Decoupling slab gliding and lattice contraction in Na layered oxides to enable high-voltage Na-ion batteries

Domain: physics
Published: 2026-01-10
Editorial outcome: major_revision

Abstract
Layered transition metal oxide cathodes (NaxTMO2) demonstrate a classic type of cathode for Sodium-ion batteries (SIBs), however their practical application faces a long-standing challenge of irreversible phase transitions at high voltages, which causes unsatisfied specific energy and cycling stability, particularly for P-type (Na+ located at prismatic sites) cathodes. This phenomenon is conventionally ascribed to the Na+ re-coordination from prismatic to octahedral (O-type) configuration upon Na+ extraction, whereby the TMO2 slab gliding and abrupt c-lattice change are always coupled, and a straightforward solution to this situation remains elusive. Here, we reveal that, the TMO2 slab gliding and the lattice contraction can be decoupled, and the rapid lattice contraction under high state-of-charge underlies the fundamental origin for their irreversible phase transitions. By pre-engineering 15.8% O-type stacking faults to a P-type Na0.7Mn0.8Ni0.2O2, the dramatic volume variation and irreversible phase transitions at high voltage (4.5 V vs. Na+/Na) can be primarily eliminated. This work advances the understanding on the phase transitions at deep desodiation states, and paves up a feasible way to realize high-energy layered oxides.

Introduction
Sodium-ion batteries (SIBs) based on earth abundant resources demonstrate great application potential in grid-scale energy storage or fast charging scenarios, attributed to their low cost and fast charging ability1–3. Layered transition-metal oxides (NaxTMO2, TM refers to transition metals) with high theoretical capacities and good processability, have been well explored as SIB cathode materials, which can be generally classified into different types according to the coordination of Na (trigonal prismatic: P-type or octahedral: O-type) and the stacking of the oxygen layers (e.g., P2: ABBA or O3: ABCABC)4–6. The P-type materials with an open prismatic structure allow fast Na+ transportation and are inherently more stable than their O-type counterparts, however, their practical applications have to deal with one primary challenge of uncompetitive specific energy, attributed to their relatively low Na content (x is typically between 0.67 and 0.78)7,8. Even worse, the full charge storage capacity of these cathode materials cannot be fully utilized primarily because they undergo deteriorated P-type to O-type phase transitions along with a very large volume change at Na-depletion states (high SOC, Supplementary Fig. 1)9,10. The irreversible P-to-O phase transition is characterized by the gliding of the transition metal oxide (TMO2) slabs associated with a re-coordination of the Na+ ions from P-type to O-type, as well as a dramatic contraction of the unit cell along the c-axis11–13. Typically, the gradual removal of Na+ will decrease the Na shielding and present face-to-face O2− ions between adjacent TMO2 layers, which causes escalated O2−–O2− electrostatic repulsion and a continuous expansion of the NaO2 layer spacing. At high SOC, this expanded spacing is maintained solely by the repulsive forces between O2− ions, which further precipitates the gliding of TMO2 slabs to form thermodynamically stable O2-type stacking faults (SF). Such a process is invariably accompanied with a rapid and significant c-drop and a volume shrinkage of ~20%, thus triggers severe particle fragmentation and the consequent capacity decay upon cycling, especially at high SOC. A more reversible phase transition from P2 to OP4/Z phase can be induced at high SOC by introducing specific elements (Li+14, Ti4+15, Mg2+16, Zn2+17, Sb5+18, Ca2+19, etc.) or TM vacancies into the TMO2 layer. These OP4 or Z phases are considered to feature with O/P-type intergrowth phase, wherein the ratios between P-type layer and O-type layer vary with the charging/discharging potentials and cycles20,21. However, from the aspect of a single slab, a sudden c-drop and volume change still happen, which can cumulatively lead to inevitable structural failure at a cut-off voltage (usually >4.3 V) or after prolonged cycling. This casts some fundamental questions: how can the abrupt P-to-O phase transition or the related lattice contraction be avoided at the level of individual slabs? Is there a straightforward way to decouple the TMO2 slab gliding from the lattice contraction?

In this work, we break away from the conventional views by pre-engineering 15.8% O-type SF to a P-type Na0.7Mn0.8Ni0.2O2, and investigate the structural origin of the P-to-O phase transition and the accompanied chemo-mechanical structural degradations. We provide substantive evidence that the TMO2 slabs gliding and the volume change can be largely decoupled, while the significant c-axis expansion upon Na+ extraction and abrupt lattice contraction at high SOC serve as the primary drivers for the irreversible phase transitions. Distinguished from the SF formed during electrochemical charge/discharge, which could damage the structural stability, the pre-engineered SF in P-type layered oxides characterized with large interlayer distance (~5.6 nm) and O-type coordination, can avoid the presence of face to face O2− ions at high SOC, leading to minimal c-axis expansion during the Na+ extraction, which eventually eradicated the irreversible P2 to O2 phase transition. The as-prepared SF-rich Na0.7Mn0.8Ni0.2O2 (SF-NaMN) delivers an specific energy of 635 Wh kg−1 (based on cathode materials), and the as constructed hard carbon||SF-NaMN full cell demonstrates a specific energy of 297 Wh kg−1 (without presodiation) with capacity retention of ~80% after 600 cycles. A pouch cell and an all-solid-state sodium batteries were also assembled (without presodiation), which demonstrate high specific energy and stable cycling performance. This work clarifies the long-standing puzzle for the origin of P to O phase transition, and paves the way to rationally design high-voltage NaxTMO2 cathodes by the virtue of stacking-faults engineering.

Results
Designing of SF rich P2-type layered oxides
This work aims to decouple the TMO2 slab gliding from the c-lattice contraction, which generally happens co-currently during the charging and discharging process. To disrupt this inherent correlation, one can consider pre-establish certain amounts of O-type slabs into P-type layered structure with unchanged interlayer distances (O–Na–O). Conventionally, O- or P-type phases can be generally achieved by adjusting the cationic potentials through introducing different hetero elements3,4,19,22,23, but such strategies inevitably lead to varied c-lattice spacing (e.g., O3: 0.53–0.54 nm, O2: 0.44 nm, P2: 0.56–0.57 nm). In this work, we focus on the typical Mn-Ni-based P2-type cathode and intentionally disrupt the Mn-Ni ordering by manipulating the co-precipitation process of the (MnNi)CO3 precursor (see details in Supplementary Note 1). We refer the concept of the “cationic potential (Φ)”1 to elucidate how these O-type SF with remained c-parameter can be created in P2-layered oxides. Previous studies have proposed an empirical relationship (y=0.4065x+1.572), which is regarded as the P–O phase dividing line (PDL). According to this equation, the compounds with different composition would exhibit distinct cationic potentials, leading to different Na coordination modes contingent upon their position relative to the PDL. Further extending this description, we can draw lines parallel to the PDL, whereby the layered oxides on these parallel lines present similar coordination environments, termed “equal coordination lines” (ECL) in this article. It can be reasonably deduced that the compounds positioned on ECLs closer to the PDL are more likely to form P/O biphasic materials, whereas those far from the PDL tend to yield pure P-type or O-type phases. To validate this hypothesis, we calculated the cationic potentials of those P/O biphasic oxides reported from literature, which are all confirmed situated in the O–P biphasic region (Supplementary Fig. 2; Supplementary Tables 1 and 2). We then applied this analysis to individual TMO2 slabs. The calculated cationic potentials of Na0.7Mn1−xNixO2 (0≤x≤0.5, Fig. 1a) reveals that the Φ values vary with the Mn:Ni ratio with such a sequence: Φ(Na0.7Mn0.7Ni0.3O2) > Φ(Na0.7Mn0.5Ni0.5O2) > Φ(Na0.7Mn0.8Ni0.2O2) > Φ(Na0.7Mn0.6Ni0.4O2) > Φ(Na0.7Mn0.9Ni0.1O2) > Φ(Na0.7MnO2) with ECLs approaching the PDL line. Therefore, if the Ni/Mn elements are unevenly distributed in Na0.7Mn1−xNixO2, those Mn-rich TMO2 slabs could induce O-type Na coordination occasionally, which demonstrate the structural origin of the large amount of O-type SF formed in SF-NaMN. With this mind, the Na0.7Mn0.8Ni0.2O2 layered oxides with pre-constructed SF (denoted as SF-NaMN) were synthesized by regulating atomic arrangement of Mn and Ni ions (more detailed information are presented in “Methods”, Supplementary Note 1; Supplementary Figs 3 and 4). Laboratory and synchrotron X-ray powder diffraction patterns (XRD) show that SF-NaMN patterns exhibit obvious diffraction peak broadening on (10l) (e.g., (102), (103), (104), and (106)) planes, but the peaks of (00l), (h00) and (hk0) remain sharp (Fig. 1b; Supplementary Fig. 5 and Supplementary Table 3), indicating that SF-NaMN has a structural integrity along the ab-plane. Generally, the diffraction peak broadening can be originated from the crystallites size decrease, the presence of microstrain or certain type of defects etc24–27. Herein, we applied Williamson–Hall equation (W–H) to probe the origins for peak broadening of (10l) facets in SF-NaMN28–30. The fitted line of (10l) facets of SF-NaMN deliver a slope of 1.51%, which is 0.1% for the other facets with a line almost parallel to the x-axis in SF-NaMN, and 0.21% for all diffraction peaks of NaMN (Fig. 1c; Supplementary Tables 4; 5 and Supplementary Note 2). This indicate that, the broadening of the (10l) peaks should be attributed to the presence of SF along the (00l) facets in SF-NaMN. Furthermore, the percentage of the SF has been calculated as 15.8% according to the equation: βcosθ=0.89λ/D+2cλcosθcosφz4f (detailed information described in “Methods”). As contrast, NaMN did not show the peak broadening on (10l) planes. The Rietveld refinement of XRD patterns of SF-NaMN and NaMN reveals that the c-axis of the SF-NaMN (11.187 Å) exhibits a slight contraction (0.456%) in comparison to the that of NaMN (11.238 Å) (Supplementary Fig. 6; Supplementary Tables 6 and 7), which maybe induced by SF. Significantly, no characteristic peaks for O3-type structure were detected, evidencing that the O–Na–O interlayer distance at the O-stacking faults remain the same with the P-type structure. Figure 1d illustrates the P- and O-type configurations with the O–Na–O interlayer distance unchanged. It can be calculated that, under the same O–Na–O interlayer distance, the O–O distance across the intercalant layer will always be longer in O-type than P-type configurations31. Henceforth, the O2−–O2− electrostatic repulsion in the O-type SF can be significantly reduced upon Na+ extraction, which may avoid abrupt lattice contraction, and offers the possibility to decouple the TMO2 slab gliding and the lattice contraction.

Structural characterization
The morphology and microstructure of the samples were examined using scanning electron microscopy (SEM) (Supplementary Fig. 7) and aberration-corrected high-angle annular dark-field scanning transmission electron microscopy (Fig. 2a–d and Supplementary Fig. 8). Grains projected along the [100] zone axis were observed, with the large bright spots representing TM atom columns, but light elements, such as Na and O, are not detectable. Notably, a series of periodic regular SF can be found in SF-NaMN (orange box) with interlayer distance of 5.56 Å at stacking fault position, and 5.65 Å for other P-type layers (Fig. 2a). The space distances of the lattice fringes for NaMN are 5.70 and 2.57 Å corresponding to the (002) and (100) plane, respectively (Fig. 2b). The atomic-level structural information along the [001] zone axis further confirmed both TMO2 octahedron of SF-NaMN and NaMN aligned in a hexagonal symmetry arrangement (Fig. 2c, d). The vision of SF in SF-NaMN can be clearly distinguished in [001] direction, which shows ordered honeycomb structure. The bright spots in blue circles represent TM atoms occupied 2a sites of P2-layered hexagonal structure, while the slight spots (green and orange lines) correspond to TM atoms at octahedral 2b sites in SF slabs (Fig. 2c, inset). These SF should be caused by the gliding of a TMO2 slab along [1/3, 2/3, z] or [2/3, 1/3, z] direction, leading to the coordination environment change of Na+ from trigonal prismatic position to octahedral positions. Besides, the horizontal distances between adjacent TMO2 octahedron of SF-NaMN under different signal intensities are similar (5.03 Å), indicating the gliding of TMO2 slabs induced by SF remain integral P2-type structure without obvious variation in atomic spacing (Fig. 2c). As contrast, the horizontal distance between TMO2 octahedron of NaMN is ~5.10 Å without extra signals generated from [001] projection, indicating single-crystalline hexagonal structure of NaMN without obvious SF (Fig. 2d). According to the above structure features observed, we built structural models to understand the special atomic arrangement of SF-NaMN (Fig. 2e, f). It indicates that the P2-Na0.7Mn0.8Ni0.2O2 lattice and O-type SF are compatibly combined by sharing identical closely packed O sublattices in a pseudo-hexagonal scheme, which matches well with the STEM atomic pattern. Notably, the interlayer distance of TMO2 slabs at stacking fault layer in SF-NaMN is close to that the P-type phase (5.56 Å v.s. 5.65 Å), which different from the noticeable decrease in interlayer distance of (002) facet observed in O2 phases (4.3–4.4 Å) during deep desodiation process13. Further characterizations were conducted to explore the possible mechanism for the formation of such SF rich structure. Atomic-resolution energy dispersive spectroscopy (EDS) mappings (Fig. 2g–j) were conducted to distinguish the distribution of Mn and Ni atoms within each sample beneath the projected atomic columns. By comparing the counts of EDX spectra, the ratios of transition metals at the two SF region are approximately Mn:Ni=0.815:0.185 and 0.789:0.211; as well as 0.789:0.211 and 0.805:0.195, respectively. These values correspond to a to d TMO2 slabs (counted from top to bottom in Fig. 2h). The ionic potential differences of the transition metals between the two layers at the stacking fault are 0.3552, and the cationic potential difference is 0.08532 with constant Na content (0.7) (Supplementary Fig. 9). It is revealed that Mn atoms demonstrate enhanced intensity at SF accompanied with weakened Ni intensity (Fig. 2h), suggesting that there is a variation of Mn-Ni contents among different TMO2 slabs in SF-NaMN. The obvious difference of Mn/Ni ratio at certain TMO2 slabs might be the fundamental origin of SF, which is consistent with our design. As contrast, NaMN sample shows homogeneous Mn and Ni distribution within each layer. (Fig. 2i, j).

The local structure of TM-ordering in SF-NaMN and NaMN was further examined by X-ray pair distribution function (PDF) and neutron diffraction (ND) patterns. The local distances of TM-O (orange curve), and TM-TM (green curve) were simulated by PDFgui software (Supplementary Fig. 10). It can be found that the short-range ordering of TM-TM pair of SF-NaMN has obvious decay compared with that of NaMN. Moreover, the TM-ordering reflections experience a noticeable decrease in SF-NaMN than that in NaMN, as evidenced by the reflections around 3.0–4.6 Å in (ND) patterns (Fig. 2k)32, which indicating that the TM (Mn-Ni) arrangement of SF-NaMN are obviously more disordered than that of NaMN. Combining the EDS-mapping and ND results, SF-NaMN synthesized by our methods could induce Mn atoms enriched in several layers of layer oxides, lead to the different TM contents and corresponding cationic potential between adjacent TMO2 slabs. The large cationic potential difference between adjacent TMO2 slabs give rise to non-equilibrium electrostatic repulsion of O2−–O2−, thus leading a couple of weakened Na–O bonds and a pair of strong Na–O bonds3. In this case, Na+ occupation in P-phase trigonal prismatic position is unstable and tend to move to octahedral positions driving by non-equilibrium O–Na–O interaction, leading to the occurrence of O-type SF.

To verify the generality of this synthesis method, samples doped with varied metallic ions, such as Na0.74Mn0.75Ni0.2Li0.05O2, Na0.64Mn0.77Ni0.18Ti0.05O2, Na0.7Mn0.75Ni0.2Mg0.05O2, Na0.67Mn0.75Ni0.21Cu0.04O2, Na0.67Mn0.75Ni0.2Al0.05O2, Na0.69Mn0.74Ni0.19Fe0.07O2, Na0.73Mn0.76Ni0.18Zn0.06O2 were synthesized using the same method. The diffraction peak of (102), (103), (104) and (106) facets of these samples were also obviously broadened compared with their control samples (Supplementary Fig. 11 and Supplementary Table 8), indicating that such a synthetic approach is viable to regulate the crystal structure of P2-phase materials with diversified compositions.

Electrochemical performance
The electrochemical performance of SF-NaMN and NaMN samples (Supplementary Data 1) was tested at varied potential ranges (vs. Na+/Na) using sodium metal as the anode (Fig. 3a–c). The as prepared SF-NaMN deliver a capacity of 130 mAh g−1 in the voltage range of 2.4–4.5 V (Fig. 3a), and 227 mAh g−1 in a wide voltage range of 1.5–4.5 V (Supplementary Fig. 12) at 0.1 C (1 C = 180 mAh g−1), corresponding to a high specific energy of 635 Wh kg−1 (calculated based on active mass of the cathode material). The peak intensities of Na+/vacancy ordering (3.78/3.64 V) are significantly reduced for SF-NaMN compared to that of NaMN, suggesting that pre-constructed SF can effectively suppress the formation of Na+/vacancy ordered structures. The CVs and dQ/dV curves of SF-NaMN and NaMN show two pairs of redox peak in the high voltage region (4.11/4.38 V) corresponding to the Ni3+/Ni4+ and O2−/On− redox reaction5,6,29 (Supplementary Fig. 13), which is confirmed by ex situ X-ray absorption spectroscopy (XAS) and X-ray photoelectron spectroscopy (XPS) (Supplementary Figs. 14; 15 and Supplementary Note 3). The galvanostatic intermittent titration techniques (GITT) (Fig. 3b and Supplementary Fig. 16) results show no obvious drop on the DNa+ value (10−9.69→10−9.89 cm2 s−1) at high voltage (4.5 V) for SF-NaMN, evidencing the improved ion diffusion capability for SF-NaMN. On the contrary, NaMN experienced two orders of magnitude decrease (10−9.89→10−11.61 cm2 s−1) of the DNa+. Consequently, SF-NaMN demonstrates a capacity retention of 86.1% after 100 cycles at 1 C, surpassing to that of NaMN (64%, Supplementary Fig. 19). The electrochemical performance comparison of cathode doped with varied metallic ions were also examined and summarized in Supplementary Figs. 20 and 21. The GCDs of Li, Ti, Mg, Zn, Cu, Al and Fe doped SF-NaMN cathodes show smoother charge-discharge curves in the high voltage range of 4.2–4.5 V, distinguishing from their corresponding control samples. Meanwhile, the rate performance, coulombic efficiency (CE) and cycling stability of these cathode oxides are further enhanced (Supplementary Fig. 22).

