Potential-dependent interfacial specific adsorption accelerates charge transfer in sodium-ion batteries

Domain: chemistry
Published: 2026-02-17
Editorial outcome: minor_revision

Abstract
Fast-charging capabilities of sodium-ion batteries have emerged as a pivotal objective within the energy storage fields. Sodium layered P2-type oxide positive electrodes are considered promising for fast charging due to their inherent fast Na+ mobility. However, their electrochemical polarization and interfacial charge transfer especially at high state of charge are limiting factors in quick kinetic response for large current. Herein, we demonstrate that a typical P2-type positive electrode (Na0.7Ni0.27Mn0.53Cu0.04Fe0.08Ti0.08O2) achieves high-rate capacities through avoiding octahedral stacking faults, maintaining lattice oxygen activity and controlling anion-specific adsorption. The intermediate Z-phase intergrowth structure mitigates kinetic polarization and thermodynamic hysteresis by simultaneously suppressing the unfavorable phase evolution from P2-type to O2-type and irreversible oxygen redox. The potential-dependent competitive adsorption mechanism between anions and solvent molecules is revealed within the inner Helmholtz plane, where optimized anion-specific adsorption elevates potential difference between electrodes and inner Helmholtz plane, accelerating charge transfer across the electrode/electrolyte interface. Furthermore, the F-rich cathode/electrolyte interphase generated from inner Helmholtz plane mitigates transition metal dissolution and surface lattice collapse for stable long-term cycling. This study highlights the synergistic coupling interaction between bulk phase stability and interfacial environment optimization in ensuring fast Na+/charge transport kinetics for sodium-ion batteries.

Introduction
Intermittent energy sources have a growing demand for energy storage, in which case sodium-ion batteries (SIBs) have significant potential to be applied to large-scale energy storage due to their low cost1,2. In particular, P2-type layered oxides with the open Na+ diffusion paths are considered a promising candidate for fast-charging SIBs3,4. Disappointingly, their energy density and power density are difficult to balance under limited 50% Na+ (de)intercalation (2.0−4.0V vs. Na|Na+)5. To tackle this issue, some researchers have chosen to extend the state of charge (SoC), mitigate the large-volume phase transition, or activate additional anionic redox reactions (ARRs), inevitably resulting in obvious potential hysteresis and irreversible structural degradation during cycling6,7.

Such phase transition from the prismatic (P) type to octahedral (O) type and oxygen-related ARRs occur at deep SoC, where the P2-type structure is confronted with undesired electrochemical polarizations during excessive deintercalation of Na+8. Specifically, the weakened shielding effect of Na+ on the electrostatic repulsion between O−O slabs causes the transition from P- to O-type stackings in the deep sodium-depleted state, and the resultant contracted Na layer within O-type phase impedes Na+ diffusion kinetics7,9. Moreover, unhybridized O 2p orbitals are placed above transition metal (TM) orbitals at high potential, where the extraction of electrons from the newly generated O 2p non-bonding states near the Fermi level (Ef) is activated for charge compensation, accompanied by obvious potential hysteresis, which seriously degrades the rate performance10,11. Critically, phase transitions and ARRs are coupled with each other under high potential, and there remains controversy about the effect of the generated deleterious electrochemical polarizations on the rate performance of P2-type positive electrodes11,12.

The activity of lattice oxygen is recognized to initiate on the positive electrode surface and react easily with electrolyte in the high SoC, leading to cathode/electrolyte interphase (CEI) generation, which also impedes electrochemical polarizations13,14. Most researchers have neglected the fact that charge transfer process across the electrode/electrolyte interface (EEI), including Na+/electron transport, is deemed to be the rate-determining step for fast-charging15–17. In general, the interfacial environment, especially the inner Helmholtz plane (IHP) consisting of solvent molecules and anions (mainly ClO4− and PF6−), is the preferred site where the charge transfer reactions primarily take place18–22. The deep desodiated positive electrodes accumulate significant positive charges on the surface at high SoC, which generates a stronger electrostatic force on anions to form the interfacial specific adsorption within the IHP23,24. The evolution of anion-specific adsorption, as well as whether it facilitates or inhibits charge transfer reactions, is critical for comprehensively understanding ion/charge interfacial transport properties and developing fast-charging positive electrodes for SIBs.

Here, we focus on the typical Ni/Mn-based P2-type positive electrode material Na0.7Ni0.27Mn0.53Cu0.04Fe0.08Ti0.08O2 (denoted as NMCFT) and the corresponding counterpart Na0.7Ni0.35Mn0.65O2 (NM). By analyzing SoC-dependent bulk structure and charge compensation, we elucidate that the Z-phase intergrowth structure, which is free from completely converting to O2 phase at deep SoC, contributes to reducing both kinetic polarization by reversible P2−Z phase transition and thermodynamic hysteresis through triggering TM−O hybridization-dominated ARRs. The boundary diffusion behavior associated with the thermodynamic barrier renders the electrodes partially capacitive, thus accelerating anion-specific adsorption in the IHP. As the positive electrode potential rises and the anion coverage on the electrode surface increases, the facilitating role of the increasing potential difference between positive electrode and IHP is diminished by the shielding effect of the anions themselves on the charge transfer. By controlling anion-specific adsorption and combining the mild P2−Z phase transition with reversible ARRs, NMCFT accelerates charge transfer and Na+ diffusion under deep SoC to achieve high-rate capacities under long-term cycling. For application, NMCFT delivers good cycling stability with approximately 80% capacity retention over 300 cycles in pouch-type full cells with specific energy of 156 Wh kg−1.

