Inhibiting inter-layer gliding in transition metal layered oxides through interphase engineering for sodium-ion batteries

Abstract

Inter-layer gliding induced phase transitions are widely recognized as the predominant cause of performance degradation in layered oxide positive electrode materials utilized in Na/Li-ion batteries. However, effectively restraining these phase transitions at a fundamental level poses a significant challenge. In this study, we elucidate that gliding at the X2/Y3 (X, Y = P or O) interphase layer can be thermodynamically inhibited through an energetically driven gliding-inhibition mechanism, by systematic structural analysis and correlated energy calculations. Building upon this insight, we propose interphase engineering as an effective approach to mitigate phase transitions. The resulting P2/P3-Na_0.46Mn_0.9Ni_0.1O₂ material, featuring dense and uniform P2/P3 interphases, exhibits notable enhancements in both cycling stability and rate capability. Detailed structure probing conducted through advanced atomic-level electron microscopy and synchrotron X-ray diffraction corroborates the role of the P2/P3 interphase structure in suppressing gliding and phase transition. Furthermore, the widespread applicability of the X2/Y3 interphase concept is validated through the successful implementation in several other extended X2/Y3 interphase materials. These findings provide further understanding of interphase phenomena and suggest a strategy to suppress phase transition in layered positive electrode materials.

Layered transition-metal oxides suffer from interlayer gliding induced phase transitions that degrade performance upon sodium (de)intercalation. Here, authors propose interphase engineering to suppress gliding-induced transitions, enhancing structural stability and electrochemical performance in sodium ion batteries.

Introduction

Layered transition-metal oxides, characterized by the general formula A_xMO₂ (where A = Li, Na, etc.; 0 < x ≤ 1; M = transition metals), have seen notable commercial success as intercalation positive electrode materials in Li (Na)-ion batteries following the pioneering discovery of LiCoO₂ by Goodenough in 1980^{1, 2–3}. The lattice structure of A_xMO₂ consists of parallel MO₂ slabs arranged along the c-axis, with A ions intercalated between them. The stacking sequence of MO₂ slabs determines the resulting phase types^{4, 5–6}. Compared to their lithium-based counterparts, sodium-based compounds (Na_xMO₂) exhibit a broader range of layered phases, including O3, P3, O2, and P2. This diversity is mainly attributed to the larger size of Na⁺ ions (~1.02 Å), which allows for a more diverse oxygen coordination environment and thereby expands the possibilities of MO₂ stacking sequences^{7, 8–9}. Although the phasic diversity of Na_xMO₂ offers considerable flexibility for structural and performance optimization^{10, 11–12}, complex phase transitions during the electrochemical intercalation/extraction of Na⁺ ions lead to significant volume fluctuations and structural degradation^{13, 14–15}. Consequently, the practical application of layered oxide continuous to be hindered by unsatisfactory performances, particularly in terms of cycling lifespan^16,17.

Considerable efforts, including widely-adopted cation doping, have been devoted to alleviate phase variations during charge/discharge process^{18, 19–20}. While some improvements have been achieved, fundamentally inhibiting phase transitions over a broad voltage range remains a significant challenging, primarily due to the energetically-driven gliding of MO₂ layers^{21, 22–23}. Specifically, duo to the the strong dependence of electrostatic interaction between MO₂ layers on Na content, the intercalation/extraction of Na ions can induce relative gliding of neighboring MO₂ layers in the a-b plane to minimize energy^24,25. For instance, in P2-type Na_xMO₂, the electrostatic repulsion between face-to-face oxygen ions in adjacent MO₂ layers gradually increases during Na extraction (charge process), resulting in the relative gliding of MO₂ layers and a transition to an energetically favorable O2 phase²⁶. Therefore, inhibiting gliding behavior of MO₂ layers, in principle, represents a promising strategy to mitigate phase transitions, and thereby enhancing the structural stability and cycling life of layered oxide positive electrodes.

Herein, we uncover a mechanism for intrinsically inhibiting MO₂ layer gliding at a coherent X2/Y3 (X, Y = P or O) phase interphase, which involves a shared MO₂ interphase layer between two phases, in the layered Na_xMO₂. This interphase tends to remain immobile during charge/discharge process, minimizing the overall system energy and thereby preventing subsequent phase transitions. Based on this insight, we prepared the Na_0.46Mn_0.9Ni_0.1O₂ positive electrode material featuring a well-designed P2/P3 interphase structure. The ex-situ atomic-level electron microscopy results clearly show a restriction of MO₂ layer gliding at the interphase region, while in-situ synchrotron X-ray diffraction results confirm a significant reduction in phase transitions. Consequently, this P2/P3 positive electrode demonstrates notable improvements in both cycling stability and rate performance compared to its pure phase counterparts. Furthermore, the universality of this interphase engineering is confirmed through successful implementation in several other extended P2/P3 and P2/O3 positive electrodes. These findings elucidate the role of interphases in inhibiting gliding and offer the design principles for mitigating phase transitions in layered oxide positive electrode materials.