As has been established that P2-type electrodes usually deliver a low initial charge capacity of ~85 mAh g−1 below 4.0 V due to the low Na content (<0.75) in the molecular9,15, and a low average voltage <3.2 V because of the irreversible P to O phase transition at high voltage33,34, thus compromising the overall specific energy of the battery. Therefore, in order to achieve a decent specific energy, a troublesome pre-sodiation process or adding self-sacrificed additives becomes necessary for the full-cell construction of P2-type cathodes35–37. Considering the incredible reversibility of SF-NaMN under high voltage could enable more Na+ ions participate in the electrochemical reactions, full cells were assembled using unpresodiated hard carbon (HCups) as anode (HCups||SF-NaMN) (Supplementary Fig. 23). The as-equipped HCups||SF-NaMN delivered a high specific capacity of 115.5 mAh g−1 (calculated based on the active mass of cathode electrode) with initial CE of 82.4% at a high average operating voltage of 3.5 V at 0.1 C (Fig. 3d), and a high rate performance with 69.2 mAh g−1 capacity remain at 4 C (Supplementary Fig. 24). Most importantly, the HCups||SF-NaMN demonstrated a improved cycling stability at varied voltage ranges (Fig. 3e and Supplementary Fig. 25), the capacity retention maintained 96.2% and 87.3% within 1–4.4 V after 50, 100 cycles, and 77.3% within 2–4.2 V after 600 cycles at 1 C respectively with negligible voltage fade of GCD curves (Supplementary Fig. 26), far surpassing that 71.9%, 56.7% within 1–4.4 V after 50, 100 cycles and 67.2% within 2–4.2 V after 229 cycles) of HCups||NaMN. The specific energy of HCups||SF-NaMN was calculated as 297 Wh kg−1 (based on the total active mass of both electrodes). Furthermore, a full cell with pre-sodiated hard carbon (HCps) and SF-NaMN (HCps||SF-NaMN) was also constructed (Supplementary Fig. 27), which delivered an specific energy of 303.7 Wh kg−1 (based on the total active mass of both electrodes). The specific energy of both HCups||SF-NaMN and HCps||SF-NaMN is comparable to full-cells constructed with P2-layered oxide cathodes and pre-sodiated HCs (Fig. 3f and Supplementary Table 9). A side-by-side comparison of full-cell performance was further conducted by re-synthesis several literature reported materials and assembling full-cells under the exact same conditions, which suggesting enhanced high voltage and full-cell performance (without presodiation) (Supplementary Fig. 28 and Supplementary Note 4). Intriguingly, the SF-NaMN exhibits optimized compatibility with NASICON (Na3Zr2Si2PO12) solid electrolyte, which achieves a high specific energy of 437.6 Wh kg−1 within 2.4–4.5 V (Supplementary Fig. 29). The all solid-state full cell can stably run at a high cutoff voltage of 4.2 V with retaining 81% of the initial capacity after 300 cycles at 0.5 C (Supplementary Fig. 29d).

To have an in-depth understanding on the electrochemical performance of SF-NaMN, the difference of the Na+ kinetic transport between SF-NaMN and NaMN at high voltage (>4.2 V) in the full cell system was systematically studied through the distribution of relaxation times (DRT) analysis, which employs a model-free approach to analyze impedance evolution. This mathematical transformation converts frequency-domain Nyquist plots into time-domain DRT spectra, enabling a model-independent assessment of electrochemical processes. Z(ω)=R1+∫0∞γ(τ)1/(1+jτω)dτ. Peaks at specific relaxation times correspond to distinct electrochemical processes, while their respective areas quantify the associated impedance values. The relaxation time (τ) is the characteristic timescale required for a system to recover from a non-equilibrium state to an equilibrium state, which reflect the rate of electrochemical reactions or ion diffusion. The electrochemical impedance spectra (EIS) are resolved by this analysis with different electrochemical processes38–41, six main peaks located at 10−6–10−4 (R0), 10−3–10−2 (Rfilm), 10−1–100 (Rct1, Rct2), 100–101 (ZW0), in accordance with ohmic resistance, resistance of solid electrolyte interphase (SEI), (refer to the film atop of anode) and cathode electrolyte interphase (CEI) (film, refer to the film atop of cathode), charge transfer resistance of anode material, charge transfer resistance of cathode material and Warburg resistance, respectively39,42,43. (high and intermediate/low frequency region etc) (Fig. 3g–j and Supplementary Fig. 30). As Na+ deep desodiated, Rct1 for HCups||SF-NaMN gradually increased and moved towards longer relaxation time, this could be attributed to the slow Na+ transport kinetics within the almost fully sodiated HC at 4.35–4.4 V39. Notably, the Warburg resistance (ZW0), which characterizes the bulk Na⁺ diffusion behavior, showed no obvious variation during the first cycle, indicating that the diffusion process remains relatively stabled during charging process. The Rct2 and ZW0 resistance could return to their original response position during the discharging process. This reversibility of the impedance-derived parameters serves as compelling evidence for the optimized electrochemical reversibility exhibited by the HCups||SF-NaMN. In contrast, for the HCups||NaMN full cell, both Rct2 (charge transfer resistance of cathode material) and Warburg resistance (ZW0) increased significantly during charging process and cannot completely recover during the discharging process, indicating the charge transfer and bulk Na+ diffusion process were severely inhibited by irreversible phase transition of NaMN.

A 0.5 Ah pouch cell are assembled to further develop the practical potential application of SF-NaMN (Fig. 3k and Supplementary Table 10). Kilograms of SF-NaMN were synthesized to fabricate a pouch cell via the same condition as that used for the gram-level reaction kettle (Supplementary Figs. 31 and 32). Moreover, HC||SF-NaMN pouch cell can operate for 340 reversible cycles with >80% capacity retention within 2.0–4.2 V at a practical level (Fig. 3l), suggesting great practical potential of SF-NaMN. These results reveal that SF-NaMN exhibit enhanced Na+ kinetic transport capability at high voltage in the SIB full cell.

Investigation of the decoupling mechanism between slab gliding and lattice contraction
To understand the structural evolution of SF-NaMN and NaMN during the charge and discharge process, in situ XRD measurements are conducted. In situ XRD patterns illustrate that, the characteristic (002’) peaks at 20° and 17–19°, corresponding to O2 and OP4 phases, do not show during the charge and discharge process of SF-NaMN, indicating irreversible P to O phase transitions was not observed in SF-NaMN (Fig. 4a). Notably, a slight decrease in the intensity of the (104) peak at 39.75 degrees at the end of charging, demonstrating relative interlayer gliding between the TMO2 slabs happened, leading to atomic displacements and broadening of the (104) peak. However, the (002) peak did not shift, and no new peaks formed, indicating that no obvious interlayer contraction of SF-NaMN occurred during this process. The c-axis variation for SF-NaMN is calculated as 0.9% at 4.0 V, and 1.45% within 2.4–4.5 V (Supplementary Figs. 33–35). As a comparison, new peaks (the (002’) peak at 18.3° and (101’) peak at 37.9°)12 related to the OP4 structure were observed in the XRD patterns of NaMN at 4.25 V, which were enhanced with the increasing voltage, meanwhile, the P-phase peaks were gradually weakened, indicating a phase transition from P2 to Z occurred at high voltage region. Consequently, the NaMN shows an obvious c-axis expansion (2.7% at 4.0 V, 12.8% within 2.4–4.5 V, Fig. 4b, more detailed expression in Supplementary Note 5). Such terrible phase transition can trigger severe irreversible structural structure degradation, eventually unsatisfied electrochemical performance. HADDF-STEM examination visually evidenced the P/O stacking mode at the atomic scale of SF-NaMN and NaMN when charged to 4.5 V. For SF-NaMN, although some TMO2 undergone relative gliding, while the TMO2 slab still maintains interlayer distance of 5.3–5.7 Å, consistent with P2-type structure (Fig. 4c). In contrast, two types of interlayer distances can be statistically depicted in NaMN, with ~5.7 Å (P2-type) and ~4.6 Å (O2-type) in the highlighted image of Fig. 4d. The formation of O-type structure of NaMN can be ascribed to the rapid c-parameter contraction and subsequent P–O phase transition during deep desodiation. Thus, NaMN undergoes a severe irreversible P2–Z phase transition at 4.5 V, while SF-NaMN remains a P2-type structure without any phase transition, which is consistent with the in situ XRD results. Moreover, geometric phase analysis (GPA) SF-NaMN exhibit significantly attenuated interlayer force changes, as evidenced by its uniformly distributed weak structural strain (Supplementary Fig. 36). In stark contrast, pristine NaMN shows severe strain localization with disordered strain vectors, which correlates with uncoordinated gliding and contraction of TMO2 slabs. Therefore, although 15.8% TMO2 slabs have been pre-glided from P to O position, no significant c-contraction was observed in SF-NaMN. In this sense, the traditional correlation between TMO2 slab gliding and c contraction has been decoupled.

Density functional theory (DFT) calculations (Supplementary Data 2–6) were conducted to understand the relationship between SF and P-to O phase transition. Two model systems, Na0.7Mn0.8Ni0.2O2 (Na21Mn24Ni6O60) with homogeneous Ni/Mn distribution structure (representing NaMN), and Na0.7Mn0.8Ni0.2O2 (Na21Mn24Ni6O60) with inhomogeneous Ni/Mn distribution (representing SF-NaMN) were established based on the structural investigation results in Figs. 1 and 2. Specifically, a Mn-rich TMO2 slab was constructed to represent the Mn-rich and Ni-deficient character of TMO2 slabs, according to the EDS mapping result of SF-NaMN. The optimal structures of NaMN and SF-NaMN before desodiation have been investigated. The structural model of SF-NaMN with perfect P2-type input structure showed SF during the structural optimization process. The stacking faults occurred near the Mn-rich TMO2 slab, and the density is 16.7%. However, no SF are observed during the structural optimization process of NaMN with a perfect P2-type input structure (Supplementary Fig. 37). These indicate that SF exist for SF-NaMN. In contrast, NaMN is free from stacking faults before desodiation. Furthermore, the stacking faults located between two adjacent Mn0.8Ni0.2O2 slabs within SF-NaMN have been examined. Interestingly, the stacking fault between two adjacent Mn0.8Ni0.2O2 slabs is unstable, and eventually evolved to reside between MnO2 and Mn0.8Ni0.2O2 slabs after the optimization process, indicating that SF in SF-NaMN might be induced by the variation in TM content between adjacent TMO2 slabs, which in accordance with HADDF-STEM and EDS results.

Then, the role of SF in SF-NaMN on the structural evolution during desodiation processes was studied. Na0.3Mn0.8Ni0.2O2 (Na9Mn24Ni6O60) was established for studying the desodiation states of SF-NaMN and NaMN, with sodium content roughly equivalent to when the material is charged to 4.25 V (Fig. 4e; Supplementary Fig. 38 and Supplementary Note 6). The most stable desodiation configuration for NaMN is achieved through the uniform extraction of Na+ ions from each sodium layer. However, the desodiation configuration for five out of twelve Na+ extracted from the SF slab shows the minimum energy (−628.90 eV) in SF-NaMN (Fig. 4e), suggesting that Na+ is prone to be extracted from the SF slabs in SF-NaMN. This phenomenon can be attributed to the varying compositional characteristics of the TMO2 slabs in relation to their proximity to the SF slab. The SF slab, which is in close proximity to the TMO2 slab with Mn3+/4+ redox couples at low voltage, would show a preference for Na+ extraction. The variation in the c-parameter for a deeper desodiation state Na0.2Mn0.8Ni0.2O2 (Na6Mn24Ni6O60) has been investigated. The variation of the c-parameter from Na0.7Mn0.8Ni0.2O2 to Na0.2Mn0.8Ni0.2O2 for SF-NaMN is 0.097%, which is significantly lower than that of NaMN (1.39%), indicating that SF slabs could effectively inhibit the expansion of c-parameter during desodiation (Fig. 4f). These results agree well with the in situ XRD results in Fig. 4a. We also constructed the models with SF of 8.33% and 12.5% and demonstrated that as stacking fault density decreases, the of c-axis variation of the structure during desodiation gradually increases (Supplementary Figs. 39; 40 and Supplementary Note 7).

The TMO2 slab gliding behavior of TMO2 slabs in Na0.2Mn0.8Ni0.2O2 (Na6Mn24Ni6O60) for SF-NaMN and NaMN along [1/3, 2/3, z] direction were further simulated. It should be noted that the slab gliding energy is calculated as the energy of the structure after gliding minus the energy of the structure before gliding, and this energy value is closely related to the propensity of the material to undergo the P–O transition: a lower energy value indicates that the slab structure is more likely to glide, and thus the material is more prone to transform from P-type to O-type. From the simulation results, the gliding energies for slab-1 and slab-2 are 0.249 and 0.375 eV of SF-NaMN, respectively, which are much greater than the energy required of slab-3 (−0.449 eV) of NaMN (Fig. 4g), indicating that NaMN is more prone to undergo a transition from P-type to O-type under deep desodiation.

Moreover, we successfully achieved a cathode with 12.1% SF by increasing the precipitation pH of the precursors to 9.62. The synthesized sample exhibited a noticeable slope change in their electrochemical profiles at the high voltage plateau of 4.2 V, and the capacity retention was obviously improved compared with NaMN (Supplementary Fig. 41 and Supplementary Note 8). Therefore, pre-constructing SF in P2-type lattice structure is an effective strategy to enhance the cycle life of the cathode material.

The Na+ desodiation behavior and evolution of the P–O phase transition for SF-NaMN and NaMN is schematically illustrated in the Fig. 4h. Previous work has reported that the escalating O2−–O2− electrostatic repulsion would lead to the expansion of the Na layer spacing with Na+ gradually extracted, and oversized Na slabs fail to sustain the P-type structure, thus lead to the re-coordination of Na+ to form thermodynamically stable O-type stacking, accompanied by the sudden c-parameter contraction4,5,8. In our work, the interlayer spacing for the pre-constructed O-type SF in SF-NaMN is ~5.56 Å as evidenced by the experimental results (Fig. 2a), which is similar with that of the P2 phase (~5.65 Å). Therefore, the O–O bond in the adjacent transition metal layer is elongated at the SF layer leading to a significant decrease in the electrostatic repulsion between O2−–O2−. As a result, the interlayer distance increasing caused by the O2−–O2− repulsion during Na+ extraction becomes not obvious in SF-NaMN, leading to negligible tensile strain between the adjacent TMO2 slabs and thus avoiding the sudden c-drop of the materials. In contrast, the control sample (NaMN) experiences strong O2−–O2− repulsion between adjacent TMO2 slabs during Na+ desodiation, resulting in the stretched prismatic Na+ coordination, which eventually induces the sudden cell contraction and subsequent phase transition upon a large amount of Na+ extraction. Therefore, we have significantly decoupled slab gliding within the material from lattice contraction by pre-engineering O–SF with large interlayer distance in the P2-type structure, demonstrating a much smaller c-axis contraction of the SF-NaMN with slabs gliding relative to NaMN (0.456% vs. 12.8%, respectively), with the contraction of SF-NaMN occurring prior to electrochemical desodiation. The phase transition at high SOC can be avoided under the presence of plenty of O-type coordination, as long as the sudden c-drop doesn’t occur.

Investigation of structural evolution during cycling
According to previous reports, the irreversible P–O phase transitions of P2-type layered oxide cathode during the high voltage region can induce grain cracking, TM ions dissolution, and severe side reactions between the cathode material and high oxidizing molecules in the electrolyte solvent4,5,44. The structural evolution of SF-NaMN and NaMN after cycling was monitored using HADDF-STEM (Fig. 5a–f and Supplementary Fig. 42). The SF-NaMN and NaMN electrode are tested with specific current of 0.2 C in 2.4–4.5 V vs. Na+/Na. It can be observed that the layered structure of SF-NaMN retained well with only a small amount of strain relaxation disclosed by the GPA after 30 charge/discharge cycles, showing an interlayer spacing of 5.67 Å (Fig. 5a–c). On the contrary, the cycled NaMN demonstrated obvious microcracks and surface degradation, as well as a severe Na+/TM intermixing rock-salt phase (Fd3m) with a perplexing strain distribution growing along the (00l) facets (Fig. 5d–f, more detailed information in Supplementary Note 9). The CEI structure of SF-NaMN and NaMN after 30 cycles were detected by the time-of-flight secondary ion mass spectrometry (ToF-SIMS). The TOF-SIMS depth profiles reveal the inorganic species (NaF−) and organic species (CHO−) showed obvious segregation and with fewer organic species than inorganic species constituting the surface of the CEI of SF-NaMN (with sputtering time <60 s, Fig. 5g and Supplementary Fig. 43), while no apparent segregation was observed between inorganic species (NaF−) and organic species (CHO−) on the surface of NaMN (Fig. 5i). Not ing that, the thickness of the CEI in the SF-NaMN is only two of third of that in NaMN according to the half-normalized maximum of the normalized intensity of the CEI species. Moreover, NiF3− and MnF3− fragments mainly comes from the cathode oxide dissolution driven by acidic attack (HF, etc). NaMN exhibit obviously stronger NiF3− and MnF3− intensity than that of SF-NaMN, confirming severe transition metal dissolution and cross-contamination of NaMN after cycling (Fig. 5h, j)45. Energy dispersive x-ray spectroscopy (EDX) of the elemental information on the Na anode reveal that the residual TM element signals in SF-NaMN (Mn: 1.37%, Ni: 1.16%) are significantly weaker than those in NaMN (Mn: 0.23%, Ni: 0.4%), also indicates that the TM dissolution in SF-NaMN has been significantly suppressed (Supplementary Fig. 44). The TEM results further confirm the robust CEI of about 2–3 nm formed on the surface of SF-NaMN cathode after 30 cycles. However, the CEI interface of NaMN is severely degraded and uneven after cycling, leading to significant surface degradation and crack formation (Supplementary Fig. 45). The electrochemical impedance spectroscopy (EIS) measurements (Supplementary Fig. 46 and Supplementary Note 10) suggest that SF-NaMN exhibit improved ion conductivity than NaMN after 50, 100, and 200 cycles, indicating that the CEI film of SF-NaMN is relatively stable and exhibits fewer side reactions with the electrolyte than that of NaMN.

To intuitively monitor the influence of high-voltage electrochemical behavior of SF-NaMN and NaMN on gas evolution in batteries, in situ differential electrochemical mass spectrometry were employed during the first charging process (Supplementary Fig. 47). Carbon dioxide (CO2) evolution was observed at 4.331 V vs. Na+/Na during the first charge of SF-NaMN, which is significantly later than that of NaMN (occurring at 4.171 V vs. Na+/Na). Noting that no obvious O2 release for SF-NaMN and NaMN, but a less considerable amount (CO2) release (308 μmol g−1) of SF-NaMN compared to that of NaMN (937 μmol g−1). According to the previous report, CO2 are generated from the reaction between the carbonate electrolyte with the highly active oxygen species on the surface or O2 evolved from the lattice oxygen of the cathode oxide46–49. Hence, the amount of CO2 is highly correlated with the release of O2. These results suggest that the degradation of P2-type layered oxide cathode and subsequent severe side reactions during high voltage region could attitude to the irreversible P–O phase transitions, resulting in unstable SEI (CEI) composition and gas evolution (Fig. 5k). SF-NaMN with pre-constructed O-type SF can effectively suppress these side reactions, addressing practical applications of positive electrode materials at the high voltage region.

In summary, we study a long-standing puzzle about the high voltage phase transition ever since the application of P-type transition layered oxides as cathode materials for SIBs: the abrupt lattice contraction underlies the fundamental origin for the irreversible phase transitions, while the TMO2 slab gliding may not be necessarily correlated. A certain amount of O-type SF was pre-established in P2-type layered oxides with interlayer distances (O–Na–O) unchanged by modulating the cationic potentials of the compounds, which can prevent sudden c-drop under high SOC. Besides, such a structure can significantly suppress the surface degradation and form a stable CEI film, thus stabilizing the Na+ (de)intercalation reaction during the cycling process. The as-prepared SF-NaMN demonstrates single P2 phase with small c-axis change (1.45%) at a high voltage up to 4.5 V, and high temperature of 50 °C. The as assembled HC||SF-NaMN full cell delivered a high specific energy of 297 Wh kg−1 without presodiation, and stably cycled over 600 cycles. A pouch-type cell and an all-solid-state full cell were also constructed based on SF-NaMN, delivers optimized cycling stability (340 cycles within 2–4.2 V) and a high specific energy of 437.6 Wh kg−1, respectively. This work sheds lights on achieving high specific energy and high-voltage P2-type NaTMO2 cathodes.