Results
Comparison of high-rate electrochemistry
The X-ray diffraction (XRD) refinements (Fig. 1a, b and Supplementary Tables 1 and 2) reveal that both NM and NMCFT positive electrodes belong to the typical P63/mmc space group. The Na layer spacing dO−Na−O of the P2-type structure decreases from 3.572 Å to 3.503 Å after substitution, and correspondingly, the weighted average ionic potential of TM decreased from 59.202 to 57.93225. The Inductively Coupled Plasma Mass Spectrometer (ICP-MS) results (Supplementary Tables 3 and 4) indicate that the prepared NM and NMCFT positive electrodes conform to the design. The high-resolution transmission electron microscopy (HR-TEM, Supplementary Fig. 1) and the scanning transmission electron microscopy (STEM, Fig. 1c and Supplementary Fig. 2) images of NMCFT demonstrate that the average interslab distance is 5.56 Å, exhibiting typical ABBA stacking along the [001] zone axis, while the neighboring distance of TM atoms measures 2.91 Å in the [010] zone axis. And X-ray absorption near edge structure results (XANES, Supplementary Fig. 3) elucidate the valence of TMs in NMCFT, consistent with the targeted substitution strategy.

When cycled at 2.2−4.0 V (vs Na|Na+) in half cells (Supplementary Fig. 4), both NM and NMCFT exhibit reversible discharge capacity of ~84 mAh g−1 at 0.1 C (1 C = 180 mA g−1), although having stable capacity retention (99.1% for NM and 100.6% for NMCFT over 100 cycles) at C/3. When the upper cutoff potential (Eup) is increased to 4.4 V, the discharge capacity of NMCFT increases to 133.2 mAh g−1 at 0.1 C (Supplementary Fig. 5, slightly lower than that of NM (143.6 mAh g−1), and demonstrates 86.9% capacity retention over 100 cycles at C/3 (Fig. 1d, versus 55.7% of NM). For NM, the ~21.3% capacity degradation observed in the first 20 cycles at high-SoC directly results in earlier failure at higher operating potential compared to NMCFT. Through practical fast-charging (3 C) and slow-discharging (C/3) tests (Fig. 1e and Supplementary Fig. 6), NMCFT maintains a reversible specific capacity of 120.7 mAh g−1 (versus 91.5 mAh g−1 of NM) and 87.0% capacity retention over 100 cycles (versus 86.0% of NM), exceeding NM by over 25 mAh g−1 from the start of cycling. Under larger current density, NMCFT delivers good cycling performance with 85.67% capacity retention over 300 cycles at 5 C (Supplementary Fig. 7) because of the fast-charging ability at high potential, while NM delivers a severe decay of rate performance (Supplementary Fig. 8, maintaining 36.7% at 5 C versus 0.1 C) and fails to cycle under high-rate condition. Quantitative analysis at different SoCs reveals that the high-rate capacity degradation of NM entirely originates from hindered Na extraction at high SoC (Fig. 1f, g). Further desodiation from NMCFT remains with fast kinetics after 50% Na+ extraction, whereas NM is severely hampered by the factors from the bulk and interface, leading to a significant deterioration in high-rate electrochemical performance.

Origin of overpotential at high SoC
From in-situ XRD results (Fig. 2a and Supplementary Fig. 9), NMCFT exhibits the reversible P2−Z phase transition in the deep SoC, which is asynchronous with the non-equilibrium potential at charging to 4.0 V. Similarly, previous studies and ex-situ XRD of NM (Supplementary Fig. 10) indicate that the non-equilibrium potential emerges prior to the O-type stackings26. In Supplementary Fig. 11, the relaxations and calculated Na+ diffusion coefficients at high SoC were revealed by the galvanostatic intermittent titration technique (GITT) measurements. The smaller potential drops and the higher average Na+ diffusion coefficient suggest that NM has “faster” high-SoC Na+ diffusion kinetics than NMCFT, despite the latter demonstrating better rate performance. The Fast-GITT measurements with lengthened relaxations and shortened pulses explain this paradox (Fig. 2b, c). Before the 4.2 V plateau, the relaxation of NM first undergoes a large potential drop of 0.4 V (4.2−3.8 V), and then, reduces to 0.08 V drops (4.2−4.12 V), which occurs preferentially (within 20 hours) and progressively occupies the limited relaxation time, while the large potential drops are concealed. In the long-period (120 hours) relaxation, the complex relaxation behaviors are further distinguished (Supplementary Fig. 12), where this prior relaxation is attributed to the contracted Na layer spacing in the O2-type phase, obstructing the Na+ diffusion. In contrast, NMCFT exhibits a consistent relaxation pattern above 4.0 V SoC (Fig. 2c). No discernible kinetic obstruction due to the P2−Z homogeneous reaction enables NMCFT to proceed directly into the sluggish relaxation of thermodynamic hysteresis27.