Results

X2/Y3 coherent interphase and gliding-inhibition mechanism

The phase structures of layered oxide Na_xMO₂ can be classified into four primary phases P2, O2, P3 and O3 by delmas’s notation, according to the Na coordination environment and the stacking sequence of the oxygen layers⁴. Some derivative phases (P′2, O′3, etc.) have also been reported recently, but the gliding behavior occurring during transformations between the primary phases is more representative (Supplementary Figs. 1 and 2). During electrochemical (de)sodiation, changes in system energy drive the transformation between various phases⁸. Notably, the transformation between the phases with repeated oxygen layers of 2 (P2, O2, etc.) and that of 3 (O3, P3, etc.) cannot occur, as they require breakage or reformation of M-O bonds^13,27. Given this, the X2/Y3 (where X, Y = P or O) type interphases are first considered, and a special X2 + Y3 intergrown crystal model was constructed to represent the X2/Y3 interphase structure (Supplementary Fig. 3). Taking P2/P3 model as an example, the model diagram (Supplementary Fig. 4) and crystallography information (Supplementary Table 1) show that the two-phase hybrid supercrystal cell consists of alternating P2 and P3 phase unit cells along the c-axis, with the shared coherent MO₂ layer defined as the P2/P3 interphase (indicated in two colors). Obviously, this concept of interphase emphasizes the phase boundaries between different phase regions within a grain, which is entirely distinct from the surface of the material. For the interphase gliding phenomena, Fig. 1a focuses on one P2/P3 interphase region (the dashed box area in Supplementary Fig. 4) and illustrates two potential gliding patterns within the a-b plane. After gliding by pattern 1, the oxygen atoms of interphase MO₂ layer glide from their original C and A sites to the A and B sites, respectively, resulting in the formation of a new structure (denoted by structure 1, Fig. 1b) accompanied by a change in the overall oxygen stacking sequence from original ABBCCAAC to ABBCABAC. In the scenario of pattern 2, the other glide-driven new structure (denoted by structure 2, Fig. 1c) presents an oxygen stacking sequence of ABBCBCAC. Further, the structural formation energies were calculated through density functional theory (DFT) to evaluate the energy variations during the gliding process. Figure 1d shows the formation energies of original structure and the two glide-driven structures (Supplementary Data 1) with a constant Na concentration of 0.67 (x = 0.67), where the multiple data points for each structure correspond to different configurations considered in the calculations. The results indicate that, regardless of the gliding pattern, the formation energy of the resulting new structure is higher than that of the original one, implying an energetically unfavorable process for such gliding of interphase MO₂ layer. Further, to better align with real charge-discharge conditions, the formation energies of the three structures with various Na concentrations were also calculated. As depicted in Fig. 1e (Supplementary Data 1), the original structure consistently remains the lowest formation energy over a wide range of Na concentrations compared to the other two glide-driven structures. These results indicate that the interphase MO₂ layer tends to remain immobile during the electrochemical process, thereby avoiding an increase in the system’s energy.

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Fig. 1

Schematic diagram of supposed P2/P3 interphase gliding and the gliding-inhibition mechanism.

a The two gliding patterns of the interphase MO₂ layer and corresponding changes of oxygen layers sequence. The top-right diagram, corresponding to the [001] perspective, depicts the three oxygen atom sites (A/B/C) and indicate the slip directions with arrows. b, c The new structure 1(2) after the interphase gliding by pattern 1(2), with oxygen layers sequence of ABBCABAC(ABBCBCAC). d The calculated formation energy of the original structure and the two new structures, with a fixed Na concentration of 0.6. e The calculated formation energy of the three structures with various Na concentration, with the highest value sets to 0. f, g The phase evolution induced by the supposed interphase gliding behaviors in the scenario of electrochemical sodiation/desodiation process. The color of sodium prismatic or octahedra signifies the energy favorability of phase transitions: red indicates unfavorable, and green denotes favorable transitions. The not-pass symbol-“X” represents energetically unfavorable phases in the current state. Evidently, in all cases, the glide of interphase leads to the emergence of at least one energetically unfavorable phase; hence, the interphase gliding is energetically inhibited. Source data are provided as a Source Data file.

This phenomenon of low-energy maintenance can be explained by a gliding-inhibition mechanism associated with phase structure. Specifically, for gliding pattern 1 (Fig. 1f) in P2/P3 structure, according to the oxygen layer stacking sequences, the transformation toward structure 1 during the desodiation (charge) process results in the phase transitions from P2 to O2 structure (...ABAC ...) and P3 to O3 structure (...BCAB ...). However, the O3 structure is energetically unfavorable under sodium-deficient state, thereby significantly increasing the formation energy as shown in Fig. 1e. During the sodiation (discharge) process, the interphase gliding results in the formation of a sodium-rich O2 phase, which is also unfavorable in energy^28,29. Similarly, in the case of pattern 2 (Fig. 1g), the glide toward structure 2 will make the original P3 structure convert to infrequent O1 structure (...BCBC ...), which is an unfavorable structure under both charge and discharge conditions^8,30. As a result, the entire P2/P3 hybrid structure suppresses the gliding of MO₂ layer during the electrochemical process. In short, the gliding behavior of the coherent interphase is jointly governed by the two phases constituting it. Due to their conflicting demands for the gliding, the interphase layer is inclined to adopt a compromise solution without gliding.