Methods
Materials preparation
We synthesized precursors of Na0.7Mn0.8Ni0.2O2 layered oxides with pre-constructed SF via a modulated co-precipitation method, leveraging the complex reactions between NH4+ and TM ions. In detail, 0.05 mol of nickel acetate (Aladdin, >99.0 %) and 0.1 mol of manganese acetate (Aladdin, >99.0%) were dissolved in deionized water in the required stoichiometric proportion and magnetically stirred at 25 °C for 1 h to yield a clear, homogeneous solution designated as A. 0.01 mol NH4HCO3 (Aladdin, >99.0%) was employed as precipitants, while 10 ml NH4OH (AR) was used to regulate the pH during precipitation, resulting in mixed (MnNi)CO3 precursors for subsequent high-temperature calcination. During the precipitation, the NH4+ ions can readily complex with Ni2+ and form [Ni(NH3)6]2+, which competes with the precipitation of NiCO3. By precisely controlling the pH to 8.9, a dynamic equilibrium between precipitation and complexation was achieved, disrupting the ordered arrangement of Mn and Ni in (MnNi)CO3 during repeated cycles of precipitation, dissolution, and reprecipitation of Ni ions. SF-NaMN is obtained after high-temperature calcination process. For comparison, a control sample with the same stoichiometric composition was synthesized using a conventional fast co-precipitation method with Na2CO3 as the precipitant. In the precipitation process, the reaction temperature was maintained at 12 °C. The precursor material was thoroughly washed several times by filtering with deionized water to remove residual sodium species, then dried in a vacuum oven at 120 °C overnight. Next, mixed the carbonate precursor of 0.01 mol with 0.004 mol Na2CO3 into planetary ball mills and milling with high energy in air with 1000 rpm at 25 °C for 6 h. Then, the powers were calcined at 500 °C for 10 h and then sintered at 900 °C for 12 h in Muffle furnace with heating rate of 5 °C/min and cooled down to room temperature. Li, Ti, Mg, Al, Cu, Fe, Zn-(SF-)/NaMN were prepared by mixing the precursor of SF-NaMN and NaMN with corresponding metal salts (Li2CO3, TiO2, MgO, Al2O3, CuO, Fe2O3, ZnO, Aladdin, >99.0%), subsequently calcined under the same thermal conditions. NASICON (Na3Zr2Si2PO12) solid-state electrolytes were purchased from HeFei Kejing company. Hard carbon (average particle size: 1.1 nm) are in-house prepared via bottom-up ZnO-assisted bulk etching strategy50. Polyvinylidenefluoride (PVDF, 99.5%), N-methylpyrrolidone (NMP, 99.5%), acetylene black (99.5%, average particle size: 300 nm), Al foil (20 μm, 99.5%) and Cu foil (18 μm, 99.5%) were all purchased from Alfa Aesar. Sodium discs (diameter: 16 mm, 0.5 mm thick) were punched from freshly rolled 99% Na metal (Alfa Aesar). The electrolyte, 1 mol NaPF6 in EC:PC with 5 % FEC (H2O≤20 ppm), is supplied by KELUDE company.

Material characterizations
Inductively coupled plasma atomic emission spectroscopy (ICP-AES, PERKIN E7300DV) was used to analyze the elemental composition of the samples. X-ray powder diffraction (PXRD) patterns were collected with Cu Kα radiation (λ=1.5418 Å). In situ XRD studies, Al foil were prepared as X-ray windows with a Swagelok cell during charge and discharge process with scan rate of 5°/min, step size is 0.02°. The in situ synchrotron XRD characterizations were also performed using the 14B1 beamline (λ=0.6887 Å) at Shanghai Synchrotron Radiation Facility Inc, China. Rietveld refinements were used to analyse the PXRD data with TOPAS software, which were collected at room temperature using a Bruker D8 ADVANCE diffractometer with Cu Kα radiation (40 kV, 40 mA) at Shanghai Institute of Ceramics. The morphology of all samples are investigated by a scanning electron microscope (SEM, 30 kV for EDX capture). Atomic-scale microstructure was examined with a JEOL JEM-Grand ARM300F, Cs-corrected STEM running at 300 kV in HAADF mode. Synchrotron XANES spectra at the Mn and Ni K-edges were collected in transmission at the Shanghai Synchrotron Radiation Facility on the IDEAS (Ge 220, IDEAS, double-crystal monochromator) and SXRMB (Si 111 double-crystal monochromator) beamlines, respectively. TOF-SIMS is conducted in negative ion mode with a region size of 80 micrometers×80 micrometers, a depth of 200 nm, and signal intensities of all species in TOF-SIMS have been normalized. To prevent air exposure, the sample was moved directly from the glove box into the ToF-SIMS vacuum chamber via a sealed transfer module. Depth profiling was then carried out by sputtering with a 2 kV, 100 nA Ar+ beam. XPS measurements were performed using an RBD-upgraded PHI-5000C ESCA system with Al Kα radiation and a testing duration of 1 h. XPS were calibrated by setting the binding energy for the hydrocarbon (C–C/C–H) in C 1s spectra to 284.8 eV. Ex situ XPS and XAS analysis was preceded by a two-step EMC rinse of the dismantled electrodes, followed by drying inside the argon glovebox. While still under the same inert atmosphere, the specimens were loaded into a hermetic analysis chamber that was immediately sealed and transported to the spectrometer.

Electrochemical measurement
The electrode slurry was prepared by dissolving the active material, acetylene black, and polyvinylidenefluoride (PVDF) at a mass ratio of 8:1:1 in N-methylpyrrolidone (NMP). It was then coated on one side of an aluminum foil with a loading density of 10–12 mg/cm2, followed by vacuum drying at 80 °C for 12 h. Metallic sodium was prepared by roll pressing followed by cutting into pieces (Ф12 mm). The 2032 coin cells were prepared in an argon-filled glove box. Using sodium tablets as counter electrode and the Glass fiber (GF/D, Whatman, Ф19 mm) was used as the separator in coin cell and PP@ceramic separator (JCK16 ceramic coated (double-face) PP-based separator (MTI); size: 16 μm×80 mm (thickness×width)) in pouch-type full cell. For coin-type full-cell equipment, the hard carbon anode was blending with PVDF 10% and acetylene black 10%, then coating on copper foil with average loading density of 4–5 mg/cm2 and dried under vacuum at 80 °C for 12 h. The N/P ratio of coin-type full cell is 1.02–1.1, and for pouch cell is 1.15–1.2. The electrolyte amount are～120 μl. Before pouch cell test, the pouch cell first be charged at a 0.1 C rate to 4.4 V. Subsequently, the pouch cell be unsealed for gas release, resealed properly, and then subjected to cycle testing, with a constant pressure of 69 kPa maintained throughout the entire cycling process. The CE is calculated as the ratio of discharge capacity divided by the charge capacity in the preceding charge cycle. EIS measurements were performed at 25 °C under open circuit potential, using a potentiostatic signal with an amplitude of 5 mV. The tests covered a frequency range of 0.1–106 Hz, with 12 data points collected per decade. The assembled cells were aged for 12 h to ensure that the system was fully wetted. Cyclic voltammetry (CV) results were collected at 0.2 mV s−1 within 2.4–4.5 V from electrochemical workstation (Autolab). Galvanostatic cycling was carried out on a Land BT2000 tester (Wuhan, China). The Na⁺ diffusion coefficient (DNa+) was extracted from GITT data with the equation DNa+ = 4/π τ (mB VB / MA S)2 (ΔEs / ΔEτ)2 (1) where D denotes the diffusion coefficient of Na+ in the cathode, VM (cm3 mol–1) is the molar volume, mB and MB are the electrode mass and molar mass of cathode material. S (1.13 cm2) is the geometric area of the electrode. τ is the duration of the current pulse in the GITT step. ΔEs is the steady state voltage change, and ΔEτ is the total voltage change during the pulse. The charging process (galvanostatic pulse) lasted for 15 min, with a data acquisition interval of 10 s, corresponding to an acquisition frequency of 0.1 Hz (0.1 data points per second).

For the hybrid solid-state batteries, a Na disc was attached to one side of the NZSP pellet and heated to 200 °C in a Swagelok cell. A pressure of about 5000 Pa was applied to ensure intimate contact between molten Na and NZSP. To ensure the intimate contact between the cathode and solid electrolyte, 10 μL of 1 mol L–1 NaPF6 dissolved in EC: PC with 5% FEC was added, and the cathode loading was around 3 mg/cm2. The thickness of a NZSP pellet is 1 mm. The testing temperature is～25±2 °C. Supplementary Data 1 consists of an additional set of electrochemical result, which is provided to demonstrate the reproducibility of the electrochemical result of half-cell and full-cells.

Neutron powder diffraction measurement
The neutron powder diffraction experiment was carried out at the multi-physics instrument, a high-intensity time-of-flight diffractometer, at the China Neutron Spallation Source (CSNS)51. Approximately 1 g of the sample was loaded into a cylindrical sample holder, which was hermetically sealed with indium wires under an argon atmosphere. The sealed holder was then transferred to the scattering chamber for room-temperature testing, covering a neutron wavelength range of 0.1–4.5 Å. Diffraction patterns were collected using six detector banks (Bank 2–7) with 2θ range from 12.5 to 170.0°.

Computational methods
DFT calculations were performed with the Vienna ab-initio simulation package (VASP). The generalized gradient approximation (GGA) with the Perdew-Burke-Ernzerhof (PBE) functional was adopted as the base method. Given that GGA fails to accurately reproduce the localized electronic states of TMO materials, the GGA+U correction was introduced, with U values set to 5.0, 6.9 eV for Mn and Ni, respectively. Core electrons were described by the plane-wave projector augmented wave (PAW) pseudo-potentials. For the DFT calculations, a plane-wave basis set cutoff energy of 520 eV was adopted. A force convergence criterion of less than 0.01 eV Å−1 per unit cell was applied, while an electronic iteration convergence threshold of 1×10–5 eV was set for the self-consistent field calculations. A 2×2×3 supercell of Mn-Ni-ordered and disordered structure was conducted. All models have undergone structural optimization.

The energy barrier of phase transition caused by gliding is obtained by the equation:
ΔE=E(O2)−E(P2) (2)
where E stands for the lowest formation energy of Na0.7Mn0.8Ni0.2O2 at each phase state at each composition. All relevant calculations were based on models with the same numbers and types of atoms.

Calculation of stacking faults probability
According to the Scherrer equation:
β = Kλ / Lcosθ (3)
β = βsize + βf (6)
For isotropic near-spherical nanocrystals: β = βsize + Acosφz (7)
when l is odd number, A = f / 2c (8)
when l is even number, A = 3f / 2c (8)

Solving for D (crystallite size) and f (SF probability) using Eqs. (6)–(8) is complex. Thus, a least squares method is proposed: substitute (3) and (7)-(8) into (6), multiply by cosθ / λ and simplify to:
βcosθ / λ = 0.89 / D + 2cλcosθcosφz4f (9)

Define variables: Y = βcosθ / λ, a = 0.89 / D, X = 1 / ccosθcosφz, F = 4f. Equation (7) becomes: Y = a + FX (10)

The regular matrix equation for least squares is:
[ n ΣX ] [ a ] = [ ΣY ]
[ ΣX ΣX2 ] [ F ] = [ ΣXY ] (11)

Solutions are: D = 0.89 / a (10), F = (n ΣXY - ΣX ΣY) / (n ΣX2 - (ΣX)2) (12)

Williamson–Hall analysis
Variations in crystallite size and microstrain lead to the broadening of the diffraction peaks of the sample’s XRD pattern28–30, which changing differently as a function of the Bragg angle (θ). These variations can be calculated using the following equation:
βstrain = 4εsinθ (13)
βsize = Kλ / Lcosθ (14)
where βstrain and βsize are lattice strain and size of crystallite, respectively. K is the crystallite size (K=0.9), and λ is the wavelength. The following equation are used to correct the instrumental broadening by the standard sample of single-crystal silicon:
βsample = sqrt((βtotal)2 - (βinstrument)2) (15)
where βtotal is the full-width at half-maximum of the sample, βinstrument is the full-width at half-maximum of the standard sample. βsample is the sum of βstrain and βsize, followed by the equation:
βsample = βstrain + βsize = 4εsinθ / cosθ + Kλ / Lcosθ (16)
Multiplying throughout by cosθ gives
βhklcosθ = Kλ / D + 4εsinθ (17)
βhkl is the line broadening (represented as FWHM) of the peak of reflecting (hkl) plane.

Calculation of cationic potential
The cationic potential is introduced by Hu et al. to distinguish and design new materials1, the weighted average ionic potential of TMs can be defined as ΦTM = Σ (ni / Ri) (18)
where wi is the content of TM having oxidation state ni and radius Ri the weighted average ionic potential anionic ion can be defined as: Φanion = Σ (ni / Ri) (19)

The weighted average ionic potential of Na is ΦNa, where x represents the Na content.
ΦNa = x / RNa (20)

Based on the weight average ionic potentials of TM, Na and anionic ions, the cationic potential can be defined as:
Φcation = ΦTM - ΦNa / Φanion (21)