The origin of the high-potential thermodynamic hysteresis of NMCFT was investigated by hard X-ray absorption spectroscopy (XAS) in Fig. 2d and Supplementary Fig. 13. The Cu−O shell peak shifts leftward (~0.032 Å) and increases in intensity at 4.4 V SoC, indicating the hybridization between Cu 3d and O 2p orbitals. The intensities of the Cu−TM and Fe−TM shell peaks dropped sharply after charging 4.4 V SoC, suggesting that the coordination environments of Cu and Fe suffered a severe distortion. In the discharged state, the Cu−TM and Fe−TM shell peaks reverted to the initial intensities and positions, when Cu(II) and Fe(III) reversibly migrate back to the octahedral sites of TMO2 slabs due to the disappearance of the tetrahedral sites related to the O-type stackings12. According to the wavelet transform-extended X-ray absorption fine structure (WT-EXAFS) results (Fig. 2e, f and Supplementary Fig. 14)28, the leftward wave vectors (k) of Cu−TM (~0.891 Å−1) and Fe−TM (~0.609 Å−1) shell peaks at 4.4 V SoC indicate that the mass numbers of the elements coordinated to Cu and Fe become smaller, further verifying the interlayer reversible Cu and Fe migration.

In the TFY mode of O K-edge soft X-ray absorption spectroscopy (sXAS), the evolution of TM 3d−O 2p orbital hybridization is investigated by comparing the pre-edges (Fig. 2g, h). For NM, the Ni 3d t2g (~529.5 eV)−O 2p orbital lack hybridization due to low valence states of Ni(III) in the low-potential region and the pre-plateau stage (2#−4# in Fig. 2g). Until the post-plateau stage and 4.4 V SoC (5# in Fig. 2g), the hybridization level rises, where the shoulder peak (~530.5 eV) representing ARR appears. In contrast to NM, the d−p hybridization of NMCFT is enhanced at lower potential (3# in Fig. 2h) accompanied by the earlier appearance of shoulder peaks. The Cu 3d and Fe 3d orbitals strengthen the d−p orbital hybridization, causing oxygen to be oxidized at lower potential via TM reductive coupling, rather than from oxygen lone-pair states (non-bonding O 2p orbitals) as in NM10. Analysis of the TEY mode of O K-edges sXAS (Supplementary Fig. 15) and O 1s X-ray photoelectron spectroscopy (XPS) results (Supplementary Fig. 16) similarly reveals the different ARR pathways for NM and NMCFT, which cause the different thermodynamic hysteresis at high potential27,29. The large ARR barrier in NM evidenced by sXAS pre-edge broadening (Fig. 2g) and XPS peak area widening (Supplementary Fig. 16), induces pronounced overpotential, explaining the steep potential drop in its third relaxation stage (Supplementary Fig. 12), whereas the smaller ARR barrier gives a mild thermodynamic hysteresis to NMCFT, which corresponds to the smooth relaxation behavior.

Together the combined analysis of structural evolution and charge compensation, the deleterious polarizations under large current in the deep SoC are significantly mitigated in the NMCFT. The ordered Na+/vacancy arrangement at certain sodium contents (x = 1/2 and 1/3) induces the polarization potentials at 3.4 V and 3.7 V to overcome the Na+ diffusion barriers (Inset I in Fig. 3a)4,30. Incorporating TM elements with different Fermi levels disorganizes the ordered Na+/vacancy arrangement so that the low-potential region becomes a solid-solution reaction with faster kinetics (Inset I in Fig. 3b). Upon charging to 4.0 V SoC, the oxygen lone-pair state participates in charge compensation (Inset II in Fig. 3a), resulting in a drastic thermodynamic hysteresis10. Broadened energy levels of the d−p hybridized orbitals lower the activation barrier for TM−O reductive coupling, facilitating oxygen oxidation at reduced potential (Inset II in Fig. 3b). With 0.4 Na+ extracted, P-type stackings begin to collapse into O-type stackings (Inset III in Fig. 3a), where the hindered Na+ diffusion triggers the kinetic polarization26. The P2−Z homogeneous reaction (Inset III in Fig. 3b) with little kinetic polarization provides a broader Na+ diffusion path to achieve rapid extraction of the remaining Na+31. From a bulk perspective, the combination of moderate ARR and reversible phase transition contributes to the enhanced rate capability of NMCFT at high SoC.