Moreover, we further validated this gliding-inhibition mechanism in the remaining X2/Y3 interphases (O2/O3, O2/P3, and P2/O3 coherent interphase) via analogous DFT calculations of formation energy and phasic evolution analyses (Supplementary Fig. 5‒7, Supplementary Data 1). Meanwhile, the non-X2/Y3 interphases (O2/P2 and O3/P3 interphase) were also studied. It was found that the non-X2/Y3 interphases do not function as inhibiting the interlayer gliding unlike the X2/Y3 interphase (Supplementary Figs. 8, 9), which is consistent with the previous studies^31,32. These results theoretically demonstrate the role of rational interphase (X2/Y3 type) in inhibiting the interlayer gliding behavior of transition metal layered oxides during the electrochemical (de)sodiation process.

Interphase density and interphase uniformity

In a realistic material, the quantity and microscopic distribution of the X2/Y3 interphases are clearly crucial factors in achieving overall gliding inhibition. In general, the interphase should be as abundant and uniformly distributed as possible. However, this vague criterion poses a challenge for the interphase design and makes it difficult to quantitatively evaluate a material. Given that, we introduced two quantifiable concepts of interphase density (α) and interphase uniformity (β) to standardize the interphase engineering parameters and quantitatively evaluate the structural stability.

As shown in Fig. 2a, N layers of X2 (orange) and Y3 (blue) are randomly stacked along the c-axis direction (with a [110] projection) to create the X2/Y3 interphase structure. The thickness of each X2 or Y3 layers corresponds to the unit cell size (X2: two MO₂ slabs; Y3: three MO₂ slabs). Clearly, the X2/Y3 interphases, highlighted by the orange sheets, form at the boundary of two different types of layers. The interphase density (α) is defined as the ratio of the number of interphases to the total number of X2 and Y3 layers (left panel of Fig. 2a), and the calculation equation is as follows:

1

$\mathrm{\alpha }=\frac{p}{N}$where $p$ represents the number of phase interphase, $N$ represents the total number of X2 and Y3 layers, and the value of α ranges from 0 to 1. Further, Interphase uniformity (β) characterizes the spatial distribution of above interphases (the right in Fig. 2a), and is calculated using a mathematical variance-based formula:

2

$\mathrm{\beta }=\frac{{\sum }_{i=1}^{p+1}{\left(x_i-\frac{1}{\alpha }\right)}^2}{p+1}$where $x_i$ denotes the number of layers in the i-th phase domain, with $x_i$ ≥ 1 (fractional values allowed; see Supplementary Fig. 10 for details). The value of β ranges from 0 to +∞. Like the concept of mathematical variance, the β value reflects the degree of deviation in interphase distribution from the ideal case.

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Fig. 2

The conception of interphase density and interphase uniformity.

a Schematic illustration and equation of interphase density (α) and interphase uniformity (β). b The schematic illustration of constructing interphase engineering. c The distribution of α and β of main reported multiphase-type Na-ion layered oxides, with corresponding values listed in Supplementary Table 2. Source data are provided as a Source Data file.

In general, for a material with X2/Y3 interphase, a high α value indicates a greater number of interphases at the microscopic level and more effective gliding inhibition, thereby achieving better structural stability in electrochemical process. Meanwhile, the interphase uniformity (β) represents the degree of spatial uniformity in interphase distribution. A low β value ensures the homogeneity of the gliding inhibition effect throughout the particle. Conversely, a high β value implies the reduction or absence of interphase in certain regions, which generally results from single-phase clusters, harming the overall stability. It should be noted that the α and β are not independent of each other, and a comprehensive consideration is required. In the ideal case, the α and β are 1 and 0, respectively (Supplementary Fig. 11), and the aforementioned P2 + P3 crystal model actually corresponds to this scenario.

Consequently, rational interphase engineering that can theoretically inhibit the interlayer gliding should meet two pivotal criteria. As illustrated in Fig. 2b, at the macroscopic level, the material must first comprise X2 and Y3 phases to form the X2/Y3 interphase structure. Subsequently, the interphase within the particle should be dense and uniform, which can be quantified by α and β. Supplementary Table 2 and Fig. 2c summarize some current multiphase reports from the interphase perspective. It can be found that systems forming the X2/Y3 interphase (i.e., composed of X2 and Y3 phases) exhibit significantly enhanced cycling stability, whereas those without such interphase show negligible improvement. Moreover, the α and β values from those reports were calculated, and most fall within the region characterized by high α and low β. These findings indirectly validate the correctness of the interphase engineering design.