Peer review rounds
[{"round": 1, "reviewer_comments": {"reviewer_1": "In this work, the authors investigated the mechanism of P to O phase transition for NaTMO2 cathode materials, and reported a method to supress the NaTMO2 phase transition by pre-processing the N2TMO2 material and adding O-type stacking faults. Structural characterizations & electrochemical test show that the sythesized SF-NaTMO2 material has minimal P to O phase transition and enables improved cycling stability and rate capability during sodiation-desodiation.\nThe central claims of the study are supported by entensive evidences, and the new SF-NaTMO2 material proposed by this work has the potential to improve the energy density, power density and cycling stability of sodium-ion batteries. Thus, I recommend the paper to be accepted for publication after the following comments being addressed:\n1. For the characterization of the cycled SF-NaMN (Fig.5):\na): What is the cycling condition? Please specify\nb): Why only after 30 cycles? Note that the authors claim that the SF-NaMN could enable stable cycling of SIB for > 600 cycles. Please explain.\n2. Fig.3b: What do the two histogram mean? Are these diffusivity values at 4.5V? Please specify in the caption.\n3. Fig.3c and Fig.3e: Please provide the voltage profiles at different cycle number.\n4. Fig.3j: Please provide the electrolyte amount in the pouch cell.\n5. Fig.S25: Details of the solid-state-battery is needed.\n6. Fig.S10b:\na): The first 3 cycles at 0.2C shows obvious capacity degradation. Please explain the reason.\nb): What is the difference between Fig.S10b and Fig.S15? The authors need to specify the details testing conditions of every single test.", "reviewer_2": "The manuscript falls short in mechanistic demonstration, benchmarking, and reproducibility. Major additional experiments—controlled variation of stacking‐fault density, side-by-side full-cell performance comparisons with state-of-the-art materials, and comprehensive three-dimensional structural analyses—are necessary to substantiate the central claims. Absent these critical data and comparative analyses, the manuscript cannot justify its key claims or its advancement over existing cathode designs.\n1. The central claim—that pre-engineered O-type stacking faults alone decouple slab gliding from abrupt c-axis contraction—is not rigorously demonstrated.\n2. The record energy density (635 Wh kg-1) and full-cell performance are not meaningfully compared against leading P2-type or OP4/Z cathodes. Without direct side-by-side data, the claimed “state-of-the-art” advance is unsubstantiated.\n3. DFT snapshots and in situ XRD show correlation, but causal proof (e.g. controlled removal of faults or systematic fault-density variation) is lacking.\n4. The manuscript attributes phase-transition suppression solely to reduced O2-–O2- repulsion at pre-fault layers, but offers no measurements of local interlayer forces or dynamic gas evolution tied to slab motion. Key questions—such as whether residual P-to-O transitions occur at microscopic scales—remain unanswered.\n5. Although stacking faults are quantified by XRD broadening and PDF analysis, their three-dimensional distribution, size distribution, and potential clustering are not characterized.\n6. The method relies on precise Mn/Ni co-precipitation control; extension to other compositions is claimed but not systematically proven. The cycle life and rate tests at practical loadings are missing.", "reviewer_3": "In this work, the authors present a strategy to inhibit the deleterious effects of phase transition from the P- to O-type crystal structures of layered transition metal oxide cathode materials for Na-ion batteries, and explore the impact of this strategy using a wide variety of physiochemical and electrochemical techniques. Through their work, the authors show that by inducing the formation of Mn-rich layers during the synthesis of a NaMn0.8Ni0.2O2 material, they are able to create introduce a small quantity of O-type slip faults into the material. These faults appear to inhibit further P to O transitions at high voltages, significantly improving the stability of the cathode over extended cycling. Numerous characterization techniques, including STEM, XRD, XPS, ToF-SIMS, XANES, in situ synchrotron XRD, and others are used to provide evidence for this conclusion, with EDX and STEM measurements providing particularly convincing evidence for the formation of these slip faults and their impact on the material stability.\nWhile the authors provide extensive characterization for their materials and thus have a well-researched and compelling story, there are areas in which the brevity of their exposition makes it difficult for the reader to understand how they came to this conclusion. We have addressed several major and minor areas below which we believe the authors should address prior to this paper being suitable for publication in Nature Communications.\nMajor Comments\n1. Please provide a better explanation for/state more clearly how the Williamson-Hall analysis shows that the peak broadening is due to strain rather than particle size. If it were purely due to particle size, would we expect all of the slopes for each material to be the same? Is there a particular significance to the measures slopes themselves beyond, i.e. a slope >0.5% indicates strain, while a slope <0.5 indicates particle size? (pg.7)\n2. The EDX analysis on pg. 12 suggests Mn enrichment in select sheets that drives cationic potential differences. What is the extent of Mn enrichment compared to the generic Mn0.8Ni0.2 stoichiometry? How large is the resulting cationic potential difference, particularly in reference to the numbers provides in Supplemental Table 2?\n3. On pg. 14, the authors’ analysis of the CV sweeps of their cathodes glosses over several peaks that may be important for their story. First, on line 255 it is stated that the 4.2V peak corresponding to the P2-O2 transition is not observed, but this peak does appear to be present, albeit shifted to slightly lower potentials. Likewise, on lines 257-258 they claim to see a single peak in the SF-NaMN at 4.11/4.38V, but this is clearly part of a doublet with the previously mentioned peak. Likewise, the authors imply that the SF-NaMN does not show a peak near 3.78/3.64V (line 263) related to Na+ vacancy ordering, but the peak (albeit diminished) is present. They ultimately conclude that the P-O transition is “missing” (pg.15, line 270). We request that the authors revisit their analysis and document that while the SF-NaMN appears to exhibit significantly reduced P-O transition and vacancy ordering, it does not prevent it. Instead of eliminating the process, the effect of the artificially-introduced slip faults appears to inhibit much of the negative impact of the P-O transition. Please supplement/support your CV analysis with dQ/dV analysis from the battery cycling data, as this may better highlight the additional peaks discussed above.\n4. Related to Comment 3 above, in Figure S29a it appears there is a very slight weakening of the P2 phase around 39.75 deg. at the end of charge that would correspond to the P-O transition. Please comment on this in relationship to the electrochemical evidence for the reduction (but not elimination) of this transition.\n5. On Pg. 24, lines 438-441 the authors argue that they have decoupled slab sliding from lattice contraction; this does not seem accurate. Rather, the evidence appears to suggest that by introducing a small quantity of slip fault defects the authors have been able to prevent additional significant volume change and slab gliding upon desodiation (Figs. S29 and 4g) in the SF-NaMN compared to the standard material. We request the authors either a) provide more evidence/a clearer explanation for why they believe that introducing a limited number of faults represents “decoupling” of the two aspects, or else b) adjust their message to better reflect that the induced faults appear to inhibit the concerted expansion and shifting of slabs from occurring instead of decoupling them.\nMinor Comments\n1. Please provide experimental and instrument details regarding your XPS and ToF-SIMS measurements.\n2. Please ensure all acronyms (e.g., CEI, SEI) are defined immediately upon introduction, and check to confirm that SEI is used to refer to the film atop the anode, while CEI is used to refer to the film atop the cathode.\n3. Please provide a better explanation for how the various equations for calculating the percentage of stacking faults are used. It is currently unclear how any of them lead to a percentage. (pg.9)\n4. On pg.14, please invert the x-axis of your XPS plots in line with standard convention of binding energies increasing to the left.\n5. On pg. 17, please define relaxation time and describe the DRT analysis more fully.\n6. On pg. 18, please provide a citation for the use of pouch cell performance as the “golden criterion”, or else remove the phrase and use a more generic phrase.\n7. On pg. 23, please provide a better description for how the slab energy is related to the propensity of the materials to undergo the P-O transition.\n8. Figure S29 is missing the c-axis data.\n9. On pg. 27, please include if the EDX data was collected using an SEM or STEM, and describe why the data indicates that TM dissolution is reduced in the SF-NaMN sample.\n10. On pg. 27, please include an FFT of the CEI of the SF-NaMN instead of just the bulk particle to provide a more fair comparison to the NaMN sample.\n11. On pg. 27 lines 489-491, please provide more detail as to why the EIS indicates a) superior ionic conductivity and b) increased stability. It is not immediately obvious from the data alone.\n12. On pg. 28 line 496, the authors mention that CO2 is evolved much earlier from SF-NaMN than from NaMN, but do not provide additional discussion. If SF-NaMN releases less CO2 overall, but begins at lower potentials, what is the implication?\n13. Please provide xy scale bars for the ToF-SIMS 3D chemical maps in Fig.5h,j."}, "author_response": "Manuscript Number: NCOMMS-25-34868-T\nTitle: Decoupling TMO Slab Gliding and Lattice Contraction to Enable High-Voltage Na-Ion\nDear Editor and Reviewers,\nThank you very much for arduous work and valuable comments. We have studied the comments\ncarefully and we are pleased to find the constructive comments in the respects of the interest and\nperformance improvement on sodium ion batteries from all reviewers. The main problem raised by\nthe reviewers, and the explanation of some scientific issues, have been fully addressed in our revised\nmanuscript. Here within the enclosure is our revised manuscript, which has been improved\naccording to your suggestions. Revised parts are highlighted with yellow color, and the point to\npoint responses to reviewers’ comments are listed below:\nComments from Reviewer#1:\nReviewer #1: In this work, the authors investigated the mechanism of P to O phase transition for\nNaTMO cathode materials, and reported a method to supress the NaTMO phase transition by pre-\n2 2\nprocessing the N2TMO material and adding O-type stacking faults. Structural characterizations &\n2\nelectrochemical test show that the sythesized SF-NaTMO material has minimal P to O phase\n2\ntransition and enables improved cycling stability and rate capability during sodiation-desodiation.\nThe central claims of the study are supported by extensive evidences, and the new SF-NaTMO\n2\nmaterial proposed by this work has the potential to improve the energy density, power density and\ncycling stability of sodium-ion batteries. Thus, I recommend the paper to be accepted for publication\nafter the following comments being addressed:\nResponse: Thank you for your positive and valuable comments. The following concerns have been\nwell addressed accordingly and the revised manuscript is highly expected to be suitable for\n1\npublishing in Nature Communications.\na): What is the cycling condition? Please specify\nb): Why only after 30 cycles? Note that the authors claim that the SF-NaMN could enable stable\ncycling of SIB for > 600 cycles. Please explain.\nResponse: Thank you for your valuable comments.\na) The SF-NaMN and NaMN electrode in Fig 5 are tested under the galvanostatic charge/discharge\nmode with current density of 0.2 C in 2.4-4.5 V vs. Na+/Na. We have revised in the manuscript and\nhighlighted in yellow.\nIn manuscript page 28, line 518: The SF-NaMN and NaMN electrode are tested with current density\nof 0.2 C in 2.4-4.5 V vs. Na+/Na.\nb) The 600 cycles correspond to the cycling performance of the full cell within 2-4.2 V. For the half\ncell, SF-NaMN retains 70% of its capacity after 300 cycles. We previously took HADDF-STEM\nimages of NaMN after 300 cycles, while the material had completely collapsed, the zone axis could\nnot be aligned, and no atomic phase structure was observable (Fig R1). SF-NaMN cloud maintain\n93% capacity retention after 30 cycles, while NaMN only retains 79% of its initial capacity (Fig\nR2). There is an obvious capacity difference between the SF-NaMN and the NaMN after 30 cycles.\nTherefore, we selected the electrode materials after 30 cycles to observe the evolution of their\nsurface and bulk phase structures.\nFig R2 The cycling performance of SF-NaMN and NaMN at 0.2C within 2.4-4.5 V.\n2. Fig.3b: What do the two histogram mean? Are these diffusivity values at 4.5V? Please specify in\nthe caption.\nResponse: Thank you for your comments. The two diffusion coefficients in Fig 3b represent the\ndiffusion coefficients of SF-NaMN and NaMN at 4.5V, respectively.\nWe have modified and highlighted them in the caption of Fig 3b.\nIn manuscript, page 20, line 367 : b Sodium ion diffusion coefficient (D\nNa\n+) of SF-NaMN and NaMN\ncalculated from GITT formula, the two histograms are the diffusion coefficients of SF-NaMN and\n3. Fig.3c and Fig.3e: Please provide the voltage profiles at different cycle number.\nResponse: Thank you for your comments. We have provided the voltage profiles at different cycles\nin supplementary information, page 20 and page 28.\nFig R3/Supplementary Fig. 18 Charge-discharge curves of SF-NaMN at 2C in 2.4-4.5 V.\nFig R4/Supplementary Fig. 26 The charge and discharge curves for different cycle numbers of\nups ups\n4. Fig.3j: Please provide the electrolyte amount in the pouch cell.\nResponse: Thank you for your comments. The amount of electrolyte used in the pouch battery is 2\nml/2.3g. We have provided this part in supplementary information page 59.\nTable R1/Supplementary Table 10. Pouch cell parameters of HC||SF-NaMN.\nComponent Mass (g)\nCathode SF-NaMN 3.7857\nSuper P 0.12\nPVDF 0.08\nAnode HC 1.537\nSuper P 0.056\nCMC 0.028\nSBR 0.028\nSeparator +Tabs Celgard separator 5.0\n+ Al/Ni tab + Al and Cu foil\nPackage Al plastic film 2.2\nElectrolyte 1 Ah/4ml 2.3\nTotal pouch battery mass Energy (mWh) Energy density\n(Wh kg-1)\n15.1347 1577 104.2\n296\n(Calculated based on total\nmass of active electrode )\n5. Fig.S25: Details of the solid-state-battery is needed.\nResponse: Thank you for your comments.\nWe have added this description in manuscript page 35, line 654. For the hybrid solid-state batteries,\n5\na Na disc was attached to one side of the NZSP pellet and heated to 200°C. A pressure of about\nintimate contact between the cathode and solid electrolyte, 10 μL of 1 mol L-1 NaPF dissolved in\nreason. b): What is the difference between Fig.S10b and Fig.S15? The authors need to specify the\ndetails testing conditions of every single test.\nResponse: Thank you for your comments.\na) The capacity decay in the first three cycles in Fig. S10 is mainly caused by the Jahn-Teller effect\nin the redox reaction of Mn3+/Mn4+ below 2.0 V,1,2 which is a general characteristic of Mn-based\nmaterials, arising from the intrinsic electronic configuration of Mn3+ (d4 high-spin state) which tends\nto induce octahedral distortion, leading to structural distortion and capacity decay (Fig R5). To avoid\nthe Jahn-Teller distortion of Mn3+/Mn4+ in the 1.5-2.4 V, we conducted tests within the 2.4-4.5\nvoltage range, therefore this adjustment effectively avoids the formation of Mn3+ by confining the\nredox activity to higher valence transitions.\nb) Fig. S10 shows the rate performance of SF-NaMN and NaMN at 1.5-4.5V, while Fig. S15\npresents the rate performance of SF-NaMN and NaMN at 2.4-4.5V.\nWe have revised this part in the caption of Fig S12 and Fig S17 in supplementary information, page\n14 and 19.\nSupplementary Fig. 12 a Initial charge/discharge profiles of SF-NaMN and NaMN in the voltage\nrange of 1.5–4.5 V. b Rate performance of SF-NaMN and NaMN within 1.5-4.5 V at room\ntemperature. The electrolyte is 1 mol L−1 NaPF dissolved in EC : PC with 5% FEC. The loading\n6\ndensity of cathodes are 10-12 mg/cm2.\nSupplementary Fig. 17 Rate performance comparison of SF-NaMN and NaMN electrodes in half-\ncell within 2.4-4.5 V at room temperature. The electrolyte is 1 mol L-1 NaPF dissolved in EC : PC\n6\nwith 5% FEC. The loading density of cathodes are 10-12 mg/cm2.\nReviewer #2 (Remarks to the Author):\nThe manuscript falls short in mechanistic demonstration, benchmarking, and reproducibility. Major\nadditional experiments—controlled variation of stacking‐fault density, side-by-side full-cell\nperformance comparisons with state-of-the-art materials, and comprehensive three-dimensional\nstructural analyses—are necessary to substantiate the central claims. Absent these critical data and\n7\ncomparative analyses, the manuscript cannot justify its key claims or its advancement over existing\ncathode designs.\nResponse: Thank you very much for your arduous work and comments. We have studied the\ncomments carefully and the following revision concerns have been well addressed accordingly.\nSpecifically, the responses to this letter are highlighted in blue, while the revised texts in the\nmanuscript and supporting information are marked in yellow.\n1. The central claim—that pre-engineered O-type stacking faults alone decouple slab gliding from\nabrupt c-axis contraction—is not rigorously demonstrated.\nResponse: Thank you for your valuable comments.\nWe believe the discrepancy regarding the term “decouple” may stem from differing understandings\nof its definition. Following the reviewer’s suggestion, we have consulted the Oxford Dictionary,\nwhose core definition of “decouple” is “to end the connection or relationship between two things.”\nSpecifically in our work, we define “decouple” as a scenario whereby stacking faults are pre-\nengineered into the P-phase lattice in the initial structure, ensuring that the interlayer spacing of\nthe P2-phase material does not contract during charge-discharge cycling.\nIn the conventional understanding, TMO slab gliding and abrupt c-axis contraction are intrinsically\n₂\ncoupled phenomena that jointly define the P2-O2 phase transition during deep sodium extraction.\nFor a typical P2-type layered oxide, this usually happens at 4.2V, when 50% Na removal is achieved,\nduring which the TMO slabs glide along [1/3, 2/3, z] or [2/3, 1/3, z] direction.3,4 In the meantime\n2\nthe Na+ coordination environment transforms from prismatic (P-type) to octahedral (O-type),\naccompanied by abrupt c-contraction. Generally, the TMO slabs gliding, or the Na+ coordination\n2\nchange are regarded as the fundamental origin of the P2-O2 phase transition, while c-contraction is\n8\nconsidered as an accompanied structural response.2,5\nHowever, we have demonstrated that the slab gliding is not always coupled with c-contraction. If\ncertain amount of TMO slabs are pre-glided from O to P positions, i.e. O-type stacking faults are\n2\npre-engineered in the initial P-phase lattice, such a structure remains stable at high voltage. No\nabrupt c-axis contraction happens even with 70% Na removed (4.5V). Therefore, distinguished from\ntraditional thought, the slab gliding is not necessarily always accompanied with c-axis\ncontraction.\nTherefore, we use the word ‘decouple’ refers to breaking the obligatory co-occurrence between\nTMO slab gliding and abrupt c-axis contraction.