Bridging interfacial environments and charge transfer
To investigate the charge transfer process on the electrode surface, galvanostatic intermittent titration technique-electrochemical impedance spectroscopy (GEIS) tests were performed using three-electrode cells (Fig. 4a, b and Supplementary Fig. 17). The galvanostatic charge/discharge curves (GCDs) of GEIS for the NM and NMCFT positive electrodes are shown in Supplementary Fig. 18, which are divided into different regions according to the GEIS results. Notice that there are anomalous inductive loops in the low-potential GEIS results. For NM, the inductive loops disappeared at slope II (Fig. 4c, thermodynamic barrier-induced overpotential), accompanied by diffusion behavior that approaches capacitance (a vertical line close to 90°). Within the subsequent high-potential plateau III (Fig. 4d), the substantial capacitive arcs in the mid-frequency region indicate the appearance of massive charge transfer resistance (Rct). However, the size of inductive loop in GEIS of NMCFT reaches a maximum at 4.2 V (Fig. 4e), then gradually diminishes until it coincides with the capacitive arc at 4.3 V (Fig. 4f). According to the timescale of reactions, the mid-frequency EIS plots characterize the charge dynamics32, suggesting that the interfacial environment transforms between inductive and capacitive due to variations in charge transfer process. Notably, this phenomenon is absent in two-electrode cells, as discussed in Supplementary Fig. 19.

In Supplementary Fig. 20, each Nyquist plot of the NM and NMCFT positive electrodes during the first charging was fitted using the corresponding equivalent circuit models, with quantitative evolutions presented in Supplementary Fig. 21 and Supplementary Tables 8–11. The inductive loop in quadrant IV is modeled as a series circuit of resistance and inductance parallel to the main circuit, which reduces the total circuit impedance. Low-frequency EIS spectra can be represented by the Warburg impedance (ZW) and the finite boundary diffusion impedance (ZT) in series33,34. Adjustment of the values for these two impedances in the W/T composite diffusion model unifies low-frequency EIS signals with angle >45°. Notice that the inductive loops and large ZT coexist. The ZT-dominated diffusion behavior forms an impermeable boundary for Na+ at the electrode surface, impeding ion and electron shuttling across the interface, resembling capacitance. The capacitive surface exerts Coulombic forces on polar groups in the electrolyte35, thereby driving anion aggregation toward the surface and subsequent specific adsorption. The anion-specific adsorption within the IHP on the electrode surface is the origin of the inductive loop. At ~4.0 V in Slope II, the NM exhibits ZT-dominated diffusion that is nearly entirely capacitive and exerts significant long-range electrostatic interactions on anions, resulting in intense anion-specific adsorption. The NMCFT consistently exhibits W/T composite diffusion behavior, whose force on anions is relatively moderate. As detailed in Supplementary Note 1, the inductive circuit (parallel) reduces total impedance, while the capacitive circuit (series) increases total impedance. The mild and persistent adsorption behavior maintains lower interfacial impedance of the positive electrode system, corresponding to better rate performance. For NM and NMCFT, the difference is most pronounced in the high-potential region. A more detailed analysis is expanded in Supplementary Note 1.

To further explore the relationships between interfacial environments, charge transfer and diffusion behavior, potentiostatic intermittent titration technique-EIS (PEIS) tests (Fig. 2g−j) were performed. The low-frequency −Z'' at plateau I and II of NM was nearly zero, yet rose abruptly to hundreds of ohms at 3.4 and 3.8 V SoC (Fig. 2g), revealing negligible Na+ diffusion kinetic hindrance at low-potential plateaus but stagnant ion transport at slope stages. The NM material manifested segregated insertion and capacitance properties, where the purely capacitive interface without Faradaic process drives rapid aggregation of anions via unshielded Coulombic forces. Notably, the inductive signal disappeared at 3.9 V SoC and a capacitive interfacial environment emerged, forcing Na+ to overcome elevated Rct during further extraction at plateau III (P2−O2 phase transition). In contrast, NMCFT maintained insertion and dispersed capacitive properties (Fig. 2i), permitting gentle anion-specific adsorption within the IHP without aggregation, as evidenced by sustained inductive signals between 3.0−4.3 V SoC region. The prolonged anion-specific adsorption process delays Rct buildup at high potential by controlling the interfacial environments, allowing NMCFT to undergo deep desodiation with minimal interfacial obstruction at deep SoC, thereby achieving good rate performance.

Evolution of interfacial anion-specific adsorption
The adsorption models of solvent molecules and anions on the surface of P2-type layered positive electrode material were developed in Fig. 5a and Supplementary Fig. 22 and the atomic coordinates of each model are shown in Supplementary Data. As shown in Fig. 5b, the anions exhibit more negative adsorption energies on the surface of positive electrode compared to the solvent molecules, suggesting that anions dominate the evolution of the interfacial environment. The plane-averaged charge density difference in the Density Functional Theory (DFT)-optimized model reveals the charge density distribution of the interface with different adsorbed groups (Fig. 5c). When an anion is adsorbed, charge transfer from the anion to superficial Ni atoms at the adsorption site, validating the reaction equation for specific adsorption, M+ClO4− → (M·ClO4) + e− (1) where M represents the positive electrode material and (M·ClO4) represents the solvated structure of Na+, while the solvent molecules are adsorbed on the Na sites without significant effect on the charge distribution36. A competitive adsorption mechanism between anions and solvent molecules within the IHP is proposed (Fig. 5d). In the initial state, the electrode surface sites are mostly occupied by solvent molecules due to the low salt concentration (1 mol L−1) in the electrolyte. During the charging process, the anions in the diffuse layer approach the IHP under the electrostatic attraction of the positive charges on the electrode surface37,38. Here, the anions gradually occupy the surface sites while the solvent molecules are displaced from the IHP via the solvated structure of Na+39.