Structure and electrochemistry of P2/P3-NaMNO

To further validate the feasibility of above interphase engineering strategy, a biphasic P2/P3-NaMNO (Na_0.46Mn_0.9Ni_0.1O₂) positive electrode material with uniform P2/P3 coherent interphases and three comparative samples (P2-NaMNO, P3-NaMNO, and P2/P3-NaMNO-PM) were prepared. The particle morphology and chemical composition of these samples were examined using the scanning electron microscope (SEM), SEM-energy dispersive spectrometer (SEM-EDS), and inductively coupled plasma mass spectrometer (ICP-MS). The comparative samples exhibit similar morphology and identical composition to the P2/P3-NaMNO sample (Supplementary Figs. 12, 13, and Supplementary Tables 3, 4). The high-angle annular dark-field scanning transmission electron microscope (HAADF-STEM) characterization brings to light the microscopic structure. As depicted in Fig. 3a, the atomic-level image in [110]-projection presents a uniform and dense P2/P3 interphase generated by alternating stacking of P2 and P3 phase structures. The average thickness along the c-axis for each phase region is confined to several MO₂ layers (at the unit cell level), closely matching the ideal interphase configuration described above. Interestingly, as shown in Fig. 3b, the interlayer spacing of P2 and P3 structure near the interphase region is remarkably close (5.54 Å vs. 5.53 Å), differing from those in individual P2 (5.51 Å) and P3 (5.55 Å) structures. The selected area electron diffraction (SAED) pattern displaying two sets of diffraction spots confirms the coexistence of P2 and P3 phases^10,33 (Fig. 3c). Moreover, given the accidentality, more microscopic areas selected randomly also been probed, and the similar results indicate a universal phenomenon for such interphase configuration (Supplementary Fig. 14). According to those images and Eqs. (1) and (2), the average values of α and β are calculated to be 0.57 and 0.17, respectively, which are very close to the ideal values (Supplementary Fig. 15). Further, synchrotron high-energy X-ray diffraction (HEXRD) patterns were analyzed using Rietveld refinements to probe the macrocosmic phase composition. The refined HEXRD results of P2/P3-NaMNO unveil a two-phase structure, consisting of P2 and P3 phases in an approximately 1:1 ratio (Fig. 3d). The refined HEXRD results of three comparative samples are also in line with expectations (Supplementary Fig. 16 and Supplementary Table 5). The shoulder/hump near the (002)/(003) peak in the P2/P3-NaMNO and P3-NaMNO samples is attributed to the formation of a hydration phase³⁴, likely caused by air exposure during the shipping and testing process (Supplementary Fig. 17). The intriguing interlayer spacing d values of separated phases, derived from the refined lattice parameters (Supplementary Table 5), were presented in Fig. 3e. Consistent with the STEM results, the d values of the P2 and P3 phases are very close (5.543 Å vs. 5.542 Å) in P2/P3-NaMNO, whereas a notable difference is observed in the comparative samples. This phenomenon may be related to the specific microscopic interphase structure and could serve as a macroscopic indicator for identifying the interphase configuration.

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Fig. 3

The characterization of P2/P3-NaMNO and contrast samples.

a, b HAADF-STEM image of P2/P3-NaMNO and the (b) is a magnified view of the white dotted region in panel a. The blue rectangle and red parallelogram present P2 and P3 phase, respectively. c SAED image of P2/P3-NaMNO; the blue rectangle represents the P2 diffraction spots and orange rectangle parallelogram is P3. d Synchrotron HEXRD refinement results of P2/P3-NaMNO using the Rietveld method. The observed patterns can be well-fitted to pure P2 and P3 phase, with a molar ratio of approximately 1:1. e Interlayer spacing histogram of P2-, P3-, P2/P3-NaMNO, and P2/P3-NaMNO-PM. For the two biphase materials, the interlayer spacing of P2 and P3 phases is displayed separately. Error is derived from the XRD refinement fitting. f The second GCD curves at 0.1 C (1 C = 180 mA g⁻¹) with voltage range of 4-1.8 V. g, Cycle performance over 500 cycles with AC negative electrode at 2 C (1 C = 180 mA g⁻¹). Source data are provided as a Source Data file.

To validate the effectiveness of proposed interphase engineering, the electrochemical performances of P2/P3-NaMNO, pure phase P2-NaMNO, and P3-NaMNO electrodes were evaluated in coin-type cells using metallic Na as negative electrode. The galvanostatic charge/discharge (GCD) (Fig. 3f) and cyclic voltammetry (CV) (Supplementary Fig. 18) curves display similar charge/discharge plateaus and redox peaks, with only slight differences in the specific capacity and redox potential, denoting the same Na-storage mechanism³⁵. Notably, P2/P3-NaMNO electrode demonstrates superior capacity retention of 81.2% after 500 cycles with activated carbon (AC) as negative electrode, outperforming P2-NaMNO (67.1%) and P3-NaMNO (69.2%) (Fig. 3g and Supplementary Fig. 19). Additionally, the P2/P3-NaMNO electrode exhibits the highest rate capability among the three electrodes (Supplementary Fig. 20). The improvement of rate capability could be attributed to the enhancement of Na⁺ diffusion kinetics (Supplementary Fig. 21). Electrochemical impedance spectroscopy (EIS) confirms a superior bulk mass transfer (Warburg impedance, Z_w) for the P2/P3-NaMNO electrode, while similar surface resistance (R_sf) and charge transfer impedance (R_ct) across all three positive electrodes suggest comparable surface properties³⁶ (Supplementary Fig. 22 and Supplementary Table 6). Moreover, the P2/P3-NaMNO retains its crystalline structure and morphology after cycling (Supplementary Figs. 23, 24). To ensure rigor, the physically mixed sample P2/P3-NaMNO-PM, which lacks a distinct P2/P3 interphase structure, was also evaluated. The results show no discernible enhancement in either cycling or rate performance compared to P2-NaMNO and P3-NaMNO electrodes (Supplementary Fig. 25). Therefore, we can reasonably attribute the improved performance and structure stability, observed exclusively in P2/P3-NaMNO, to its unique P2/P3 interphase configuration. Furthermore, seven additional P2/P3 or P2/O3 positive electrode materials, which possess similar X2/Y3 interphase configurations to that of P2/P3-NaMNO, were also prepared. The detailed structural characterization and electrochemical performances of the seven positive electrodes are displayed in Supplementary Fig. 26 and 27. A summary of the pivotal cycling performance is shown in Supplementary Fig. 28. It reveals that all seven X2/Y3 interphase positive electrodes present superior capacity retention after 200 cycles than that of their pure-phase counterparts. These results strongly confirm the universality of X2/Y3 interphase engineering.