\ncycling is significantly smaller than that of NaMN, and most importantly, no abrupt c-axis\ncontraction is observed despite the occurrence of slab gliding. This outcome directly demonstrates\nthat we have successfully decoupled transition metal gliding from the c-axis contraction that is\ntraditionally considered inevitable in the P-O phase transition. By pre-constructing O-type stacking\nfaults in the P2 layered cathode, we effectively suppress both excessive volume expansion and the\nabrupt c-axis contraction during charge-discharge processes, validating our definition of \"decouple\".\nTo help readers better understand our concept of decoupling, we have added these sentences in the\nmain text. In manuscript, page 5 , line 104:\nThe \"decouple\" here does not refer to the traditional scenario where the gliding between transition\nmetal layers during the charging and discharging process is necessarily accompanied by cell\ncontraction. Instead, our definition of \"decouple\" refers to the scenario where stacking faults are\npre-engineered into the lattice of the P-phase in the initial structure, the interlayer spacing of the P2-\nphase material will not contract during the charge-discharge process.\nTherefore, although 15.8% TMO slabs have been pre-glided from P to O position, no significant\n2\nc-contraction was observed in SF-NaMN. In this sense, the traditional correlation between TMO\n2\nslab gliding and c contraction has been decoupled.\n2. The record energy density (635 Wh kg-1) and full-cell performance are not meaningfully\ncompared against leading P2-type or OP4/Z cathodes. Without direct side-by-side data, the claimed\n“state-of-the-art” advance is unsubstantiated.\nResponse: Thank you very much for your valuable comment, which has helped us further improve\nthe rigor of our work. We fully understand your concern regarding the comparison of our material’s\nperformance. Initially, we chose to compare our cathode with previously reported high-voltage\ncathode materials (operating voltage ≥ 4.2 V) published in internationally leading journals. This\napproach is consistent with the common practice in most relevant literatures6-12, as it ensures\nobjectivity and reliability by focusing on a homogeneous category of materials with comparable\ncore characteristics (i.e., high-voltage operation). Through a comparison of recent literature, for P2-\ntype layered cathode materials operating in a wide voltage range, the energy density of our half-cell\n(635 Wh kg-1) was slightly lower than P2-type Li-Mn based cathodes, where Li+ stimulated more\nanion (O2-) to participate redox reaction. Accordingly, we have revised the relevant section in the\nabstract of this manuscript. However, although some P2-type cathodes exhibit high specific capacity,\nthey suffer from sluggish Na+ transport kinetic at high voltages, this consequently indicates\nadditional sodium sources (sacrificial cathode presodiation agents, presodiated anode etc), are still\nrequired to enhance their energy density (> 250 Wh kg-1) in full cells. In this work, we have\nsignificantly enhanced the Na+ diffusion kinetics of Na Mn Ni O at high-voltage (>4.2 V) by\n0.7 0.8 0.2 2\n10\npre-introducing stacking faults in P-type structure, enabling the full cell to be directly assembled\nwith unpresodiated hard carbon and achieving a high energy density of 297 Wh kg-1 (without\npresodiation, calculated based on the active masses of both the cathode and anode). We have also\ncompiled a table, in supplementary information page 58, to compare the energy density, operating\nvoltage, and cycling stability of the full cell we assembled with those of the materials reported in\nthe literature. Accordingly, we have revised the description in the abstract of revised manuscript:\n“Consequently, the full cell assembled with unpresodiated hard carbon demonstrates an exceptional\nenergy density of 297 Wh kg-1 (based on the total active mass of electrodes), outperforming the\nstate-of-the-art literature.”\nRegarding your point that “direct side-by-side data is lacking,” it is challenging to directly reproduce\nboth the materials and the reported performance described in much of the literature due to\nvariations in synthesis protocols, testing setups, and other experimental details. However, we\nhave now synthesized several representative P2-type and O3 cathodes as suggested several\nsamples [Na Zn ]Ni Cu Mn O , Na Ni Mg Al Mn Ti O ,\n0.67 0.05 0.18 0.1 0.67 2 0.67 0.28 0.04 0.02 0.64 0.02 2\nNaNi Cu Fe Mn Ti O , Na Li Cu Ni Fe Mn O F , (Fig R6a-e) and\n0.3 0.1 0.2 0.3 0.1 2 0.89 0.05 0.11 0.11 0.3 0.43 1.97 0.03\nconducted comparative tests under the same experimental conditions (including electrode\npreparation, electrolyte composition, testing potential window, and charge-discharge rate) to\neliminate systematic discrepancies.\nBy comparing the charge-discharge curves of all full cells, the average operating voltage of the\nHC||SF-NaMN is 3.51 V, which is significantly higher than that of other full cells (Fig R6f-j). This\nexplains why, despite SF-NaMN having a lower specific capacity than O3-phase samples, it exhibits\na higher specific energy density. In addition, SF-NaMN also exhibits the good capacity retention\n11\nrate of 87% after 100 cycles, which is higher than that of other materials, indicating that it is suitable\nfor cycling under high voltage. As a result, the energy density of the SF-NaMN full cell is increased\nto 297 Wh/kg, thus addressing a key limitation of P2-phase materials for the necessity of assembling\nwith pre-sodiated hard carbon anodes to improve energy density of full-cell.\nWe have supplemented this set of direct comparison data in supplementary information page 28,\nline 149, note 4, thereby providing solid support for the claimed \"state-of-the-art\" advance.\nFig R6/Supplementary Fig. 28 a-e The XRD patterns of SF-NaMN,\n[Na Zn ]Ni Cu Mn O Na Ni Mg Al Mn Ti O ,\n0.67 0.05 0.18 0.1 0.67 2 0.67 0.28 0.04 0.02 0.64 0.02 2\nNaNi Cu Fe Mn Ti O , and Na Li Cu Ni Fe Mn O F , f-j The charge and\n0.3 0.1 0.2 0.3 0.1 2 0.89 0.05 0.11 0.11 0.3 0.43 1.97 0.03\ndischarge curves of full cell of SF-NaMN, [Na Zn ]Ni Cu Mn O\n0.67 0.05 0.18 0.1 0.67 2\nNa Ni Mg Al Mn Ti O , NaNi Cu Fe Mn Ti O , and\n0.67 0.28 0.04 0.02 0.64 0.02 2 0.3 0.1 0.2 0.3 0.1 2\nNa Li Cu Ni Fe Mn O F , k The cycling performance of SF-NaMN,\n0.89 0.05 0.11 0.11 0.3 0.43 1.97 0.03\n[Na Zn ]Ni Cu Mn O Na Ni Mg Al Mn Ti O ,\n0.67 0.05 0.18 0.1 0.67 2 0.67 0.28 0.04 0.02 0.64 0.02 2\nNaNi Cu Fe Mn Ti O , and Na Li Cu Ni Fe Mn O F within 1-4.4 V.\n0.3 0.1 0.2 0.3 0.1 2 0.89 0.05 0.11 0.11 0.3 0.43 1.97 0.03\nperformance\nIn order to maintain the consistency of the testing voltage range, type and amount of electrolyte,\nand experimental details, thereby enabling, we synthesized four representative cathode materials\n([Na Zn ]Ni Cu Mn O Na Ni Mg Al Mn Ti O ,\n0.67 0.05 0.18 0.1 0.67 2 0.67 0.28 0.04 0.02 0.64 0.02 2\nNaNi Cu Fe Mn Ti O , Na Li Cu Ni Fe Mn O F ) reported in the literature\n0.3 0.1 0.2 0.3 0.1 2 0.89 0.05 0.11 0.11 0.3 0.43 1.97 0.03\nto assemble full cells (Supplementary Fig. 28). The operating voltage of SF-NaMN is significantly\nhigher than that of other materials. Moreover, its capacity retention is also enhanced during cycling\nwithin the voltage range of 1-4.4 V at 1C after 100 cycles.\nthe recently published works with the P2- and O3-type cathodes.\nCathode Anode Voltage Capacity Energy Average Ref.\nRange retention density Voltage\n(V) (%) (Wh (V)\nkg-1)a\nP2-Na Mn Ni O HC 1-4.4 87.3%, 100 297 3.5 This\n0.7 0.8 0.2 2\ncycles, 1C work\n2-4.2 77.3%, 600\ncycles, 1C\nP2- Presodiated 1.4-3.9 91.3%, 200 274 2.3 10\nNa\n0.72\nMg\n0.1\nMn\n0.9\nO\nP2- Presodiated 1.9-4.3 93%, 150 292.3 3.41 6\nNa Cu Li Mn HC cycles, 1C\n0.8 0.22 0.08 0.67\nO\nNa Ni Mg Al HC cycles, 0.2C\n0.67 0.28 0.04 0.0\nMn Ti O\n2 0.64 0.02 2\nP2- Presodiated 2.0-4.3 82.6%, 100 228.4 2.78 8\nNa\n0.8\nCu\n0.22\nLi\n0.08\nMn\n0.67\nHC cycles, 0.2C\nO\nNa Zn ]Ni Cu HC cycles, 1C\n0.67 0.05 0.18 0.\nMn O\n1 0.67 2\nP2- Presodiated 1.4-4.2 92.4%, 100 273 3.31 14\nNa Ca [Li Ni HC cycles,\n0.61 0.05 0.1 0.23\nP2-type Mn ]O F ~0.5C\n0.67 1.95 0.05\ncathode P2- Presodiated 1-4.4 ~80%, 50 286.8 3.17 15\nNa Li Mg Ni HC cycles,\n0.8 0.03 0.05 0.28\nFe Mn Ti O ~0.5C\n0.05 0.51 0.05 2\nP2- Presodiated 2.3-4.14 84%, 300 202 3.5 16\nNa Mn Ni Nb HC cycles, 5C\n0.78 0.67 0.31 0.\nO\n02 2\nP2-Na Ca [Ni Presodiated 1.9-4.2 80.4%, 50 257.6 3.25 17\n0.76 0.05 0.23\n□ Mn ]O HC cycles, 0.5C\n0.08 0.69 2\nP2- Presodiated 1-4.4 85.1%, 60 296 3.3 18\nNa Li Mg Ni HC cycles, 0.2C\n0.7 0.03 0.03 0.27\nMn Ti O\n0.6 0.07 2\nP2- HC 1.5-4.3 / 270 2.54 19\nNa Li Fe Mn\n0.67 0.1 0.37 0.53\nO\n0.7 0.1 0.75\nO\n0.15 2\nP2-Na MnO -coated Presodiated 1.5–3.9 83.6%, 200 274.8 2.55 21\n2/3 2\nP2/O3- O3-NaNi Mn O HC cycles, 3C\n0.5 0.5 2\ntype P2/O3- Presodiated 0.5–4.4 130 294.6 3.18 22\ncathode Na Ni Mn Ti HC\n0.85 0.34 0.33 0.3\nO\n3 2\nO3- Presodiated 1.0-4.1 70.4%, 500 255 2.82 23\nNaNi\n0.6\nCo\n0.2\nMn\n0.2\nO\nO3- HC 0.5-4.0 80%, 682 269 ~3.1 24\nNaNi Cu Fe Mn cycles, 1C\n0.3 0.1 0.2 0.3\nTi ]O\n0.1 2\nO3- HC 1.0 4.0 81.6%, 200 268 3.2 25\nO3-type\nNa Li Cu Ni cycles, 1C\n0.95 0.07 0.11 0.11\ncathode Fe −\nMn O F\n0.3 0.41 1.97 0.03\nO3- HC 0.5-4.4 80%, 700 260 3.2 12\nNa Li Ni Mn cycles, 3C\n5/6 2/27 8/27 11/2\nTi O\n7 6/27 2\n14\n3. DFT snapshots and in situ XRD show correlation, but causal proof (e.g. controlled removal of\nfaults or systematic fault-density variation) is lacking.\nResponse: Thank you for your valuable comments.\na) Following your recommendations, we have built models with stacking fault probability of 8.33%\nand 12.5%, respectively (Fig R7). These models were constructed to investigate how the controlled\nremoval of faults affects the c-axis variation of the structure during the desodiation process.\nWhen desodiated from Na to Na , for the 12.5% stacking fault model, the c-axis increased by\n0.7 0.2\n0.925% (Fig R7c). Besides, the 8.33% stacking fault model exhibited a more significant c-axis\nexpansion of 1.26% under the same desodiation (Fig R5d). For reference, the original model in the\npaper featured a 16.7% stacking fault content, which only showed a minimal c-axis variation of\n0.097% during desodiation. These results indicate that as the stacking fault density decreases, the c-\naxis variation of the material during desodiation gradually increases, suggesting a progressive\nweakening of the stacking faults’ stabilizing effect on the interlayer spacing (Fig R8a). Moreover,\nthe gliding energy of the TMO slabs in the 12.5% and 8.33% stacking fault models are 0.08\n2\nand -0.279 eV, respectively, which is lower than that of 0.249 eV in the 16.7% stacking fault model\n(Fig R8b). This finding implies that stacking faults with lower densities are not effective in\nsignificantly suppressing interlayer gliding.\nb) Following your suggestions, we have made efforts to synthesize NaMN with different amount of\nstacking faults. However, the amount of the pre-engineered faults are quite trcky, which demands a\ncarful modulation of the synthesis parameters including stirring speed, the temperature and humidity\nof the air etc. which together determine the deposition rate of TM ions and the stacking fault density.\nIncorrect co-precipitation parameters will result impurities in the layered oxides (Fig R9). After\n15\nseveral trials, a sample with 12.1%. stacking faults was successfully achieved (Fig R10a-b).\nCompared to the NaMN material, the synthesized sample exhibited a noticeable slope change in\ntheir electrochemical profiles at the high voltage plateau of 4.2 V (Fig R10c). A clear trend was\nfurther observed: as stacking fault density increased, the high-voltage plateau gradually diminished.\nThis phenomenon likely corresponds to a reduction in the extent of the P-O phase transition. In\naddition, these materials demonstrated varying cyclic performance at a current rate of 0.2C.\nSpecifically, the SF-NaMN sample retained 93% of its initial capacity, while the samples with 12.1%\nstacking fault densities maintained 86% of their specific capacities after 30 cycles, respectively (Fig\nR10d). All these results are significantly higher than the capacity retention rate of 79% observed for\nNaMN.\nHowever a quantitative relationship between stacking faults and the high-voltage performance\ncannot be obtained in the current stage, which can be further investigated in future studies.\nWe have supplemented this description in manuscript, page 24, line 453.\nWe also constructed the models with stacking faults of 8.33% and 12.5% and demonstrated that as\nstacking fault density decreases, the of c-axis variation of the structure during desodiation gradually\nincreases (Supplementary Fig. 39-40, Supplementary note 7).\nWe have supplemented this description in manuscript, page 25, line 468.\nMoreover, we successfully achieved a cathode with 12.1% stacking faults by increasing the\nprecipitation pH of the precursors to 9.62. The synthesized sample exhibited a noticeable slope\nchange in their electrochemical profiles at the high voltage plateau of 4.2 V, and the capacity\nretention was obviously improved compared with NaMN (Supplementary Fig. 41, Supplementary\nnote 8). Therefore, pre-constructing stacking faults in P2-type lattice structure is an effective\n16\nstrategy to enhance the cycle life of the cathode material.\nIn supplementary information, page 67.\nSupplementary Note 7. DFT calculation results of models with different stacking faults.\nTo gain deep insights into how stacking faults influence interlayer spacing changes during the\ndesodiation process, we further constructed models with stacking fault densities of 8.33% and\n12.5%, respectively (Supplementary Fig. 39). DFT results showed that for the model with 12.5%\nstacking faults, the c-axis expanded by 0.925% when desodiated from Na to Na . While the\n0.7 0.2\nmodel with 8.33% stacking faults exhibited a more obviously c-axis increase of 1.26% under the\nsame desodiation conditions. These findings indicate a clear trend: as stacking fault density\ndecreases, the of c-axis variation in the material during desodiation gradually increases\n(Supplementary Fig. 40a). Moreover, the gliding energy of the TMO slabs in the 12.5% and 8.33%\n2\nstacking fault models are 0.08 and -0.279 eV, respectively, were lower than that in the 16.7%\nstacking fault model. This suggests that, the ability of stacking faults to suppress TMO slabs gliding\n2\nis enhanced as stacking fault density rises (Supplementary Fig. 40b).\nIn supplementary information, page 68.\nSupplementary Note 8. Synthesis and electrochemical properties of cathodes with different\nstacking fault densities.\nTo verify the DFT calculations results, and the relationship between the density of pre-engineered\nstacking faults (SFs) and the P-O phase transition, a cathode material with different full widths at\nhalf maximum (FWHM) for the (10l) diffraction peak was achieved (Supplementary Fig. 41). This\nwas accomplished by increasing the pH value of the transition metal precursor solution to the 9.62\nduring the decomposition process. Through calculations, the stacking fault densities of this material\n17\nis determined to be 12.1%, respectively. Compared to the NaMN, the synthesized sample exhibited\na noticeable slope change in their electrochemical profiles at the high voltage plateau of 4.2 V\n(Supplementary Fig. 41c). A clear trend was further observed: as stacking fault density increased,\nthe high-voltage plateau gradually diminished. This phenomenon likely corresponds to a reduction\nin the extent of the P-O phase transition. In addition, these materials demonstrated varying cyclic\nperformance at a current rate of 0.2C. Specifically, the SF-NaMN sample retained 93% of its initial\ncapacity, while the sample with 12.1% stacking fault densities maintained 86% of their specific\ncapacities, respectively (Supplementary Fig. 41d). All these results are significantly higher than the\ncapacity retention rate of 79% observed for the reference sample. These results collectively confirm\nthat the pre-construction of stacking faults in P2-type lattice structure is an effective strategy to\nenhance the cycle life of the cathode material.\nstacking faults and NaMN from Na Mn Ni O (Na Mn Ni O ) to Na Mn Ni O\n28 32 8 80 0.7 0.8 0.2 2 8 32 8 80\n(Na Mn Ni O ), Na Mn Ni O (Na Mn Ni O ) to Na Mn Ni O\n0.3 0.8 0.2 2 42 48 12 120 0.7 0.8 0.2 2 12 48 12 120\n(Na Mn Ni O ).\n0.3 0.8 0.2 2\nFig R8/Supplementary Fig. 40 a The c-axis variation of SF-NaMN with different stacking faults\nprobability during desodiation process. b Theoretical calculation of TM gliding energy of optimized\nSF-NaMN structure with different stacking faults probability with the stoichiometry of\nNa Mn Ni O .\n0.2 0.8 0.2 2\nFig R9 The XRD results of SF-NaMN with impurity phase when the precursors pH is 9.25 and 9.51.\nstacking faults probability. c The charge and discharge curves of NaMN and SF-NaMN with\ndifferent stacking faults probability d The cycling performance of NaMN and SF-NaMN with\ndifferent stacking faults probability.\n4. The manuscript attributes phase-transition suppression solely to reduced O2-–O2- repulsion at\npre-fault layers, but offers no measurements of local interlayer forces or dynamic gas evolution tied\nto slab motion. Key questions—such as whether residual P-to-O transitions occur at microscopic\nscales—remain unanswered.\nResponse: Thank you for your valuable comments. To elucidate the fundamental driving forces\nbehind the P-O phase transition, we focus on the evolution of interlayer oxygen-oxygen (O2--O2-)\nrepulsive forces as the primary governing mechanism. We couldn’t have measured the local\ninterlayer forces, because it remains experimentally challenging.\nstrain evolution in both SF-NaMN and NaMN when charged to 4.5V. We observed that for both the\nuncycled SF-NaMN and the NaMN, interlayer strain was negligible due to their large interlayer\nspacing (Fig R11). Critically, SF-NaMN exhibits significantly attenuated interlayer force\nchanges, as evidenced by its uniformly distributed weak structural strain. In stark contrast,\npristine NaMN shows severe strain localization with disordered strain vectors (Fig\nR12/Supplementary Fig. 36). This distinct difference confirms that the pre-engineering stacking\nfaults drastically suppresses dynamic interlayer force variations during cycling. Consequently, the\nchaotic strain in NaMN directly correlates with uncoordinated gliding and contraction of TMO\n2\nslabs-a hallmark of P-O phase transition-while SF-NaMN maintains structural stability.\nFurthermore, to intuitively monitor the influence of high-voltage electrochemical behavior of SF-\nNaMN and NaMN on gas evolution in batteries, in situ differential electrochemical mass\nspectrometry (DEMS) were employed during the first charging process (Fig R13/Supplementary\nFig. 47). A critical divergence emerged: Carbon dioxide (CO ) release in SF-NaMN onset at 4.331\nexhibited detectable O emission. However, SF-NaMN generated only 308 μmol g-1 of CO ,merely\n2 2\n33% of NaMN’s 937 μmol g-1. According to the previous report, CO are generated from the reaction\n2\nbetween the carbonate electrolyte with the highly active oxygen species on the surface or O evolved\n2\nfrom the lattice oxygen of cathode oxide.26-29 Hence, the amount of CO is highly correlated with\n2\nthe release of O . Therefore, pre-engineering stacking faults in SF-NaMN suppresses oxygen\n2\nreactivity, the delayed CO onset voltage (ΔV = +160 mV) and drastically reduced CO emission\n2 2\n(↓67%) in SF-NaMN collectively indicate attenuated lattice oxygen activation. This aligns with\n21\nstabilized structure of SF-NaMN during high-voltage charging.\nRegarding the question, whether residual P-to-O transitions occur at microscopic scales, it has been\nbe investigated through in-situ XRD analysis. In situ XRD data reveal that the (002) peak of SF-\nNaMN show negligible shift toward lower angles when charged to 4.5 V (Fig. 5c), indicating no\nmacroscopic contraction of interlayer spacing, and no typical O2-type characteristic peaks (e.g. 2-\ntheta at 20o) are observed. This is further corroborated by atomic-scale HAADF-STEM imaging\n(Fig. 4a), where the transition metal interlayer distance in SF-NaMN at 4.5V remains identical to\nits pre-charged state (measured d-spacing: 5.7 ± 0.1 Å vs. 5.6 ± 0.1 Å). Critically, the combined\nabsence of interlayer contraction at both macro- and atomic scales conclusively demonstrates the\nsuppression of P-O phase transitions in SF-NaMN.