Based on the DFT-optimized model, the comparison of the electron escape work functions of the electrode surface with different adsorbed species yields higher work functions in the presence of anion-specific adsorption (Fig. 5e), which indicates that anions generate a barrier that prevents electrons from transferring across the interface38,40. On the other hand, the ab initio molecular dynamics (AIMD) simulations with complete electrolyte models demonstrate that the average lengths of Na−O bonds around the anion-adsorption sites are shorter than those at molecule-adsorption sites (Fig. 5f and Supplementary Fig. 23, the initial and final atomic configurations of AIMD simulations are shown in Supplementary Data), consistent with DFT results in Supplementary Fig. 24 and Supplementary Table 5. Interfacial charge redistribution alters the coordination environment of superficial Na, and shorter Na−O bonds after anion-specific adsorption reflect stronger Na−O binding energy41, thereby hindering Na+ extraction. Systematically, Fig. 5g summarizes the charge transfer process at the molecule-adsorbed and the anion-adsorbed interfaces. At molecular-adsorbed sites, the deintercalated Na+ form solvated structures directly with solvent molecules. At anion-adsorption sites, however, the anion-rich IHP disrupts the Na+ solvation by shielding solvent access and prevents superficial Na+ from deintercalation. Meanwhile, the charge escape barrier increases the difficulty of charge transfer across the interface. Therefore, the charge transfer process is obstructed at anion-adsorbed sites, which explain the substantial Rct shown by the layered positive electrode under deep SoC (Fig. 4d, f).

It is necessary to elucidate the relationship between anion-specific adsorption and interfacial charge transfer reactions. The state variable θ (anion coverage on the electrode surface) is introduced to characterize the specific adsorption process, and combined with the simplified equivalent circuit with inductance (Supplementary Fig. 25), the Y0F (interfacial Faraday admittance) can be expressed by IF (Faraday current), θ, and E (electrode potential):
Y0F = IF/Rt + 1/(Ro + 1/jωL) = IF/Rt + (∂IF/∂θ)ss * (dθ/dE) (2)
in which a = (∂Ξ/∂θ)ss > 0, denoting the partial differential of the change rate (Ξ) of the state variable with respect to the state variable and being constant positive due to the stability requirement of the small signal system. Thus, the physical significance about the inductive signal in the EIS was obtained. When the trend of the effect of E on the θ and the effect of θ on the IF are the same, the inductive signal appears (Fig. 4e), and if the trend is opposite, the capacitive signal appears (Fig. 4d, f). Obviously, as the potential increases, the anion coverage increases, so the key is how θ affects the IF. Dividing the IF into I1 and I2 that represent the Faraday current densities of the solvent-adsorbed sites and anion-adsorbed sites respectively and assuming (∂IF/∂θ)ss = 0, we obtain the anion coverage threshold θ0 for the transition from inductive loop to capacitive arc at high SoC according to the Butler-Volmer equation42,
θ0 = I1 / (I1 + 2I2) (3)
The relationship between anion-specific adsorption and charge transfer reaction is clarified (Supplementary Fig. 26), where anion-specific adsorption plays a facilitating role at θ < θ0, whereas the inhibition of anions within the IHP will dominate at θ > θ0. Detailed mathematical derivations are given in Supplementary Note 2.

CEI formation and lattice reconstruction at surface
The surfaces of cycled NM and NMCFT electrodes were dissected by XPS (Fig. 6a), where the Na−F bond signal (~685.5 eV, contributed by the inorganic CEI component of the metal fluoride, such as NaF) is stronger for NMCFT, while the TM−O bond signal (~529 eV, contributed by the bare active particles) is more pronounced for NM, suggesting that the NMCFT surfaces generate more F-rich inorganic CEI that protects the active particles from direct contact with the electrolyte. The 3D rendering images and normalized depth profiles of Time-of-flight secondary-ion mass spectrometry (ToF-SIMS) (Fig. 6b, c and Supplementary Fig. 27) manifest that the NMCFT surface is distributed with thin inorganic NaF2− groups (~25 nm thickness) and relatively thicker organic groups (tens of nm thickness), whereas the surfaces of NM (represented by MnO− groups) are bare, and neither organic nor inorganic groups can completely encapsulate the active surface. The EIS results of cycled electrodes and corresponding distribution of relaxation times (DRT) spectra (Supplementary Fig. 28) also demonstrate the dense CEI on the NMCFT surface. Anion-specific adsorptions in the IHP on the NM and NMCFT surfaces induce the different CEIs after cycling17,21. At high SoC, the anion-rich IHP on the NM surface lacks solvent molecules decomposed into CEIs, whereas solvent molecules are still present on the NMCFT surface until 4.3 V SoC and eventually form the F-rich inorganic CEI.

