c/sup\u003e eV. During the structural optimization process, the current configuration was regarded as a minimum point on the potential energy surface when the average atomic force was less than 0.05 eV/\u0026Aring;. The Partial Density of States (PDOS) of the optimized systems was calculated, and visualization of the results was conducted using the VESTA\u003csup\u003e\u003cspan citationid=\"CR60\" class=\"CitationRef\"\u003e60\u003c/span\u003e\u003c/sup\u003e.\u003c/p\u003e \u003c/div\u003e"},{"header":"Declarations","content":"\u003cp\u003e\u003cstrong\u003eAcknowledgements\u003c/strong\u003e\u003c/p\u003e\n\u003cp\u003eH.L. and Y.L. acknowledge the support from the Swedish Energy Agency (P2022-00055 and P2023-00603) and STandUP for Energy. P.A. and K.A.M. thank the Bundesministerium f\u0026uuml;r Bildung und Forschung (BMBF) for funding over project KAROFEST (03XP0498A) and TRANSITION (03XP0533C). P.A. also acknowledges support from the German Research Foundation (DFG, grant no. 501562980). J.W. and W.Z. are grateful for the financial support from the National Natural Science Foundation of China (52371225 and 92472115). The authors acknowledge DESY (Hamburg, Germany), a member of the Helmholtz Association HGF, for the provision of experimental facilities and beamtime (proposals I-20220574, I-20211173) with the kind support of Dr. Martin Aaskov Karlsen and Dr. Volodymyr Baran. This work was performed at DESY/PETRA III beamline P02.1, and Y.L. and Y.S. also acknowledge DESY for support with travel costs. Soft XAS and part of the \u003cem\u003eoperando\u003c/em\u003e XRD experiments were carried out at the FlexPES and DanMAX beamlines (Proposals no. 20241703 and 20241814), MAX IV Laboratory, Sweden. The authors gratefully acknowledge Dr. Jiefang Zhu, Dr. Dan Li, Mohammad Baghban Shemirani and Hao Wang from Uppsala University, as well as Dr. Frederik Holm Gj\u0026oslash;rup and Dr. Eleanor Frampton (MAX IV) for their valuable support during the beamtime. Hard XAS and RIXS were carried out at the KMC-2 and U41-PEAXIS beamlines at the BESSY II electron storage ring operated by the Helmholtz-Zentrum Berlin f\u0026uuml;r Materialien und Energie, GmbH, which is also acknowledged for provision of beamtime. We acknowledge Myfab Uppsala for providing facilities and experimental support. Myfab is funded by the Swedish Research Council (2020-00207) as a national research infrastructure.\u003c/p\u003e\n\u003cp\u003e\u003cstrong\u003eAuthor Contributions\u003c/strong\u003e\u003c/p\u003e\n\u003cp\u003eY.L., P.A. and H.L. conceived the idea, Y.L. synthesized and characterized the materials, performed Rietveld refinement studies, and all electrochemical measurements. \u003cem\u003eOperando\u003c/em\u003e XRD measurements were done by Y.L. and K.A.M. PDF measurements were performed by Y.L. and Y.S. Hard XAS experiments were performed by Y.L., K.A.M., Y.S., C.B., and M.Y. with the help of G.S. RIXS measurements were done by D.W. and Y.L. DFT was calculated by Q.Z. and J.W. STEM was performed by W.Z. H. Z., and J.W. The data and manuscript were organized by L.Y. and H.L. The manuscript was written by Y.L., H.L., J.W. and P.A. with input from all authors. All authors contributed to discussion of the results.\u003c/p\u003e"},{"header":"References","content":"\u003col\u003e\u003cli\u003e\u003cspan\u003eNayak PK, Yang L, Brehm W, Adelhelm P (2018) From Lithium-Ion to Sodium-Ion Batteries: Advantages, Challenges, and Surprises. Angew Chem Int Ed 57:102\u0026ndash;120\u003c/span\u003e\u003c/li\u003e \u003cli\u003e\u003cspan\u003eVaalma C, Buchholz D, Weil M, Passerini S (2018) A cost and resource analysis of sodium-ion batteries. Nat Rev Mater 3:18013\u003c/span\u003e\u003c/li\u003e \u003cli\u003e\u003cspan\u003eUsiskin R et al (2021) Fundamentals, status and promise of sodium-based batteries. 