\nWe have revised in manuscript, page 22, line 406:\n“Moreover, Geometric phase analysis (GPA) SF-NaMN exhibits significantly attenuated interlayer\nforce changes, as evidenced by its uniformly distributed weak structural strain (Supplementary Fig.\n36). In stark contrast, pristine NaMN shows severe strain localization with disordered strain vectors,\nwhich correlates with uncoordinated gliding and contraction of TMO slabs.”\n2\nc The GPA analysis of SF-NaMN from a. d The GPA analysis of SF-NaMN from b.\nFig R12/Supplementary Fig. 36. STEM-HAADF images of a SF-NaMN and c NaMN after 10\ncycles and stopped at 4.5 V. b The GPA analysis of SF-NaMN from a. d The GPA analysis of SF-\nNaMN from b.\ndischarge process with cutoff of 4.5 V.\n5. Although stacking faults are quantified by XRD broadening and PDF analysis, their three-\ndimensional distribution, size distribution, and potential clustering are not characterized.\nResponse: Thank you for your valuable comments. A stacking fault is a crystallographic defect in\nwhich the regular sequence of atomic planes in a crystal lattice is interrupted. It occurs when the\nnormal stacking order of close-packed planes (e.g., ABCABC in FCC or ABAB in HCP structures)\nis locally disrupted, creating a region with a different stacking sequence.30-33 In our study, stacking\nfaults are defined relative to the pristine P2 phase, pertaining to the relative displacement between\nadjacent layers. Three-dimensional reconstruction and nano-CT cannot accurately observe the\n24\natomic-scale distribution of stacking faults. Therefore, we supplemented atomic HAADF-STEM\nimages along the [100] and [001] zone axes, enabling clear observation of interlayer changes, sizes\nand distribution of stacking faults, and potential clustering from 3D perspective. From Fig. R14a,\nit can be observed that the O-type stacking faults are randomly distributed in the P-phase lattice\nwithout obvious clustering. It can be seen from Fig. 2a that the interlayer spacing at the stacking\nfault is 0.55 nm, so there is no significant reduction in the interlayer spacing at the stacking fault\nsites. In addition, from the zone axis along the [001] direction (Fig R14 d-e), it can also be observed\nthat the transition metal layer at the stacking fault has relatively gliding by [1/3, 2/3, 0] fault vector,\nwhich is consistent with our simulation results (Fig R14c, Fig R14f). Therefore, we have\ndemonstrated from a three-dimensional perspective that the stacking faults in SF-NaMN do not form\nclusters to generate a continuous O-phase structure, but are randomly distributed within the P-type\nlattice. This is also consistent with the XRD pattern, where only the broadening of the (10l) peak\noccurs without the generation of new characteristic peaks of the O-phase. We have supplemented\nFig R14a and d in supplementary information, page 10.\nalong the [100] zone axis. c Corresponding simulated model of NaMN and SF-NaMN along the\n[100] zone axis. d-e STEM-HAADF images of SF-NaMN and NaMN projected along the [001]\nzone axis. f Corresponding simulated model of NaMN and SF-NaMN along the [001] zone axis.\n6. The method relies on precise Mn/Ni co-precipitation control; extension to other compositions is\nclaimed but not systematically proven. The cycle life and rate tests at practical loadings are missing.\nResponse: Thank you for your suggestion. In the current stage, we have extended this method to\nother composition (Na Mn Ni Li O , Na Mn Ni Ti O , Na Mn Ni Mg O ,\n0.74 0.75 0.2 0.05 2 0.64 0.77 0.18 0.05 2 0.7 0.75 0.2 0.05 2\nNa Mn Ni Cu O , Na Mn Ni Al O , Na Mn Ni Fe O ,\n0.67 0.75 0.21 0.04 2 0.67 0.75 0.2 0.05 2 0.69 0.74 0.19 0.07 2\nNa Mn Ni Zn O ). We performed systematical charge-discharge profile tests, cycling tests\n0.73 0.76 0.18 0.06 2\nand rate capability measurements of above electrodes with mass loading of are 10-12 mg/cm2, and\nthe results confirmed that all the controlled samples exhibit obviously smoother charge-discharge\ncurves compared with uncontrolled samples (Fig R15-17/Supplementary Fig 20-22, in\nsupplementary information, page 22-24.).\nmaximum (FWHM) for the (10l) diffraction peak (Fig R10a-b) investigated the electrochemical\nproperties of the material. This was accomplished by adjusting the pH value of the transition metal\nprecursor solution to 9.62 during the decomposition process.\nThrough calculations, the stacking fault densities of this material were determined to be 12.1%,\nrespectively. Compared to the NaMN material, both synthesized samples exhibited a noticeable\nslope change in their electrochemical profiles at the high voltage plateau of 4.2 V (Fig R10c). A\nclear trend was further observed: as stacking fault density increased, the high-voltage plateau\ngradually diminished. This phenomenon likely corresponds to a reduction in the extent of the P-O\nphase transition. In addition, the materials demonstrated varying cyclic performance at a current\nrate of 0.2C. Specifically, the SF-NaMN sample retained 93% of its initial capacity, while the\nsamples with 12.1% stacking fault densities maintained 86% of their specific capacities,\nrespectively (Fig R10d). All these results are significantly higher than the capacity retention rate of\n79% observed for the reference sample. These results collectively confirm that the pre-construction\nof stacking faults in P2-type lattice structure is an effective strategy to enhance the cycle life of the\ncathode material.\nThe following results are presented in manuscript page 15 line 282, supplementary information,\nsupplementary Fig 20-22\nThe GCDs of Li, Ti, Mg, Zn, Cu, Al and Fe doped SF-NaMN cathodes show smoother charge-\ndischarge curves in the high voltage range of 4.2-4.5 V, distinguishing from their corresponding\ncontrol samples. Meanwhile, the rate performance, coulombic efficiency and cycling stability of\nthese cathode oxides are further enhanced (Supplementary Fig. 22).\nLi/Ti/Mg/Zn/Cu/Al/Fe-NaMN material and Li/Ti/Mg/Zn/Cu/Al/Fe-NaMN. a, SF-NaMNLi and\n0.05\nNaMNLi ; b, SF-NaMNCu and NaMNCu ; c, SF-NaMNMg and NaMNMg ; d, SF-\n0.05 0.04 0.04 0.05 0.05\nNaMNAl and NaMNAl ; e, SF-NaMNTi and NaMNTi ; f, SF-NaMNFe and\n0.05 0.05 0.05 0.05 0.07\nNaMNFe\n0.07\n; g, SF-NaMNZn\n0.06\nand NaMNZn\n0.06\n.\n0.05\nNaMNLi . c-d Electrochemical performance of SF-NaMNMg and NaMNMg . e-f\n0.05 0.05 0.05\nElectrochemical performance of SF-NaMNTi and NaMNTi . g-h Electrochemical\n0.05 0.05\nperformance of SF-NaMNAl and NaMNAl . i-j Electrochemical performance of SF-\n0.05 0.05\nNaMNCu and NaMNCu . k-l Electrochemical performance of SF-NaMNFe and\n0.04 0.04 0.07\nNaMNFe . m-n Electrochemical performance of SF-NaMNZn and NaMNZn .\n0.07 0.06 0.06\nFig R17/Supplementary Fig. 22 Electrochemical performance comparisons of the layered oxides\nafter doping with Li, Ti, Mg, Zn, Cu, Al and Fe in SF-NaMN and NaMN.\nIn this work, the authors present a strategy to inhibit the deleterious effects of phase transition from\nthe P- to O-type crystal structures of layered transition metal oxide cathode materials for Na-ion\nbatteries, and explore the impact of this strategy using a wide variety of physiochemical and\nelectrochemical techniques. Through their work, the authors show that by inducing the formation\nof Mn-rich layers during the synthesis of a NaMn0.8Ni0.2O2 material, they are able to create\nintroduce a small quantity of O-type slip faults into the material. These faults appear to inhibit\nfurther P to O transitions at high voltages, significantly improving the stability of the cathode over\nextended cycling. Numerous characterization techniques, including STEM, XRD, XPS, ToF-SIMS,\nXANES, in situ synchrotron XRD, and others are used to provide evidence for this conclusion, with\nEDX and STEM measurements providing particularly convincing evidence for the formation of\nthese slip faults and their impact on the material stability.\nWhile the authors provide extensive characterization for their materials and thus have a well-\nresearched and compelling story, there are areas in which the brevity of their exposition makes it\ndifficult for the reader to understand how they came to this conclusion. We have addressed several\nmajor and minor areas below which we believe the authors should address prior to this paper being\nsuitable for publication in Nature Communications.\nResponse: We sincerely thank the reviewer for their thoughtful and positive assessment of our work,\nas well as for your constructive comments. In response to the points raised regarding the brevity of\ncertain explanations, we have carefully addressed each of the major and minor concerns in the point-\nby-point response below. We have provided additional experimental details, clarifications, and\nexpanded discussions that the logical flow and evidence supporting our conclusions are clearly\n30\narticulated. We believe these revisions substantially enhance the manuscript and hope that the\nresponses satisfactorily resolve all the issues raised.\nMajor Comments\n1. Please provide a better explanation for/state more clearly how the Williamson-Hall analysis\nshows that the peak broadening is due to strain rather than particle size. If it were purely due to\nparticle size, would we expect all of the slopes for each material to be the same? Is there a particular\nsignificance to the measures slopes themselves beyond, i.e. a slope >0.5% indicates strain, while a\nslope <0.5 indicates particle size? (pg.7)\nResponse: Thank you for your valuable comments. The Williamson-Hall (W-H) method is a\npowerful tool in X-ray diffraction (XRD) to separate peak broadening effects caused by crystallite\nsize and lattice strain in polycrystalline materials. Crystallite size broadening is determined by the\nScherrer equation: where is the broadening due to size, is a shape fact (0.9),\n𝐾𝐾λ\n𝛽𝛽𝑠𝑠𝑠𝑠𝑠𝑠𝑠𝑠 = 𝐿𝐿𝐿𝐿𝐿𝐿𝑠𝑠𝐿𝐿 𝛽𝛽𝑠𝑠𝑠𝑠𝑠𝑠𝑠𝑠 𝐾𝐾\nis the X-ray wavelength, L is crystallite dimensions and is Bragg angle. Peak broadening due to\n𝜆𝜆 𝜃𝜃\nlattice strain is determined by the equation: , 34 where is strain on various crystal\n𝑠𝑠𝑠𝑠𝑠𝑠𝐿𝐿\n𝛽𝛽𝑠𝑠𝑠𝑠𝑠𝑠𝑠𝑠𝑠𝑠𝑠𝑠 =4ε𝐿𝐿𝐿𝐿𝑠𝑠𝐿𝐿 𝜀𝜀\nplanes (hkl). Hence, the total broadening is determined by β =β + β = +\nsample strain size\n𝑠𝑠𝑠𝑠𝑠𝑠𝐿𝐿 𝐾𝐾𝐾𝐾\n4𝜀𝜀𝐿𝐿𝐿𝐿𝑠𝑠𝐿𝐿 𝐿𝐿𝐿𝐿𝐿𝐿𝑠𝑠𝐿𝐿\nTranslate the formula into a linear form: y = β = c + m·x, where c = , m = 4ε, x = ,\nsample\n𝐾𝐾𝐾𝐾 𝑠𝑠𝑠𝑠𝑠𝑠𝐿𝐿\n𝐿𝐿𝐿𝐿𝐿𝐿𝑠𝑠𝐿𝐿 𝐿𝐿𝐿𝐿𝑠𝑠𝐿𝐿\nIf peak broadening is entirely due to particle size broadening, all characteristic peaks will exhibit\nbroadening. There is no strain component ε = 0. The W-H plot will have a zero slope m = 0, y = c.\nThe connecting lines between various crystal planes should be parallel to the x-axis. If peak\nbroadening is induced by pure strain broadening (no size effect), the W-H plot will have a non-zero\nslope (ε ≠ 0). And ε is proportional to the slope. The greater the strain, the larger the absolute value\nof the slope. Therefore, for SF-NaMN, in the Williamson-Hall equation plot, the stresses on the\n31\n(102), (104), (106) crystal planes can be connected as a straight line passing through the origin, with\na slope of 1.51% and c = 0. This indicates that the peak broadening of these crystal planes is entirely\ncaused by stress strain, not size effect. In contrast, the straight line formed by connecting all crystal\nplanes of NaMN has a slope of 0.21% and an intercept of 0 on the y-axis, which suggests that the\ncrystal plane of NaMN has slight stress strain. However, compared with SF-NaMN, the influence\nof stress on the broadening of crystal planes of NaMN is significantly reduced.\nWe have added this description in Supplementary information, page 61, Supplementary Note 2.\n2. The EDX analysis on pg. 12 suggests Mn enrichment in select sheets that drives cationic potential\ndifferences. What is the extent of Mn enrichment compared to the generic Mn0.8Ni0.2\nstoichiometry? How large is the resulting cationic potential difference, particularly in reference to\nthe numbers provides in Supplemental Table 2?\nResponse: Thank you for your valuable comments. By comparing the counts of EDX spectra, a\nrelatively accurate value was obtained. The ratios of transition metals at the two stacking faults\nregion are approximately Mn:Ni = 0.815:0.185 and 0.789:0.211; as well as 0.789:0.211 and\n0.805:0.195, respectively (Fig R18). The ionic potential differences of the transition metals between\nthe two layers at the stacking fault are 0.3552, and the cationic potential difference is 0.08532, if the\nNa content is assumed to remain constant (0.7). We have revised in the manuscript and attached the\nrelevant figure in supplementary information.\nIn manuscript, page 11 line 215:\nBy comparing the counts of EDX spectra, The ratios of transition metals at the two stacking faults\nregion are approximately Mn:Ni = 0.815:0.185 and 0.789:0.211; as well as 0.789:0.211 and\n0.805:0.195, respectively. The ionic potential differences of the transition metals between the two\n32\nlayers at the stacking fault are 0.3552, and the cationic potential difference is 0.08532 with constant\nNa content (0.7) (Supplementary Fig. 9).\nFig R18/Supplementary Fig. 9. The cation potential phase map of Na Mn Ni O (0≤x≤0.5).\n0.7 1-x x 2\n3. On pg. 14, the authors’ analysis of the CV sweeps of their cathodes glosses over several peaks\nthat may be important for their story. First, on line 255 it is stated that the 4.2V peak corresponding\nto the P2-O2 transition is not observed, but this peak does appear to be present, albeit shifted to\nslightly lower potentials. Likewise, on lines 257-258 they claim to see a single peak in the SF-\nNaMN at 4.11/4.38V, but this is clearly part of a doublet with the previously mentioned peak.\nLikewise, the authors imply that the SF-NaMN does not show a peak near 3.78/3.64V (line 263)\nrelated to Na+ vacancy ordering, but the peak (albeit diminished) is present. They ultimately\nconclude that the P-O transition is “missing” (pg.15, line 270). We request that the authors revisit\ntheir analysis and document that while the SF-NaMN appears to exhibit significantly reduced P-O\ntransition and vacancy ordering, it does not prevent it. Instead of eliminating the process, the effect\n33\nof the artificially-introduced slip faults appears to inhibit much of the negative impact of the P-O\ntransition. Please supplement/support your CV analysis with dQ/dV analysis from the battery\ncycling data, as this may better highlight the additional peaks discussed above.\nResponse: Thank you for your valuable comments. We sincerely appreciate the reviewer’s incisive\nobservations regarding our CV sweep analysis, which have highlighted critical nuances we\noverlooked in our initial discussion. We supplemented dQ/dV analysis to accurate convey the\nelectrochemical behavior of the SF-NaMN cathodes. Below, we address each point raised and\noutline the revisions to our analysis:\n1. Re-evaluation of the redox reaction peak at 4.11/4.38 V.\nTypically, P2-phase cathodes, which usually undergo a P2-O2 phase transition at 4.2V, this transition\ncorresponds to a very flat plateau in charge-discharge curves. Consequently, the presence of 4.2V\nplateau is empirically used to estimately determine the occurrence of a P2-O2 phase transition. In\nfact, the redox peaks in CV curves are defined as current peaks generated when redox couples\nundergo electron transfer (oxidation or reduction) during the potential scan. As long as the\nconsumption rate of reactants is much higher than the diffusion rate, resulting in the formation of a\nconcentration gradient on the electrode surface, there will be a net current. Then the characteristic\npeaks will show up in the CV test. Therefore, the phase transitions should be determined through\nstructural analysis such as XRD, but not on the electrochemical results of GCD or CV curves.\nTo confirm whether a phase transition occurred in our material, we ultimately relied on\ncharacterizations such as in-situ XRD and transmission electron microscopy (HADDF-STEM). In\nthe in-situ XRD results, only a slight decrease in the (104) peaks of P2-phase of SF-NaMN was\nobserved, with no new peaks appearing; similarly, no interlayer contraction of SF-NaMN were\n34\ndetected via HADDF-STEM (More detailed response in comment 4). Thus, P-O phase transition is\nnot observed in SF-NaMN via in-situ XRD and HADDF-STEM measurements, this leads to the\ndisappearance of plateaus in the charge-discharge curves of SF-NaMN, which instead transform\ninto sloped curves. Upon re-examining the CV data and dQ/dV results (Fig R19), the two redox\npeaks of SF-NaMN at 4.11 and 4.38 V corresponding to the redox reactions of Ni3+/Ni4+, O2-/On- (as\nevidenced by XAFS and XPS results), 35-37 while a single peak (4.28 V) is exhibited in the CV and\ndQ/dV curves of NaMN. This difference may be attributed to the changes in the coordination\nenvironment of these two materials at high voltage.\n2. Persistence of the Na⁺ vacancy ordering peak (≈3.78/3.64 V)\nWe concur with the reviewer that a diminished peak near 3.78/3.64 V-associated with Na⁺ vacancy\nordering-is present in SF-NaMN. In SF-NaMN, Na+/vacancy ordering still exists, but its degree is\nsignificantly reduced compared to that in NaMN. This result strengthen the conclusion that pre-\nconstructed stacking faults significantly reduce the Na+/vacancy ordering.\nBased on this reanalysis, we have revised our conclusion in manuscript, page 14 line 263:\nThe peak intensities of Na+/vacancy ordering (3.78/3.64 V) are significantly reduced for SF-NaMN\ncompared to that of NaMN, suggesting that pre-constructed stacking faults can effectively suppress\nthe formation of Na+/vacancy ordered structures. The CVs and dQ/dV curves of SF-NaMN and\nNaMN show two pairs of redox peak in the high voltage region (4.11/4.38 V) corresponding to the\nNi3+/Ni4+ and O2-/On- redox reaction 5,6,29 (Supplementary Fig. 13), which is confirmed by ex situ\nX-ray absorption spectroscopy (XAS) and X-ray photoelectron spectroscopy (XPS) (Supplementary\nFig. 14-15, supplementary Note 3).\ns−1. b dQ/dV curves of SF-NaMN and NaMN\n4. Related to Comment 3 above, in Figure S29a it appears there is a very slight weakening of the P2\nphase around 39.75 deg. at the end of charge that would correspond to the P-O transition. Please\ncomment on this in relationship to the electrochemical evidence for the reduction (but not\nelimination) of this transition.\nResponse: Thank you for your valuable comments. The P-O phase transition is characterized by the\ngradual decrease in intensity and a shift toward higher angles of the (002) diffraction peak in XRD\npatterns during the charging process, which reflects the decrease of interlayer spacing.3,4,38 In this\nwork, we observed a slight reduction in the intensity of the (104) peak of SF-NaMN at the end of\ncharging, demonstrating relative interlayer gliding between the transition metal slabs, leading to\natomic displacements and broadening of the (104) peak. Generally, the formation of the O2 phase\nis determined by the new characteristic (002’) peak, that appears around 20 degrees.19,20 Meanwhile,\nthe formation of OP4 phase generally depends on whether there is a phenomenon of continuous\nhigh-angle shift (17-19°) in the in-situ XRD of (002) peak of the P2 phase.38\nHowever, the persistence of the P2-phase (002) peak and the absence of characteristic (002’) peak\nfrom the O2-phase indicate that no interlayer contraction of SF-NaMN occurred during this process\n36\n(Fig R20/Fig 4a-b). This can further provide evidence that the gliding of transition metal slabs and\nthe contraction of interlayer spacing are not inherently correlated, and can be decoupled by pre-\nintroducing stacking faults in P-type lattice structure.\nTherefore, by combining the microstructural evolution data with charge/discharge profiles, dQ/dV\nanalysis and CV curves, we demonstrate that the SF-NaMN still undergoes Ni3+/Ni4+ and oxygen\nredox reactions above 4.2V. However, unlike conventional P2-O2 phase transitions, the sodium\nextraction in SF-NaMN does not induce interlayer contraction, indicating a fundamentally different\nstructural evolution mechanism.