After 50 cycles at C/3, severe cracks and TM dissolution appeared on the NM surface, as demonstrated in Fig. 6d, resulting in more active surfaces being exposed to the electrolyte. In comparison, the particles of NMCFT maintained an intact surface (Fig. 6e) that provided a stable carrier for a dense CEI. Further sXAS characterization of the NM electrode surface (left in Fig. 6f) reveal that the pre-edge peak of the O K-edge spectra shows an overall shift and decay after cycling. This shift indicates severe lattice distortion and a phase transition from the layered structure to the redox-inactive rock-salt structure (Fm3m space group, Fig. 6d)43. The weakening of the peak strength indicates the surface O-sublattice accompanied by TM rearrangement, leading to a worse unfavorable transition near cracks and TM dissolution regions44,45. In contrast, the NMCFT shows that only a slight decrease of the t2g peaks after cycling (right in Fig. 6f), indicating that the surface lattice reconstruction is alleviated considerably, as evidenced by the ~6 nm thin rock-salt-mixed layered phase (Fig. 6e). The retention of oxygen redox activity that proved by preserved shoulder peaks is one reason for the stable high-potential cycling stability. According to the XRD (Supplementary Fig. 29) and XAS (Supplementary Figs. 30 and 31) results, the long-range layered structure and the TM coordination environment of NM remained well within 50 cycles and is disrupted until 100 cycles by mechanical stress (resulting from P2−O2 phase transition), which indicates that the oxygen redox activity is the first to deteriorate and causes the capacity of NM to decay rapidly. Furthermore, due to the absence of CEI, stress corrosion cracking accelerates the failure of NM under long-term cycling conditions44. Rather, the F-rich inorganic CEI and reversible ARR pathway enable NMCFT to maintain both of the TM-provided and oxygen-provided capacity, demonstrating long-term cycling stability in a wide operating potential window.

In view of the high rate and cycling performance of NMCFT and in order to obtain evaluation under industrial conditions, we scaled up the synthesis of ~10 kg per batch and assembled lab-scale (6 mAh at C/3) and Ah-scale (5.6 Ah at C/10) pouch-type full cells against hard carbon (HC) negative electrodes as the experimental batch (Supplementary Fig. 32). When cycled between the full-cell voltage range of 1.5−4.3 V (versus HC), the lab-scale pouch cell with NMCFT demonstrated the cycling performance with 94.37% capacity retention after 100 cycles at the current density of 60 mA g−1 (Fig. 6h). Without any presodiation, the Ah-scale pouch cell delivered high initial-cycle Coulombic efficiency (81.4%) that is close to the theoretical level of HC, a high specific energy (156 Wh kg−1, calculated based on initial discharge capacity of 17.75 Wh and entire pouch weight of 113.81 g after secondary sealing) and sustained cycling stability (~80%, over 300 cycles at 120 mA g−1) (Fig. 6i). The assembly processes of the pouch-type full cells are shown in Supplementary Fig. 33 and detailed parameters are described in Supplementary Tables 6 and 7.

Methods
Synthesis of NM and NMCFT
The Na0.7Ni0.35Mn0.65O2 (NM) and Na0.7Ni0.27Mn0.53Cu0.04Fe0.08Ti0.08O2 (NMCFT) positive electrodes were synthesized by the solid-state reaction. Na2CO3 (Sinopharm Chemical Reagent Co., Ltd.), NiO, MnO2, CuO, Fe2O3, TiO2 were weighed in stoichiometric amounts with 5% excess of Na2CO3. All oxides were purchased from Aladdin. The powders were added with a small amount of ethanol and ball milled at 400 rpm for 8 h. The mixed powders were dried and pressed at 16 MPa to pellets with a diameter of 12 mm. The pellets were calcined in air at 950 °C for 15 h and after natural cooling, they were milled into powder and kept in a glove box under argon atmosphere (O2 < 1 ppm, H2O < 0.1 ppm; all subsequent argon atmospheres are maintained at these O2 and H2O concentrations) for electrode preparation.