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J Appl Cryst 44:1272\u0026ndash;1276\u003c/span\u003e\u003c/li\u003e\u003c/ol\u003e"}],"fulltextSource":"","fullText":"","funders":[],"hasAdminPriorityOnWorkflow":false,"hasManuscriptDocX":true,"hasOptedInToPreprint":true,"hasPassedJournalQc":"","hasAnyPriority":true,"hideJournal":false,"highlight":"","institution":"","isAcceptedByJournal":false,"isAuthorSuppliedPdf":false,"isDeskRejected":"","isHiddenFromSearch":false,"isInQc":false,"isInWorkflow":false,"isPdf":false,"isPdfUpToDate":true,"isWithdrawnOrRetracted":false,"journal":{"display":true,"email":"[email protected]","identity":"nature-portfolio","isNatureJournal":true,"hasQc":false,"allowDirectSubmit":false,"externalIdentity":"","sideBox":"","snPcode":"","submissionUrl":"","title":"Nature Portfolio","twitterHandle":"","acdcEnabled":false,"dfaEnabled":false,"editorialSystem":"ejp","reportingPortfolio":"","inReviewEnabled":true,"inReviewRevisionsEnabled":false},"keywords":"sodium ion batteries, layered oxide cathode, cycling current-dependent, degradation pathway, oxygen redox, phase transition, synchrotron characterization","lastPublishedDoi":"10.21203/rs.3.rs-8109344/v1","lastPublishedDoiUrl":"https://doi.org/10.21203/rs.3.rs-8109344/v1","license":{"name":"CC BY 4.0","url":"https://creativecommons.org/licenses/by/4.0/"},"manuscriptAbstract":"\u003cp\u003eHigh-voltage operation boosts the energy density of sodium layered oxides by activating oxygen redox, but the underlying degradation mechanisms remain controversial, hindering rational materials design. Here we demonstrate that the applied cycling current determines which degradation pathway dominates, thereby resolving this long-standing inconsistency. Using synchrotron-based X-ray probes, advanced microscopic characterizations and first-principle calculation on a compositionally controlled set of Sc-, Sc\u0026ndash;Ti-, and Sc\u0026ndash;Ti\u0026ndash;B-doped O3-type oxides, it is revealed that low current preserves oxygen-redox participation and drives oxygen-loss-induced degradation, whereas high current suppresses oxygen redox and induces a transition to a structure-governed evolution pathway along simplified P3-type routes. This compositional set makes the current-dependent degradation behaviour experimentally distinguishable, allowing the oxygen-redox-driven and structure-governed modes to be clearly disentangled. These findings establish a unified current-governed mechanistic framework and provide guidance for designing durable high-voltage sodium-ion cathodes.\u003c/p\u003e","manuscriptTitle":"Identifying cycling current-dependent degradation pathways in high-voltage layered sodium–ion cathodes","msid":"","msnumber":"","nonDraftVersions":[{"code":1,"date":"2025-12-12 07:51:15","doi":"10.21203/rs.3.rs-8109344/v1","editorialEvents":[],"status":"published","journal":{"display":true,"email":"[email protected]","identity":"nature-communications","isNatureJournal":true,"hasQc":false,"allowDirectSubmit":false,"externalIdentity":"NCOMMS","sideBox":"Learn more about [Nature Communications](http://www.nature.com/ncomms/)","snPcode":"","submissionUrl":"https://mts-ncomms.nature.com/","title":"Nature Communications","twitterHandle":"","acdcEnabled":true,"dfaEnabled":true,"editorialSystem":"ejp","reportingPortfolio":"Nature Communications","inReviewEnabled":true,"inReviewRevisionsEnabled":false}}],"origin":"","ownerIdentity":"f55933b5-2a14-42b5-93a2-3b197f5b3491","owner":[],"postedDate":"December 12th, 2025","published":true,"recentEditorialEvents":[],"rejectedJournal":[],"revision":"","amendment":"","status":"under-review","subjectAreas":[{"id":59489897,"name":"Physical sciences/Chemistry/Electrochemistry/Batteries"},{"id":59489898,"name":"Physical sciences/Energy science and technology/Energy storage/Batteries"},{"id":59489899,"name":"Physical sciences/Materials science/Materials for energy and catalysis/Batteries"}],"tags":[],"updatedAt":"2026-02-06T13:00:36+00:00","versionOfRecord":[],"versionCreatedAt":"2025-12-12 07:51:15","video":"","vorDoi":"","vorDoiUrl":"","workflowStages":[]},"version":"v1","identity":"rs-8109344","journalConfig":"researchsquare"},"__N_SSP":true},"page":"/article/[identity]/[[...version]]","query":{"redirect":"/article/rs-8109344","identity":"rs-8109344","version":["v1"]},"buildId":"8U1c8b4HqxoKbykW_rLl7","isFallback":false,"isExperimentalCompile":false,"dynamicIds":[84888],"gssp":true,"scriptLoader":[]}