\nWe have revised in manuscript page 21, line 378:\nTo understand the structural evolution of SF-NaMN and NaMN during charge and discharge process,\nin-situ XRD measurements are conducted. In-situ XRD patterns illustrate that, the characteristic\n(002’) peaks at 20o and 17-19o, corresponding to O2 and OP4 phases, do not show during the charge\nand discharge process of SF-NaMN, indicating irreversible P to O phase transitions was not\nobserved in SF-NaMN. Notably, a slight decrease in the intensity of the (104) peak at 39.75 degrees\nat the end of charging, demonstrating relative interlayer gliding between the TMO slabs happened,\n2\nleading to atomic displacements and broadening of the (104) peak. However, the (002) peak did not\nshift and no new peaks formed indicating that no obvious interlayer contraction of SF-NaMN\noccurred during this process. The c-axis variation for SF-NaMN is calculated as 0.9% at 4.0 V, and\n1.45% within 2.4-4.5 V.\ndischarge curves at 20 mA g-1 between 2.4-4.5 V of Swagelok cell of SF-NaMN and NaMN\n5. On Pg. 24, lines 438-441 the authors argue that they have decoupled slab sliding from lattice\ncontraction; this does not seem accurate. Rather, the evidence appears to suggest that by introducing\na small quantity of slip fault defects the authors have been able to prevent additional significant\nvolume change and slab gliding upon desodiation (Figs. S29 and 4g) in the SF-NaMN compared to\nthe standard material. We request the authors either a) provide more evidence/a clearer explanation\nfor why they believe that introducing a limited number of faults represents “decoupling” of the two\naspects, or else b) adjust their message to better reflect that the induced faults appear to inhibit the\nconcerted expansion and shifting of slabs from occurring instead of decoupling them.\nResponse: We sincerely appreciate the reviewer’s recognition of our work. The traditional view\nholds that during the deep sodium extraction process, the P2-O2 phase transition inevitably involves\ntransition metal slab sliding accompanied by corresponding c-axis contraction.\nWe believe the discrepancy regarding the term “decouple” may stem from differing understandings\nof its definition. Following the reviewer’s suggestion, we have consulted the Oxford Dictionary,\nwhose core definition of “decouple” is “to end the connection or relationship between two things.”\nSpecifically in our work, we define “decouple” as a scenario whereby stacking faults are pre-\nengineered into the P-phase lattice in the initial structure, ensuring that the interlayer spacing of\nthe P2-phase material does not contract during charge-discharge cycling.\nIn the conventional understanding, TMO slab gliding and abrupt c-axis contraction are intrinsically\n₂\ncoupled phenomena that jointly define the P2-O2 phase transition during deep sodium extraction.\nFor a typical P2-type layered oxide, this usually happens at 4.2V, when 50% Na removal is achieved,\nduring which a large amount of TMO slabs glide along [1/3, 2/3, z] or [2/3, 1/3, z] direction.3,4 In\n2\nthe meantime the Na+ coordination environment transforms from prismatic (P-type) to octahedral\n(O-type), accompanied by abrupt c-contraction. Generally, the TMO slabs gliding, or the Na+\n2\ncoordination change are regarded as the fundamental origin of the P2-O2 phase transition, while c-\ncontraction is considered as an accompanied structural response.2,5\nHowever, we have demonstrated that the slab gliding is not always coupled with c-contraction. If\ncertain amount of TMO slabs (15.8%) are pre-glided from O to P positions, i.e. O-type stacking\n2\nfaults are pre-engineered in the initial P-phase lattice, such a structure remains stable at high voltage.\nNo abrupt c-axis contraction happens even with 70% Na removed (4.5V). Therefore, distinguished\nfrom traditional thought, the slab gliding is not necessarily always accompanied with c-axis\ncontraction. Therefore, we use the word ‘decouple’ refers to breaking the obligatory co-occurrence\nbetween TMO slab gliding and abrupt c-axis contraction.\ncycling is significantly smaller than that of NaMN, and most importantly, no abrupt c-axis\ncontraction is observed despite the occurrence of slab gliding. This outcome directly demonstrates\nthat we have successfully decoupled transition metal gliding from the c-axis contraction that is\ntraditionally considered inevitable in the P-O phase transition. By pre-constructing O-type stacking\nfaults in the P2 layered cathode, we effectively suppress both excessive volume expansion and the\nabrupt c-axis contraction during charge-discharge processes, validating our definition of \"decouple\".\nThis aligns with the core definition of “decouple” (ending the inherent connection between two\nphenomena), and we have included this specific definition in the paper.\nWe revised in manuscript, page 5 , line 103:\nThis work aims to decouple the TMO slab gliding from the c-lattice contraction, which generally\n2\nhappen co-currently during the charging and discharging process. To disrupt this inherent\ncorrelation, one can consider pre-establish certain amounts of O-type slabs into P-type layered\nstructurewith unchanged interlayer distances (O-Na-O).\nTo help readers better understand our concept of decoupling, we have added this sentence in the\nmain text. In manuscript, page 22, line 411.\nTherefore, although 15.8% TMO slabs have been pre-glided from P to O position, no significant\n2\nc-contraction was observed in SF-NaMN. In this sense, the traditional correlation between TMO\n2\nslab gliding and c contraction has been decoupled.\nMinor Comments\n1. Please provide experimental and instrument details regarding your XPS and ToF-SIMS\nmeasurements.\nResponse: Thank you for your comments. We have modified and highlighted with yellow in method\nsection.\nIn manuscript, Page 34, line 630:\nTOF-SIMS is conducted in negative ion mode with a region size of 80 micrometers × 80\nmicrometers, a depth of 200 nm, and signal intensities of all species in TOF-SIMS have been\n40\nnormalized. X-ray photoelectron spectroscopy (XPS) measurements were performed using an RBD-\nupgraded PHI-5000C ESCA system with Al Kα radiation.and a testing duration of 1 hour.\n2. Please ensure all acronyms (e.g., CEI, SEI) are defined immediately upon introduction, and check\nto confirm that SEI is used to refer to the film atop the anode, while CEI is used to refer to the film\natop the cathode.\nResponse: Thank you for your comments. We have revised the description in the DRT analysis.\nIn manuscript, Page 18, line 335:\nThe electrochemical impedance spectra (EIS) are resolved by this analysis with different\nelectrochemical processes39-42 , six main peaks located at 10-6 to 10-4 (R ), 10-3 to 10-2 (R ), 10-1\n0 film\nto 100 (Rct , Rct ), 100 to 101 (Z ), in accordance with ohmic resistance, resistance of solid\n1 2 W0\nelectrolyte interphase (SEI, refer to the film atop of anode) and cathode electrolyte interphase (CEI\nfilm, refer to the film atop of cathode), charge transfer resistance of anode material, charge transfer\nresistance of cathode material and Warburg resistance, respectively.\n3. Please provide a better explanation for how the various equations for calculating the percentage\nof stacking faults are used. It is currently unclear how any of them lead to a percentage. (pg.9)\nResponse: Thank you for your comments. According to the Scherrer equation: (1),\n𝐾𝐾𝐾𝐾\n𝛽𝛽𝑠𝑠𝑠𝑠𝑠𝑠𝑠𝑠 = 𝐿𝐿𝐿𝐿𝐿𝐿𝑠𝑠𝐿𝐿\nand the characteristics of broadening in hexagonal close-packed (HCP) diffraction lines43, when h-\nk=3n±1 and l=even, the diffraction line broadening is significant, whereas for l=odd, broadening is\nminimal. When only stacking faults exist, the integral breadth β for an (hkl) reflection relates to\nstacking fault broadening as follows: (2)\n𝛽𝛽𝑓𝑓𝐿𝐿\n𝛽𝛽𝑓𝑓 𝛽𝛽 =cos𝜑𝜑𝑧𝑧\n𝑓𝑓𝐿𝐿𝐿𝐿𝑠𝑠𝜑𝜑𝑧𝑧 (3)\n𝑤𝑤ℎ𝑒𝑒𝑛𝑛 𝑙𝑙 𝑖𝑖𝑖𝑖 𝑂𝑂𝑂𝑂𝑂𝑂 𝑛𝑛𝑛𝑛𝑛𝑛𝑛𝑛𝑒𝑒𝑛𝑛,𝛽𝛽𝑓𝑓 = 2𝐿𝐿\nℎ−𝑘𝑘 =3𝑛𝑛±1 3𝑓𝑓𝐿𝐿𝐿𝐿𝑠𝑠𝜑𝜑𝑧𝑧\n𝑤𝑤ℎ𝑒𝑒𝑛𝑛 𝑙𝑙 𝑖𝑖𝑖𝑖 𝑒𝑒𝑒𝑒𝑒𝑒𝑛𝑛 𝑛𝑛𝑛𝑛𝑛𝑛𝑛𝑛𝑒𝑒𝑛𝑛,𝛽𝛽𝑓𝑓 = 2𝐿𝐿\nwhere is the angle between the diffracting plane and the (001) plane, f is the stacking fault\n𝜑𝜑𝑠𝑠 41\nprobability, and c is the lattice parameter along the c-axis.\nWhen both crystallite size and stacking fault effects coexist, the Lorentzian approximation gives the\ntotal breadth as: (4)\n𝛽𝛽 = 𝛽𝛽𝑠𝑠𝑠𝑠𝑠𝑠𝑠𝑠 +𝛽𝛽𝑓𝑓\nFor isotropic near-spherical nanocrystals: (5)\n𝛽𝛽 = 𝛽𝛽𝐿𝐿+𝐴𝐴cos𝜑𝜑𝑠𝑠\n𝑓𝑓 (6)\n𝑤𝑤ℎ𝑒𝑒𝑛𝑛 𝑙𝑙 𝑖𝑖𝑖𝑖 𝑂𝑂𝑂𝑂𝑂𝑂 𝑛𝑛𝑛𝑛𝑛𝑛𝑛𝑛𝑒𝑒𝑛𝑛,𝐴𝐴 =2𝐿𝐿\nℎ−𝑘𝑘 =3𝑛𝑛±1 3𝑓𝑓\n𝑤𝑤ℎ𝑒𝑒𝑛𝑛 𝑙𝑙 𝑖𝑖𝑖𝑖 𝑒𝑒𝑒𝑒𝑒𝑒𝑛𝑛 𝑛𝑛𝑛𝑛𝑛𝑛𝑛𝑛𝑒𝑒𝑛𝑛,𝐴𝐴 = 2𝐿𝐿\nSolving for D (crystallite size) and f (stacking faults probability) using Equations (4)-(6) is complex.\nThus, a least squares method is proposed: Substitute (1) and (5)-(6) into (4), multiply by\ncos𝐿𝐿\n𝐾𝐾\nand simplify to: (7)\ncos𝐿𝐿 0.89 1 cos𝐿𝐿\n𝛽𝛽 𝐾𝐾 = 𝐷𝐷 +2𝐿𝐿 𝐾𝐾 𝑐𝑐𝑐𝑐𝑖𝑖𝜑𝜑𝑠𝑠4𝑓𝑓\nDefine variables: Y=\ncos𝐿𝐿 0.89 1 cos𝐿𝐿\n𝛽𝛽 𝐾𝐾 , a= 𝐷𝐷, X= 2𝐿𝐿 𝐾𝐾 𝑐𝑐𝑐𝑐𝑖𝑖𝜑𝜑, 𝑠𝑠F=4f.\nEquation (7) becomes: (8)\n𝑌𝑌 =𝑎𝑎+𝐹𝐹𝐹𝐹\nThe regular matrix equation for least squares is: (9)\n𝑛𝑛 ∑𝐹𝐹 𝑎𝑎 ∑𝑌𝑌\n 2  = \nSolutions are: D = (10), F= ∑𝐹𝐹 ∑𝐹𝐹 (11𝑛𝑛) ∑𝐹𝐹𝑌𝑌\n2 2\n𝑠𝑠∑𝑋𝑋 −(∑𝑋𝑋) 𝑠𝑠∑𝑋𝑋𝑌𝑌−∑𝑋𝑋∑𝑌𝑌\n2 2 2\n0.89×∑𝑋𝑋 ∑𝑌𝑌−∑𝑋𝑋∑𝑋𝑋𝑌𝑌 𝑠𝑠∑𝑋𝑋 −(∑𝑋𝑋)\nFor SF-NaMN, the stacking fault probability is 15.8%.\n𝛽𝛽\nWe have supplemented this description in the method of manuscript, page 37, line 682.\n4. On pg.14, please invert the x-axis of your XPS plots in line with standard convention of binding\nenergies increasing to the left.\nResponse: Thank you for your comments. We have inverted x-axis of our XPS results.\nto 4.2 V, 4.5 V and discharged to 3.2 V, 2.4 V.\n5. On pg. 17, please define relaxation time and describe the DRT analysis more fully.\nResponse: Thank you for your comments. The relaxation time (τ) is the characteristic time scale\nrequired for a system to recover from a non-equilibrium state to an equilibrium state, which reflect\nthe rate of electrochemical reactions or ion diffusion. We have revised the discussion regarding the\nDRT analysis.\nIn manuscript, Page 17, line 323:\nTo have an in-depth understanding on the superior electrochemical performance of SF-NaMN, the\ndifference of the Na+ kinetic transport between SF-NaMN and NaMN at high voltage (> 4.2 V) in\nthe full cell system was systematically studied through the distribution of relaxation times (DRT)\nanalysis, which employs a model-free approach to analyze impedance evolution. This mathematical\ntransformation converts frequency-domain Nyquist plots into time-domain distribution of relaxation\ntimes (DRT) spectra, enabling a model-independent assessment of electrochemical processes.\n(1)\n∞ 𝛾𝛾(𝜏𝜏)\n𝑍𝑍(𝜔𝜔)= 𝑅𝑅∞+∫0 1+𝑗𝑗𝑗𝑗𝜏𝜏𝑂𝑂𝑑𝑑\nPeaks at specific relaxation times correspond to distinct electrochemical processes, while their\n43\nrespective areas quantify the associated impedance values. The relaxation time (τ) is the\ncharacteristic time scale required for a system to recover from a non-equilibrium state to an\nequilibrium state, which reflect the rate of electrochemical reactions or ion diffusion.\nAs Na+ deep desodiated, R for HC ||SF-NaMN gradually increased and moved towards longer\nct1 ups\nrelaxation time, this could be attributed to the slow Na+ transport kinetics within the almost fully\nsodiated HC at 4.35-4.4 V. Notably, the Warburg resistance (Z\nW0\n), which characterizes the bulk Na⁺\ndiffusion behavior, showed no obvious variation during the first cycle, indicating that the diffusion\nprocess remains relatively stable during charging process. The R and Z resistance could return\nCt2 W0\nto their original response position during the discharging process. This reversibility of the\nimpedance-derived parameters serves as compelling evidence for the exceptional electrochemical\nreversibility exhibited by the HC ||SF-NaMN. In contrast, for the HC ||NaMN full cell, both R\nups ups ct2\n(charge transfer resistance of cathode material) and Warburg resistance (Z ) increased significantly\nW0\nduring charging process and can not completely recover during the discharging process, indicating\nthe charge transfer and bulk Na+ diffusion process were severely inhibited by irreversible phase\ntransition of NaMN.\n6. On pg. 18, please provide a citation for the use of pouch cell performance as the “golden criterion”,\nor else remove the phrase and use a more generic phrase.\nResponse: Thank you for your comments. We have replaced “Pouch cell performance are\nrecognized as the “golden criterion” to evaluate the practical potential of electrode materials.” with\n“A 0.5 Ah pouch cell are assembled to further develop the practical potential application of SF-\nNaMN” , and highlighted with yellow in manuscript page 19, line 356.\n7. On pg. 23, please provide a better description for how the slab energy is related to the propensity\n44\nof the materials to undergo the P-O transition.\nResponse: Thank you for your comments. We have revised the description in the manuscript.\nIn manuscript, Page 25, line 457:\nThe TMO slab gliding behavior of TMO slabs in Na Mn Ni O (Na Mn Ni O ) for SF-\n2 2 6 24 6 60 0.2 0.8 0.2 2\nNaMN and NaMN along [1/3, 2/3, z] direction were further simulated. It should be noted that the\nslab gliding energy is calculated as the energy of the structure after gliding minus the energy of the\nstructure before gliding, and this energy value is closely related to the propensity of the material to\nundergo the P-O transition: a lower energy value indicates that the slab structure is more likely to\nglide, and thus the material is more prone to transform from P-type to O-type. From the simulation\nresults, the gliding energies for slab-1 and slab-2 are 0.249 eV and 0.375 eV of SF-NaMN,\nrespectively, which are much greater than the energy required of slab-3 (-0.449 eV) of NaMN (Fig.\n4g), indicating that NaMN is more prone to undergo a transition from P-type to O-type under deep\ndesodiation.\n8. Figure S29 is missing the c-axis data.\nResponse: Thank you for your comments. The variation of the c-axis in in-situ XRD is shown in\nFig. 4a and 4b.\n9. On pg. 27, please include if the EDX data was collected using an SEM or STEM, and describe\nwhy the data indicates that TM dissolution is reduced in the SF-NaMN sample.\nResponse: Thank you for your valuable comments. the EDX (energy-dispersive X-ray spectroscopy)\ndata were all collected using an EDX detector attached to a scanning electron microscope (SEM).\nThe reason for choosing SEM-EDX is that it is suitable for rapid characterization of the large-area\nelemental distribution on the surface and cross-section of electrode materials, enabling a more\n45\ncomprehensive capture of possible dissolution of transition metals (TMs). On the electrode surface\nand separator of NaMN, distinct characteristic peaks of transition metal (TM) elements (Mn, Ni)\nwere detected, with their atomic percentages (Mn: 1.37%, Ni: 1.16%) significantly higher than those\nof the SF-NaMN sample (Mn: 0.23%, Ni: 0.4%). This indicates that more TM ions dissolved from\nNaMN and migrated to the separator during cycling process.\nWe have revised the description in manuscript page 29, line 538.\nEnergy dispersive x-ray spectroscopy (EDX) of the elemental information on the Na anode reveals\nthat the residual TM element signals in SF-NaMN (Mn: 1.37%, Ni: 1.16%) are significantly weaker\nthan those in NaMN (Mn: 0.23%, Ni: 0.4%), also indicates that the TM dissolution in SF-NaMN\nhas been significantly suppressed\nFig R22/Supplementary Fig. 44 Energy Dispersive X-ray spectroscopy (EDX) spectra of a SF-\nNaMN, b NaMN of anodes site of the separator, which stopped at 2.4 V in discharge state after 30\ncycles.\n46\n10. On pg. 27, please include an FFT of the CEI of the SF-NaMN instead of just the bulk particle\nto provide a more fair comparison to the NaMN sample.\nResponse: Thank you for your valuable comments. We have redone the FFT of HR-TEM results\nincluding the CEI interface film.\nFig R23/Supplementary Fig. 45 The TEM image of SF-NaMN a and NaMN c after 30th cycles. b,\nd FFT transform of boxes in a and c.\n11. On pg. 27 lines 489-491, please provide more detail as to why the EIS indicates a) superior ionic\nconductivity and b) increased stability. It is not immediately obvious from the data alone.\nResponse: Thank you for your valuable comments. We have supplemented this description in\nsupplementary information, page 70, note 10.\na) By comparing the impedance spectra of SF-NaMN with that of NaMN, it is found that the arc in\nthe high-frequency region of SF-NaMN is significantly smaller. We fitted the impedance spectra.\nSEI\nΩ and 5.804 Ω, respectively, which are significantly lower than those of NaMN (750 Ω and 249 Ω,\nrespectively). This indicates that the interface charge transfer resistance of SF-NaMN is lower than\nthat of NaMN, which is beneficial to ionic conductivity.\nb) Through comparison of EIS results after cycling tests, we found that the variation of impedance\nspectra (Rct and R ) of SF-NaMN is obviously lower than that of NaMN, indicating that the\nSEI\nchanges in the interfacial layer and charge transfer of the SF-NaMN are relatively slight, and the\nion transport capability remains stable with the increase in cycle numbers.\nFig R24/Supplementary Fig. 46 EIS curves of SF-NaMN and NaMN electrodes at pristine state,\nafter 50, 100 and 200 cycles.\n48\n12. On pg. 28 line 496, the authors mention that CO2 is evolved much earlier from SF-NaMN than\nfrom NaMN, but do not provide additional discussion. If SF-NaMN releases less CO2 overall, but\nbegins at lower potentials, what is the implication?\nResponse: Thank you for your careful attention. After rechecking the experimental data and original\nrecords, we found that there was an typo in the description of the CO₂ evolution timing in line 496\non page 28, and we sincerely apologize for this. The correct result is: CO₂ evolution from SF-NaMN\nstarts later than from NaMN, rather than earlier. We have corrected this content in the revised\nmanuscript.\nIn manuscript, page 30, line 553: Carbon dioxide (CO ) evolution was observed at 4.331 V vs.\n(occurring at 4.171 V vs. Na+/Na).\n13. Please provide xy scale bars for the ToF-SIMS 3D chemical maps in Fig.5h,j.\nResponse: Thank you for your comments. We have revised Fig 5 h,j and provided xy scale bars for\nthe ToF-SIMS 3D chemical maps.\nFig R25/Fig. 5h-j TOF-SIMS profiles and 3D visualization of architecture evolution of cycled SF-\nNaMN and NaMN.\n1. House, R. A., et al. Superstructure control of first-cycle voltage hysteresis in oxygen-redox cathodes.\nNature 577, 502-508 (2020).\n2. Chen, Z., et al. 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Separation of predominant processes in electrochemical impedance\nspectra of lithium-ion batteries with nickel-manganese-cobalt cathodes. J. Power Sources 425, 121-129\n(2019).\n41. Plank, C., et al. A review on the distribution of relaxation times analysis: A powerful tool for process\nidentification of electrochemical systems. J. Power Sources 594, 233845 (2024).\n42. Raccichini, R., Amores, M. &Hinds, G. Critical Review of the Use of Reference Electrodes in Li-\nIon Batteries: A Diagnostic Perspective. Batteries 5, 12 (2019).\n43. Cheng, G.-f., Cheng, C.-z. &Huang, H.-y. Characterization and Research of Hexagonal Closed\nPacking Structured ZnO Nano-powders by X-ray Diffraction Method. Journal of Inorganic Materials 23,\n(2008).\n51\n52"}, {"round": 2, "reviewer_comments": {"reviewer_1": "We would like to thank the authors for their detailed response to the questions raised by us. All the questions have been appropriately addressed in the revised manuscript. Thus, we recommend the paper to be accept for publishment in Nature Communications.", "reviewer_2": "The authors clearly answered overall queries raised by the reviewer. The revised manuscript now meets the general standards of Nature Communications.", "reviewer_3": "Revised Review of Q. Shi et al., “Decoupling TMO Slab Gliding and Lattice Contraction to Enable High-Voltage Na-Ion Batteries”\nVerdict: Accept with Minor Revisions\nThe authors have made a reasonable attempt to address most of the comments from out previous review. In particular, we believe that the additional details regarding the WH measurements and the CV analysis improve the manuscript. That said, we would like to highlight a couple of areas that could use a final revision, after which point we believe the manuscript is suitable for publication in Nature Communications.