Electrochemical characterizations
The 2032-type coin cells were applied to test the half-cell electrochemical performance of NM and NMCFT. The cathodic electrodes were prepared by 75 wt% positive electrode powders with 15 wt% Super P (TIMCAL, purchased from Shenzhen Youyan Technology Co., Ltd) as the conductive agent and 10 wt% polyvinylidene difluoride (PVDF, ARKEMA HSV900, purchased from Shenzhen Youyan Technology Co., Ltd) as the binder. All the mass was dispersed in N-Methylpyrrolidone (NMP, 99% purity, purchased from Aladdin) by ball milling to form a slurry and homogeneously coated onto aluminum foil (99.9% purity, purchased from Shenzhen Youyan Technology Co., Ltd, 12 μm thick, without any surface treatment) and dried in a vacuum oven at 80 °C for 12 h. Each dried positive electrode had a diameter of 10 mm, a thickness of 85−105 μm (including aluminum foil) and the loading mass of the active materials was 3−4 mg cm−2. The half-cells were assembled in the glove box under argon atmosphere with fresh sodium metal (99.8% purity, purchased from Thermo Fisher Scientific, with oxide layer scraped clean using a knife) rolled by hand to a thickness of ~0.5 mm and a diameter of 12 mm as the negative electrode, and Whatman GF/D glass fiber (16 mm diameter) as separator. The composition of the electrolyte was 1 M NaClO4 in propylene carbonate (PC) with an additional 5% fluoroethylene carbonate (FEC), purchased from Suzhou DodoChem Technology Co., Ltd and stored in aluminum bottles under an argon atmosphere. Electrolyte transfer was performed in a glove box using polypropylene pipette tips, with 150 μL added to each coin cell.

The lab-scale pouch cells used the same positive electrode, separator and electrolyte as coin cells, but with HC (Kuraray) as the negative electrode. The slurry of negative electrode was prepared with 94 wt% of HC, 3 wt% of carbon nanotube (CNT, TUBALL TM BATTH O, purchased from Shenzhen Youyan Technology Co., Ltd), 2 wt% of sodium carboxymethyl cellulose (CMC, DAICEL CMC 2200, purchased from Shenzhen Youyan Technology Co., Ltd), and 1 wt% of styrene butadiene rubber (SBR, Zeon BM-451B, purchased from Shenzhen Youyan Technology Co., Ltd) using water as solvent, and the slurry was coated on an aluminum foil and dried in a vacuum oven at 80 °C for 12 h. The cathodic loading of the lab-scale pouch cell was 3.72 mg cm−2 and the anodic loading was 1.316 mg cm−2. The N/P ratio was controlled at 1.10 ± 0.02 based on the specific capacities at 0.1 C (1 C = 180 mA g−1). The separator and electrolyte used in the lab-scale pouch cell are the same as those in the coin cell. Electrolyte transfer was performed in a glove box using polypropylene droppers, with ~2 g added to each lab-scale pouch cell. The pouch cells first underwent a formation process: the fresh pouch cells were first charged to 3.2 V at 0.05 C, then to 4.3 V at 0.1 C, and then degassed and second-sealed when the cells were discharged to 3.8 V at 0.1 C.

The cathodic electrodes of Ah-scale pouch cells were prepared by NMCFT (96.3 wt%), PVDF (1.7 wt%), Super P (1 wt%) and CNT (1 wt%). The anodic electrodes were prepared by HC (94 wt%), CMC (2 wt%), SBR (2 wt%), Super P (1 wt%) and CNT (1 wt%). All electrodes loaded with active materials were stacked after compression (compacted density of 3.1 g cm−3 for positive electrodes and 1.0 g cm−3 for negative electrodes). The separator was purchased from Dongguan Zhongyu New Material Co., Ltd. and consisted of a PE/boehmite ceramic composite material coated with PVDF on both sides. The thickness parameters are as follows: polyethylene 12 μm, ceramic layer 4 μm, PVDF layer 1 μm, with a total thickness of 18 μm. The total separator area per each Ah-scale pouch cell is 1827 cm2 (width 12.6 cm, length 145 cm), with the ceramic layer in contact with the positive electrode. The electrolyte used in Ah-scale pouch cell was purchased from Zhuhai Smooth Way Electronic Materials Co., Ltd., model number SWYM-A003, and was always sealed in argon-filled aluminum bottles. Electrolyte transfer was performed in a glove box using polypropylene droppers, with ~21 g added to each Ah-scale pouch cell. The N/P ratio was controlled at 1.10 ± 0.01. The Ah-scale pouch cells first underwent a formation process: the fresh pouch cells were first charged for 15 h at 0.05 C, then to 3.8 V at 0.05 C, then to 4.3 V at 0.1 C, and then degassed and second-sealed when the cells were discharged to 3.8 V at 0.1 C. All devices and detailed processes for kilogram-scale scale-up of NMCFT and assembly of Ah-scale pouch cells are presented in Supplementary Fig. 33. Throughout all processes, except for certain operations in the glove box, the environmental dew point must be maintained below –40 °C.