The role of interphase in P2/P3-NaMNO

The microstructure evolutions of P2/P3 interphase in the P2/P3-NaMNO material cycled between 1.8 and 4.0 V was observed by atomic-resolution HAADF-STEM technology (Fig. 4). As depicted in Fig. 4a–c, over the whole cycle, the P2/P3 interphase maintained the stability of the adjacent phase structure, thereby preserving the interleaved P2 and P3 phase configuration. In regions distant from the interphase, a slight phase transition from P3 into O′3 or/and O3 phases³⁷ (indicated by deep red parallelogram) occurs at the discharged state. For a more comprehensive insight, Fig. 4d–f illustrates the structural changes near the interphase layer region in greater detail. From the perspective of MO₂ interlayer spacing, it can be found that the spacing adjacent to the interphase layer consistently remains at ~5.53 Å (both P2 and P3 phase), whereas the spacing in the region far from the interphase undergoes more significant changes throughout the cycle. Especially for P3 phase, when charged to 4 V, the interlayer spacing shows an increasing tendency with distance from the interphase, reaching 6.08 Å, indicating a transformation toward P′3. Conversely, during discharge, the spacing decreases to 5.29 Å, corresponding to the O′3 or O3 phase (the reason for this phase transition are detailly discussed in later discussion section).

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Fig. 4

The micro-structural evolution of P2/P3 interphase.

HAADF-STEM images with [110]-projection at charging cut-off voltage 4 V (a, d), discharging middle voltage 2.25 V (b, e), and discharging cut-off voltage 1.8 V (c, f). The (a–c) indicate the persistent presence of the P2 + P3 interphase and P2/P3 interleaving structure over the whole cycle; the (e, f) illustrate changes in interlayer spacing near the interphase layer. g The theoretical angles (indicated as θ) between adjacent layers based on the M atoms for P2, P3, and O3 phase. h, i The change of θ value near the interphase at 1.8 V for different interphase distribution situations. Comparing with the theoretical θ values, it can be found that the P2 phase was almost unchanged, and the remote P3 phase region was transformed into O3 phase with a deviated θ value, while that can be avoided for a more rational interphase distribution (i).

Moreover, the lateral shift of MO₂ layers provided a more intuitive representation of the gliding behaviors at the interphase region. Figure 4g displays primary diffraction angles between the MO₂ layers in various phase structures, from which the angles of P2, P3, and O3 are 90°, 80°, and 73°, respectively. The glide of MO₂ layer can be reflected through the change in the corresponding angle. Figure 4h, I show such angle changes in two selected regions with different interphase configurations at the discharged state (1.8 V). In the former, which only contains one interphase, the P2 and P3 layers nearby the interphase retain their initial angles, while the angles of more distant P3 layers gradually decrease toward 73° (indicative of the O3 phase). In the later, with two interphases arranged similarly to the aforementioned P2/P3 model, the P3 layer angles remain almost unchanged, close to the theoretical values, indicating that no gliding occurred. Consequently, based on above results, it can be confirmed that the P2/P3 interphase remains immobile and simultaneously prevents the gliding of adjacent MO₂ layers or phase transitions within a certain region during the electrochemical process, which is the origin of the performance improvement. Notably, when beyond the effective range of the interphase, phase transition occurs, accounting for the emergence of the P′3, O′3, and O3 phases. This also indirectly confirms the role of interphase in inhibiting gliding of both itself and the adjacent MO₂ layers.