\nOur remaining concerns largely center on the discussion revolving around the ”decoupling” of slab gliding and lattice expansion. We feel that the data presented clearly demonstrates the value of the reported work, but that the interpretation/discussion of the results is still incorrectly stated. This said, we do not want to reject this paper on purely semantic grounds when the readers will eventually be able to draw conclusions of the accuracy of the authors’ interpretations for themselves. Thus, while we would recommend the authors revisit their results and accompanying discussion, we believe that the manuscript is suitable for publication after addressing Comments 1 and 2 (both of which are minor comments).\nComments\n1. Regarding the improved description of the Mn/Ni quantification on pgs. 11-12, please improve the clarity of what slab the four ratios detailed are referring to. It appears that they may refer to 4 of the 5 boxed line maps in Fig. 2h from the top down (skipping 1). If this is the case, please clarify this in the text. Additionally, the ratios provided by the authors seem to indicate that the slabs adjacent to the Mn-rich slabs identified are actually Ni-rich; can the authors comment on this?\n2. Please provide figure references to the new discussion on pg. 21. The usefulness of the discussion is undercut with the absence of figures.\n3. With regard to our previous Comment 5, we thank the authors for attempting to address our question and for the additional context for their claim of decoupling the slab gliding and lattice contraction. We still disagree with the strength of the authors’ claim for two reasons; one is related to the pre-cycled material results and one is related to the post-cycled material results. [1] First, comparing the lattice spacings provided in Figs. 2a and 2b, one can see that the O-type faults are corelated to a ~0.14Å decrease in the lattice spacing relative to the NaMN material. We agree that this is smaller than the change observed under the full electrochemically-induced P-O transition and is thus a significant improvement, but the change is still present. If the authors can provide evidence that 15.8% of slabs gliding is significant relative to the number that slide during electrochemical desodiation, this would support their decoupling argument. [2] Second, the authors convincingly demonstrate that the classical P-O transition does not occur during desodiation using their in situ XRD, but they less convincingly show that the slab gliding does occur during desodiation. The XRD evidence (Fig. 33a and description on pg. 21) is very weak and unsupported by references, while the visual evidence by STEM (Fig. 4c) is oriented along the wrong axis to provide sufficient support that additional gliding occurred during the desodiation process. Indeed, the new DFT simulations of slab gliding behavior (pg. 25) provide further evidence for our argument; according to the authors, the SF-NaMN is less likely to glide and undergo the P-O transition, and thus would be less likely to undergo lattice contraction. If the authors can provide stronger evidence that slab gliding occurs during electrochemical desodiation, this would also support their decoupling argument. Based on the data and the updated definition(s) provided by the authors in their response, and in lieu of the additional evidence for points 1 and 2 above, we can suggest one of the following constructions which might be more semantically accurate:\na. (modifying the author’s stated definition in their response to Comment 5) “Specifically in our work, we define “decouple” as a scenario whereby stacking faults are pre-engineered into the P-phase lattice in the initial structure, ensuring that the interlayer spacing of the P2-phase material does not contract during charge-discharge cycling.”\ni. The authors have decoupled the (de)sodiation of the material from the changes in interlayer spacing the typically accompany this electrochemical process during cycling.\nb. (softening the author’s stance, given the pre-cycling structural data)\ni. The authors have largely/significantly decoupled slab gliding within the material from lattice contraction, demonstrating a much smaller contraction of the SF-NaMN lattice with gliding relative to NaMN (2.5% vs. 12.8%, respectively), with most of the contraction occurring prior to electrochemical desodiation.\nc. (based on the post-cycling structural data)\ni. The authors have decoupled the desodiation of the SF-NaMN material from the sharp contraction of the lattice that occurs during the P-O transition."}, "author_response": "Manuscript Number: NCOMMS-25-34868-B\nTitle: Decoupling TMO Slab Gliding and Lattice Contraction to Enable High-Voltage Na-Ion\nDear Editor and Reviewers,\nThank you very much for arduous work and valuable comments. We have studied the comments\ncarefully and we are pleased to find the constructive and positive comments and recommendation\non this work from all reviewers. The main problem raised by the reviewers, and the explanation of\nsome scientific issues, have been fully addressed in our revised manuscript. Here within the\nenclosure is our revised manuscript, which has been improved according to your suggestions.\nRevised parts are highlighted with yellow color, and the point to point responses to reviewers’\ncomments are listed below:\nComments from Reviewer#1:\nReviewer #1: We would like to thank the authors for their detailed response to the questions raised\nby us. All the questions have been appropriately addressed in the revised manuscript. Thus, we\nrecommend the paper to be accept for publishment in Nature Communications.\nResponse: Thank you very much for your positive comment and your recommendation on this\narticle. We also highly appreciate your constructive and valuable advices and suggestions.\nReviewer #2: The authors clearly answered overall queries raised by the reviewer. The revised\nmanuscript now meets the general standards of Nature Communications.\nResponse: Thank you very much for your positive comment and your recommendation on this\narticle. We also highly appreciate your constructive and valuable advices and suggestions.\nReviewer #3 (Remarks to the Author):\nRevised Review of Q. Shi et al., “Decoupling TMO Slab Gliding and Lattice Contraction to Enable\n2\nVerdict: Accept with Minor Revisions\nThe authors have made a reasonable attempt to address most of the comments from out previous\nreview. In particular, we believe that the additional details regarding the WH measurements and the\nCV analysis improve the manuscript. That said, we would like to highlight a couple of areas that\ncould use a final revision, after which point we believe the manuscript is suitable for publication in\nNature Communications.\nOur remaining concerns largely center on the discussion revolving around the ”decoupling” of slab\ngliding and lattice expansion. We feel that the data presented clearly demonstrates the value of the\nreported work, but that the interpretation/discussion of the results is still incorrectly stated. This said,\nwe do not want to reject this paper on purely semantic grounds when the readers will eventually be\nable to draw conclusions of the accuracy of the authors’ interpretations for themselves. Thus, while\nwe would recommend the authors revisit their results and accompanying discussion, we believe that\nthe manuscript is suitable for publication after addressing Comments 1 and 2 (both of which are\nminor comments).\nResponse: Thank you for your positive and valuable comments. We greatly appreciate your\nrecognition of our efforts to address most of the previous feedback, particularly your positive\nassessment of the additional details regarding WH measurements and CV analysis.\nComments\n1. Regarding the improved description of the Mn/Ni quantification on pgs. 11-12, please improve\nthe clarity of what slab the four ratios detailed are referring to. It appears that they may refer to 4 of\nthe 5 boxed line maps in Fig. 2h from the top down (skipping 1). If this is the case, please clarify\n2\nthis in the text. Additionally, the ratios provided by the authors seem to indicate that the slabs\nadjacent to the Mn-rich slabs identified are actually Ni-rich; can the authors comment on this?\nResponse: Thank you for your valuable comments. As suggested, we have calculated the transition\nmetal ratios across TMO slabs a to d (counted from top to bottom in Fig 2h) and updated both the\n₂\ndescription in the main text.\nIt should be noted that, to visually emphasize the inhomogeneous distribution of Mn and Ni, the\nintensity map shown in Fig. 2h was normalized. Consequently, although the ratio of Ni content\n(0.211) in the layer adjacent to the Mn-rich slab slightly exceeds the nominal ratio value (0.2) in\nNa Mn Ni O formula, the absolute amount of Mn remains substantially higher than that of Ni\n0.7 0.8 0.2 2\n(Fig R2). Therefore, this region should not be termed “Ni-rich”; it merely exhibits a slight\nenrichment in Ni relative to the stoichiometric composition. This clarification has been incorporated\ninto the caption of Figure 2i.\nHAADF images of SF-NaMN and NaMN projected along the [100] zone axis. c-d STEM-HAADF\nimages of SF-NaMN and NaMN projected along the [001] zone axis. e Corresponding simulated\nmodel of NaMN and SF-NaMN along the [100] zone axis. f Simulated structure of NaMN and SF-\nNaMN from [001] direction. Atomic EDS mapping of element distributions in the g, SF-NaMN and\ni, NaMN, and corresponding normalized element signal density of h, Mn and j, Ni atoms along the\n[100] zone axes. k A section of neutron diffraction (ND) peaks of SF-NaMN and NaMN.\nzones.\nIn manuscript, page 11, line 214.\nBy comparing the counts of EDX spectra, the ratios of transition metals at the two stacking faults\nregion are approximately Mn:Ni = 0.815:0.185 and 0.789:0.211; as well as 0.789:0.211 and\n0.805:0.195, respectively. These values correspond to a to d TMO slabs (counted from top to\n2\nbottom in Fig 2h).\nIn manuscript, page 14, line 258. The caption of Fig 2h.\nAtomic EDS mapping of element distributions in the g, SF-NaMN and i, NaMN, and corresponding\nnormalized element signal density of h, Mn and j, Ni atoms along the [100] zone axes.\n2. Please provide figure references to the new discussion on pg. 21. The usefulness of the discussion\nis undercut with the absence of figures.\nResponse: Thank you for your valuable comments. We have added the figure reference to the\ndiscussion on page .21.\nIn manuscript, page 21, line 386:\nin-situ XRD measurements are conducted. In-situ XRD patterns illustrate that, the characteristic\n(002’) peaks at 20o and 17-19o, corresponding to O2 and OP4 phases, do not show during the charge\nand discharge process of SF-NaMN, indicating irreversible P to O phase transitions was not\nobserved in SF-NaMN (Fig. 4a). Notably, a slight decrease in the intensity of the (104) peak at 39.75\ndegrees at the end of charging, demonstrating relative interlayer gliding between the TMO slabs\n2\nhappened, leading to atomic displacements and broadening of the (104) peak. However, the (002)\npeak did not shift and no new peaks formed indicating that no obvious interlayer contraction of SF-\nNaMN occurred during this process. The c-axis variation for SF-NaMN is calculated as 0.9% at 4.0\nV, and 1.45% within 2.4-4.5 V (Supplementary Fig. 33-35).\n3. With regard to our previous Comment 5, we thank the authors for attempting to address our\nquestion and for the additional context for their claim of decoupling the slab gliding and lattice\ncontraction. We still disagree with the strength of the authors’ claim for two reasons; one is related\nto the pre-cycled material results and one is related to the post-cycled material results. [1] First,\ncomparing the lattice spacings provided in Figs. 2a and 2b, one can see that the O-type faults are\ncorelated to a ~0.14Å decrease in the lattice spacing relative to the NaMN material. We agree that\nthis is smaller than the change observed under the full electrochemically-induced P-O transition and\nis thus a significant improvement, but the change is still present. If the authors can provide evidence\nthat 15.8% of slabs gliding is significant relative to the number that slide during electrochemical\ndesodiation, this would support their decoupling argument. [2] Second, the authors convincingly\ndemonstrate that the classical P-O transition does not occur during desodiation using their in situ\nXRD, but they less convincingly show that the slab gliding does occur during desodiation. The XRD\n6\nevidence (Fig. 33a and description on pg. 21) is very weak and unsupported by references, while\nthe visual evidence by STEM (Fig. 4c) is oriented along the wrong axis to provide sufficient support\nthat additional gliding occurred during the desodiation process. Indeed, the new DFT simulations\nof slab gliding behavior (pg. 25) provide further evidence for our argument; according to the authors,\nthe SF-NaMN is less likely to glide and undergo the P-O transition, and thus would be less likely to\nundergo lattice contraction. If the authors can provide stronger evidence that slab gliding occurs\nduring electrochemical desodiation, this would also support their decoupling argument. Based on\nthe data and the updated definition(s) provided by the authors in their response, and in lieu of the\nadditional evidence for points 1 and 2 above, we can suggest one of the following constructions\nwhich might be more semantically accurate:\na. (modifying the author’s stated definition in their response to Comment 5) “Specifically in our\nwork, we define “decouple” as a scenario whereby stacking faults are pre-engineered into the P-\nphase lattice in the initial structure, ensuring that the interlayer spacing of the P2-phase material\ndoes not contract during charge-discharge cycling.”\ni. The authors have decoupled the (de)sodiation of the material from the changes in interlayer\nspacing the typically accompany this electrochemical process during cycling.\nb. (softening the author’s stance, given the pre-cycling structural data)\ni. The authors have largely/significantly decoupled slab gliding within the material from lattice\ncontraction, demonstrating a much smaller contraction of the SF-NaMN lattice with gliding relative\nto NaMN (2.5% vs. 12.8%, respectively), with most of the contraction occurring prior to\nelectrochemical desodiation.\nc. (based on the post-cycling structural data)\ni. The authors have decoupled the desodiation of the SF-NaMN material from the sharp contraction\nof the lattice that occurs during the P-O transition.\nResponse: Thank you for your positive and valuable comments. We agree that precise expression\nis critical to conveying our findings clearly, and we will promptly revisit accompanying discussion\nto refine the description of “decoupling.” Our revisions will be guided by your insights to ensure\nthe interpretation aligns with experimental observations and avoids any potential ambiguities.\nFor the pre-cycled results, according to the XRD refinement results, the c-axis of the SF-NaMN is\n11.187 Å, which exhibits a slight contraction in comparison to the NaMN (whose c-axis is 11.238\nÅ).\nWith respect to your concern about the evidence that “15.8% of slabs gliding is significant relative\nto the number that slide during electrochemical desodiation”, we noted the following: The literature\ngenerally reports that TMO slab gliding and c contraction occur concomitantly, but to our\n2\nknowledge no prior study has defined a quantitative threshold of slab gliding required to trigger a\npronounced c contraction. Moreover, reliable calculation of sliding events from stacking fault\nanalysis demands very high XRD signal quality and sample purity, which makes it remain\nchallenging to accurately calculated the number of gliding of TMO slabs during the charge-\ndischarge process of material.\nAbout the STEM axis: We have further checked with TEM specialists and previous literature: For\nP2-phase cathodes, the atomic arrangement of transition metals can be clearly observed along either\nthe [100] or [010] zone axis direction, and it is also feasible to determine whether gliding of\ntransition metal layers has occurred along [100] zone axis direction, which has been reported in\nprevious studies.1-3 In this work, the atomic structure of SF-NaMN was projected along the [100]\n8\nzone axis, which allowed for clear visualization of atomic arrangements as well as the relative\ninterlayer gliding. Following your suggestion, we revisited the HADDF-STEM results in Fig. 4c\nand additionally provided a set of result with a larger field-of-view. The results show that the\nstacking fault density in Fig. 4c is 26.7%, while that in Fig. R4 is 27.4%. Compared with the pre-\ncycling stacking fault density (15.8%) of SF-NaMN, both values exhibit a significant increase\n(~11%) in stacking fault probability, indicating that certain amount of TMO slab gliding does occur\n2\nfor SF-NaMN during charge/discharge process, while no sharp contraction of interlayer spacing was\nobserved. However, since STEM only reflects the changes in the local microstructure of the material,\nthe stacking fault density obtained in this way may be not accurate.\nBased on our results and your valuable suggestions, we have adopted suggestion (b) and revised the\nmanuscript text to improve the precision of the wording. It is noted that, the 2.5% c-contraction for\nSF-NaMN suggested by the reviewers should be derived from the STEM measurement, whereas the\n12.8% contraction for NaMN is obtained from in-situ XRD. Because STEM probes local regions\nwhile XRD yields an averaged c value, these values are not directly comparable. To enable a fair\ncomparison, we re-calculated the c contraction of both samples based on XRD and STEM data,\nrespectively.\nSTEM derived values: the c-axis contraction of SF-NaMN is calculated based on the equation\n, and = 1.59%. For NaMN, the contracted interlayer spacing\n𝑑𝑑𝑃𝑃2−𝑑𝑑𝑆𝑆𝑆𝑆 5.65−5.56\n𝑑𝑑𝑃𝑃2 ×100 5.65 ×100\nmeasured in Fig 4d are 4.4 and 4.7 A, respectively. Therefore, the c-axis contraction of NaMN is a\nrange from = 17.544% to = 22.8%.\n5.7−4.7 5.7−4.4\n5.7 ×100 5.7 ×100\nXRD Rietveld refinement derived results: the c-axis contraction of SF-NaMN is\n11.238−11.187\n11.238 ×100\n= 0.454%, and for NaMN is 12.8% from the results of in-situ XRD in Fig 4b.\nrepresentation of the materials, the XRD results were adopted and the manuscript have been updated\naccordingly.\nFig R3/Fig 4c STEM-HAADF images of SF-NaMN after 10 cycles and stopped at 4.5 V along [100]\nprojection.\nFig R4 STEM-HAADF images of SF-NaMN after 10 cycles and stopped at 4.5 V along [100]\nprojection.\nIn manuscript, page 8, line 155.\nThe Rietveld refinement of XRD patterns of SF-NaMN and NaMN reveals that the c-axis of the SF-\nÅ) (Supplementary Fig. 6, Supplementary Table 6-7), which may be induced by stacking faults.\nIn manuscript, page 4, line 87.\nWe provide substantive evidence that the TMO slabs gliding and the volume change can be largely\n2\ndecoupled, while the significant c-axis expansion upon Na+ extraction and abrupt lattice contraction\nat high SOC serve as the primary drivers for the irreversible phase transitions.\nIn manuscript, page 26 , line 496.\nTherefore, we have significantly decoupled slab gliding within the material from lattice contraction\nby pre-engineering O-stacking faults with large interlayer distance in P2-type structure.\nConsequently, SF NaMN demonstrates a substantially reduced c-axis contraction as compared to\n-\nthat of NaMN (0.454% vs. 12.8%, respectively), with the contraction in SF-NaMN observed prior\nto electrochemical desodiation.\nReference\n1. Somerville, J. W., et al. Nature of the “Z”-phase in layered Na-ion battery cathodes. Energy Environ.\nSci. 12, 2223-2232 (2019).\n2. Wang, P.-F., et al. Na+/vacancy disordering promises high-rate Na-ion batteries. Sci. Adv. 4,\neaar6018 (2018).\n3. Zhao, C., et al. Revealing High Na-Content P2-Type Layered Oxides as Advanced Sodium-Ion\nCathodes. Journal of the American Chemical Society 142, 5742-5750 (2020).\n11"}]

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