At the macroscopic scale, in-situ synchrotron XRD was employed to monitor the structural evolution of P2/P3-NaMNO. An illustrative set of comparative data is presented in Fig. 5, which consists of the evolutionary contour plots of P2/P3-NaMNO (Fig. 5a) and P2/P3-NaMNO-PM (Fig. 5b) samples and the corresponding GCD curves. In the charge and discharge process, P2/P3-NaMNO maintained the initial P2 + P3 phase structure without the appearance of new diffraction peaks, suggesting a favorable structural stability. By contrast, in P2/P3-NaMNO-PM, the P3 phase transforms into the P′3 phase (highlighted by the black dashed circle) at the end of charging and the beginning of discharging^38,39, confirming the effective suppression of phase transition in P2/P3-NaMNO. The in-situ XRD results of two pure phase electrodes also indirectly support this conclusion (Supplementary Fig. 29). Further, Fig. 5c, d depict the c-axis evolution of the P2 and P3 phases in both P2/P3-NaMNO and P2/P3-NaMNO-PM samples, based on the in-situ results. Comparative analysis reveals that the c-axis of the P2 phase exhibits a similar trend in both samples. However, the c-axis of the P3 phase in P2/P3-NaMNO-PM undergoes a sharp change during the P′3 phase transition, while that in P2/P3-NaMNO changes more gradually. These results suggest that the P2 and P3 phases can mutually reinforce each other’s structural stability through the P2/P3 interphase, as previously discussed. Moreover, the change in the c-axis of P2/P3-NaMNO is less than 1% for both the P2 and P3 phases, indicating a low-strain property. Overall, the in-situ synchrotron XRD characterization further confirms that the P2/P3-NaMNO sample exhibits a better structural stability, which should be attributed to the distinctive P2/P3 interphase configuration. Hence, it can be concluded that the interphase engineering can effectively inhibit the gliding of MO₂ layers and associated phase transitions.

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Fig. 5

The structural evolution in macro-level.

a, b Contour plot of in-situ synchrotron XRD patterns of P2/P3-NaMNO (a) and P2/P3-NaMNO-PM (b) during the first cycle. The right figure shows the corresponding charge-discharge curves. The Red to blue coloration represents decreasing peak intensity. The white dotted line marks the region of phase transition, and the black dotted line ellipse marks the signal of new phase. c, d A comparison of the lattice parameters of c-axis between P2/P3-NaMNO and P2/P3-NaMNO-PM; c corresponds the individual P2 phase and (d) corresponds the individual P3 phase. The lattice parameters were obtained by the corresponding in-situ synchrotron XRD refined data. Source data are provided as a Source Data file.

Discussion

For the interphase engineering, adopting the “X2 + Y3” (X2/Y3) phase combination pattern is the basis for constructing the gliding-inhibition interphases. The underlying reason for the “X2 + Y3” principle may arise from the fact that X2-type and Y3-type phases cannot interconvert via electrochemical (de)sodiation, which involves breakage or reformation of M-O bonds^13,27, thereby ensuring the structural stability of the interphase. Moreover, for Na insertion/extraction, the X2- and Y3-type phases inherently exhibit contradictory gliding behaviors of MO₂ layers (Supplementary Fig. 30), thus, the shared interphase tends to remain motionless. It is worth noting that, although the resultant material from the interphase engineering also belongs to a multiphase structure at the macroscopic scale, its core lies on the microscopic phase configuration associated with two specific phases. In contrast, widely reported multiphase strategies primarily focus on the integration of multiple phases themselves based on the principle of complementary advantages, without intentionally designing phase types or considering microscopic interactions. However, the significant inconsistency in electrochemical performance observed in many multiphase studies suggests the limitations of such principle^10,40,41, and the interphase engineering proposed in our work may serve as further direction for multiphase materials. Additionally, the interphase discussed in this work specifically refers to that generated by two layered structures along the c-axis, and the requirement for X2/Y3 type only applies only in this context. Other types of interphases, such as “three-phase” interphase, “tunnel-layer” interphase, and “a-axis direction” interphases, are beyond the scope of this article.

Despite the evidence clearly elucidates that the MO₂ layer gliding could be inhibited at a designed X2/Y3 interphase, P2/P3-NaMNO still displays slight phase transitions, especially during deep charge and discharge states. This gap between ideal and reality mainly stems from the fact that the quantity and distribution of interphases (α and β) in real materials may not align perfectly with the model due to the inevitable disordered distribution of single-phases clusters. This observation leads to some MO₂ layers being too far from the interphase to be affected, and thus, sliding and phase transition still occur, as demonstrated in the regions of P2/P3-NaMNO distant from the interphase. While achieving the perfect interphase configuration is challenging, we can still strive for an optimal interphase distribution (higher α and lower β). Additionally, these two parameters of α and β collectively determine the overall effectiveness in inhibiting gliding and phase transitions, and thus predict the electrochemical performances. It can also explain the discrepancy in cycling improvement among the multiphase reports, even when they possess the same X2/Y3 interphase.

Lastly, in reality, the main challenge of interphase engineering lies in synthesizing materials with high-quality X2/Y3 interphase configuration. Hence, we offer some perspectives on how to achieve the eligible X2/Y3 interphase in a layered oxide material (Supplementary Fig. 31). First, achieving uniform pre-mixing of various elements is crucial to prevent the clustering of specific phases and to ensure high interphase density and uniformity. In addition to the coprecipitation method, both the sol-gel and solid-state methods are applicable; however, the solid-state method requires a preliminary step, such as ball milling, to ensure thorough mixing of the raw materials (as demonstrated in Supplementary Figs. 32–34). Secondly, a moderate Na content is required to achieve a multiphase structure, as high Na content tends to form the O3 phase, while low Na content favors the formation of P2 or P3 phases. Lastly, the calcination temperature also affects the formation of different phase types, as shown by the variable temperature XRD results (Supplementary Fig. 35), where lower temperatures favor Y3-type phases and higher temperatures favor the X2-type group phase.

In summary, our study demonstrates that interphase engineering can effectively address MO₂ layers gliding issues associated with phase transition in layered oxide materials. Exemplifying the Na-based layered oxides, our findings revealed that the gliding of MO₂ layers near a X2/Y3 coherent interphase (a shared MO₂ layer) can be fundamentally inhibited through an energetically driven gliding-inhibition mechanism. This suggests that the interphase MO₂ layer tends to remain immobile during charge/discharge process to minimize the overall system energy. Building on this understanding, we designed and synthesized a multiphase positive electrode material, P2/P3-NaMNO, with an eligible P2/P3 interphase configuration. This material certainly exhibited enhanced cycling stability and rate performance compared to pure-phase counterparts. Detailed structural analyses, including HAADF-STEM and synchrotron XRD, confirmed the role of the P2/P3 interphase layer in restricting MO₂ layer gliding. Moreover, the broad applicability of interphase engineering also has been demonstrated by several other X2/Y3 interphase materials. In addition, we provide complementary discussions on the interphase and offer perspectives on material synthesis. In a broader context, our findings pave the way for addressing gliding issues and advancing structural stability in the electrochemical processes of sustainable battery layered positive electrode materials.

Methods

Materials

NiSO₄·6H₂O (≥98%), MnSO₄·H₂O ( ≥ 98%), NaOH (≥99%), disodium citrate (≥98%), metallic sodium (≥99%), and activated carbon (AC) were purchased from Aladdin. Na₂CO₃ (≥ 99.9%) was purchased from Alfa Aesar. All chemicals were used as received. All the solutions were prepared with deionized water. [Mn_0.9Ni_0.1](OH)₂ precursor was firstly prepared by a coprecipitation method. Specifically, NaOH was used as precipitating agent, and disodium citrate was used as chelating agent. Stoichiometric amounts of NiSO₄·6H₂O and MnSO₄·H₂O with a total M concentration of 2.0 mol L⁻¹ were dissolved in deionized water to form a homogeneous solution. Such solution was then slowly pumped into a tank reactor purged with N₂. Simultaneously, NaOH and disodium citrate solutions were separately pumped into the reactor under carefully controlled pH (≈11.0), temperature (50 °C), and stirring speed (1000 rpm). The obtained precipitate was washed, filtered, and vacuum dried at 80 °C for 24 h. To obtain the transition metal oxides, the mixture of [Mn_0.9Ni_0.1](OH)₂ precursor and Na₂CO₃ (with a 5% excess) was calcined for 12 h in air. The calcination temperatures for P2/P3-NaMNO, P3-NaMNO, and P2-NaMNO material were 850, 700, and 900 °C, respectively. P2/P3-NaMNO-PM materials were synthesized through mixing the pure P2-NaMNO and P3-NaMNO with mass ratio of 1:1, followed by calcination at 850 °C for 12 h in air. The seven additional samples were synthesized using the same coprecipitation method.

Electrochemical tests

Active materials (80 wt.%), carbon black (10 wt.%), and polyvinylidene difluoride (10 wt.%) were thoroughly mixed in N-methyl pyrrolidone solvent (Aladdin, 99.9%). Then the obtained homogeneous slurry was pasted on an aluminum foil (15μm, 99.6%, Hubei Fengsheng Electronic Technology Co., Ltd) to fabricate the electrodes and dried at 120 °C overnight in a vacuum oven. The slurry for the AC negative electrode consisted of 70 wt% commercial AC, 20 wt% acetylene black, and 10 wt% polytetrafluoroethylene (PTFE). All electrodes were cut into circular pieces with a diameter of 12 mm. The average mass loading of positive electrode material was 4–5 mg cm⁻². CR2025 coin cells (containing positive and negative stainless steel shells, a piece of single coated positive electrode, negative electrode, separator, electrolyte, one round stainless steel (15 mm ×0.15 mm), and one stainless steel spring) were assembled in an Ar-filled glovebox with less than 0.01 ppm H₂O and O₂. The negative electrode was made up of fresh sodium metal (~0.5 mm in thickness, 14 mm in diameter) or AC electrode (~2 mm in thickness, 14 mm in diameter). The separator was made up of glass fiber separator (Whatman GF/D, 16 mm in diameter, 675μm in thickness). The electrolyte consisted of 1 M NaPF₆ in diglyme (Duoduo Chemical Technology Co., Ltd). The calculations of specific capacity and specific current are based on the active positive electrode material content. All electrochemical measurements were conducted on coin cells at room temperature (25 ± 2 °C). Galvanostatic charge–discharge tests were carried out between 1.8 and 4.0 V using a Land BT2000 battery testing system (Wuhan, China) under constant current conditions. The same testing system was used for galvanostatic intermittent titration technique (GITT), applying a 10-minute current pulse followed by a 60 min rest period. Cyclic voltammetry (CV) and electrochemical impedance spectroscopy (EIS) were performed on a CHI760D electrochemical workstation (Chenhua, China). CV scans were recorded at 0.1 mV s⁻¹ within 1.8–4.0 V. EIS was conducted with a 5 mV amplitude over a frequency range of 10⁻² to 10⁵Hz at open-circuit voltage, and data were fitted using Zview software. Each electrochemical test was repeated on at least three independently assembled cells, and the results presented correspond to a representative cell exhibiting typical or median performance.

HAADF-STEM and SEM measurement

The as-prepared ingots were examined by focused ion beam (FIB) milling (Helios Nanolab G3 UC, FEI) for structural characterization. Double Cs-corrected transmission electron microscopy (Titan Themis G2 60–300, FEI) was used for STEM test with HAADF techniques. All HAADF images in this paper were Fourier-filtered to minimize the effect of the contrast noise, which did not have any effect on our measurement results. The field-emission scanning electron microscope (SEM) and SEM-energy dispersive spectrometer (SEM-EDS) were collected on JEOL-7100F. The ex-situ samples at different charge/discharge states were prepared via charging and discharging in the coin cells beforehand. The coin cells were disassembled in an Ar-filled glove box with water and oxygen content less than 0.01 ppm. The extracted electrodes were thoroughly washed with dimethyl carbonate (Aladdin, 99.9%) and dried at 100 °C to remove solvents in the glove box. For the transfer of these samples, the corresponding sample holder was first brought into the glove box for sample preparation, then rapidly transferred to the instrument for testing. The entire transfer process takes less than 3 min.

X-ray diffraction measurements

High-energy X-ray diffraction (HEXRD) measurements were carried out at beamline 11-ID-C of the Advanced Photon Source (APS), Argonne National Laboratory. Samples were hermetically sealed in an argon-filled glovebox prior to transport to avoid air exposure. A focused X-ray beam (λ = 0.1173 Å, 0.2 × 0.2 mm²) was employed, and two-dimensional diffraction patterns were collected in transmission (Laue) geometry using a Perkin-Elmer area detector positioned 1800 mm from the sample. For in situ experiments, custom-designed CR2032 coin cells with a 3 mm aperture were used to enable X-ray transmission. The aperture was sealed with Kapton tape to maintain an inert atmosphere. Time-resolved diffraction patterns were acquired every 10 min under operating conditions using the same beamline setup. The laboratory XRD data were measured by a Bruker D8 Advance X-ray diffractometer with Cu Kα radiation (λ = 1.5418 Å) in the scan range (2θ) of 10°–80°. Structural analysis was performed via Rietveld refinement using the GSAS-II software package.

Theoretical calculations

All the first-principles calculations were performed within the density-functional theory (DFT) framework⁴², and were carried out in the CASTEP module of Materials Studio 2018. The exchange-correlation potential was treated using Perdew-Burke-Ernzerhof (PBE) functional in generalized gradient approximation (GGA) method⁴³. The value of the K point was set to 2 × 5 × 1, and Norm conserving pseudopotentials were employed to solve the electronic structure of the whole system. The GGA + U method was used with U-J value of 3.9 and 6.2 for Mn 3 d and Ni 3 d electrons, respectively. A cutoff energy of 500 eV was used, and the energy convergence criterion for ionic relaxation was set to <10⁻⁵eV per atom. During optimization process, the maximum force was set to 0.05 eV Å⁻¹, and the maximum displacement was limited to 0.002 Å. Furthermore, the correction of van der Waals forces and non-bonded interactions were carried out by the DFT-D3 model developed by Professor Grimme’s group⁴⁴. For the calculation of the formation energy, the following formula is followed:${\mathrm{\Delta }\mathrm{E}}_{\mathrm{Formation\; Energy}}={\mathrm{E}}_{\mathrm{Na}x\mathrm{Mn}y\mathrm{Ni}z\mathrm{O}2}-{x\mathrm{E}}_{\mathrm{Na}}-{y\mathrm{E}}_{\mathrm{Mn}}-{z\mathrm{E}}_{\mathrm{Ni}}-2{\mathrm{E}}_{\mathrm{O}}$where, E_NaxMnyNizO2 is the total energy (single point energy) of the structure with different components, 0 <x < 1, y = 0.9, and z = 0.1. E_Na is energy of one Na atom. E_Mn is energy of one Mn atom. E_Ni is energy of one Ni atom. E_O is half energy of oxygen molecules.

Acknowledgements

This work was supported by the National Key Research and Development Program of China (2023YFB2406100) and National Natural Science Foundation of China (Grant No. 52172234).

Author contributions

Y.Y., J.L., and X.Z. conceived the idea and designed the experiments. proposed and planned the project. X.Z., C.Y., and Y.Z. carried out all sample synthesis and electrochemical measurements. T.L. and X.Z. performed the synchrotron HEXRD, in-situ synchrotron XRD, and data analysis. X.Z. and X.P. carried out the ex-situ HAADF-STEM, SAED, and data analysis. X.L. conducted the DFT theoretical calculations. X.Z., Y.Y., L. W., and J.L. wrote and revised the manuscript. All authors discussed the results and commented on the manuscript.

Peer review

Peer review information

Nature Communications thanks the anonymous reviewer(s) for their contribution to the peer review of this work. [A peer review file is available].

Data availability

The source data generated in this study are provided in the Source Data file.  are provided with this paper.

Competing interests

The authors declare no competing interests.

Supplementary information

The online version contains supplementary material available at https://doi.org/10.1038/s41467-025-61065-w.

Publisher’s note Springer Nature remains neutral with regard to jurisdictional claims in published maps and institutional affiliations.

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