All coin cells and lab-scale pouch cells underwent electrochemical measurements on the Land CT2001A tester system, and Ah-scale pouch cells were tested on the Neware CT-4008Q-5V12A-S1 tester system. The specific capacities of coin cells and pouch cells are calculated by dividing the charging or discharging capacity by the mass of the positive electrode active material in the cells. All cells were tested under controlled temperatures of 25 ± 1 °C. All pouch cells were tested within explosion-proof boxes, ensuring only ambient temperature control without local temperature measurements. Galvanostatic cycling was applied in the electrochemistry measurements, unless specified otherwise. The rate measurements were performed by directly applying the specified currents after 3 cycles of activation at 0.1 C. GITT tests were composed of periodic constant current pulses and relaxation processes with a pulse time of 30 min and a relaxation time of 10 h. All data points for GITT tests were acquire data frequency of 0.5 Hz (once every 2 s). The Na+ diffusion coefficients were calculated based on the standard equation (Equation S30 in Supplementary Information), assuming a linear relationship between the transient voltage change and the square root of time (τ1/2) during the pulse. Note that in the high-potential region, the calculated values represent ineffective diffusion coefficients, as the electrochemical behavior is heavily influenced by thermodynamic hysteresis and interfacial adsorption as discussed in the main text. The Fast-GITT tests had a pulse time of 15 min and a relaxation time of 20 h. The long-period relaxation tests had a relaxation time of 120 h. All data points for GITT tests were acquire data frequency of once every 2 s.

EIS measurements were performed on the Autolab Electrochemical Workstation (Metrohm Autolab). GEIS and PEIS tests were conducted using a three-electrode Swagelok cell with Na metal as the reference electrode, while standard EIS tests were conducted using coin cells with Na metal as both the reference and counter electrodes. All EIS measurements utilized potentiostatic perturbations.

For standard EIS tests, the applied AC small signal amplitude was 10 mV and the testing frequency range was 50 kHz−0.01 Hz. 20 data points were sampled per decade. The EIS tests were conducted around the quasi-stationary potential (for charged or discharged cells) or open-circuit potential (for fresh cells).

For GEIS tests, the test parameters were identical to those of standard EIS (10 mV amplitude, 50 kHz−0.01 Hz, 20 points per decade). The three-electrode cells were subjected to a current density of 18 mA g−1 (0.1 C) to reach the target potentials and underwent a 1-hour relaxation before each EIS test to reach steady state. Each EIS test was conducted around the quasi-static potential after the relaxation.

For PEIS tests, the potentials of the applied DC signals were varied between 2.2−4.4 V in the steps of 2 mV, and after stabilizing for 1 min at each step an AC small signal was applied for the mid-frequency EIS tests, with an amplitude of 5 mV and a frequency range of 10 Hz −0.1 Hz. 20 data points were sampled per decade. Each mid-frequency EIS test was conducted around the quasi-static potential. The small potential steps and constant-potential processes allow the electrode to be approximated as steady state during each EIS test in the PEIS tests.

Material characterization
Powder XRD and ex-situ XRD were performed on a Bruker D2 X-ray diffractometer fitted with Cu Kα X-ray (wavelength λ=1.5406 Å) radiation, with the test step set at 0.02°, 0.25 s per scanning step, and a scanning range of 10–80°. For in-situ XRD characterization, we used a self-made cell mold with a window and tested on a Bruker D8 X-ray diffractometer while charging and discharging. One scan ranged from 10−80° and took 30 min. All XRD results were fitted by Rietveld refinement using GSAS II software52, and CIF card of the standard P63/mmc space group was used for fitting.

FAULTS simulation results were performed on the FAULTS software53. By varying the different P/O ratios in the structure, the FAULTS software simulated realistic X-ray diffraction spectra. The pure P2-type phase consisted of 100% P-type stackings and the pure O2-type phase consisted of 100% O-type stackings. The structural evolution of P2−Z−OP4 was modeled by the gradual addition of O-type slip units in the P2-type phase structure. The final OP4-type phase structure had a P/O ratio of 1:1. The corresponding models were developed according to the above structures and the spectra were output through FAULTS software. The X-ray diffraction spectrum of the P2/O2 composite structure was plotted by superposition of the spectra of pure P2 and pure O2-type phase.

ICP-MS results were performed on NexION 350D ICP-MS. Nitric acid was used to dissolve the powder samples and microwave digestion was performed on NMCFT powder to completely dissolve the Ti element.

SEM was conducted on an S4800 field emission scanning electron microscope. SEM was conducted on the COXEM EM-30AX+ scanning electron microscope. Powder samples were stuck on conductive adhesive for SEM test.

HRTEM and EDS mapping were conducted on Talos F200X transmission electron microscope. STEM-HAADF and ABF were conducted on the JEM-ARM300F2 scanning transmission electron microscope with double spherical aberration correctors, and image optimization was performed using the Filter plugin with the optimization step of 5. Powder samples were milled in ethanol and then dispersed by sonication. The supernatant was taken and dropped on the ultra-thin carbon film and dried completely using IR lamp for TEM and STEM tests. The cycled electrode samples were first immersed in NMP to dissolve the active materials and then added to ethanol for ultrasonic dispersion. A drop of supernatant was taken on a dual networking and left to dry completely for STEM tests.

XAS data was collected on the BL20U1 beamline at Shanghai Synchrotron Radiation Facility (SSRF) and sXAS data was collected on beamline MCD-B (Soochow Beamline for Energy Materials) at National Synchrotron Radiation Laboratory (NSRL).

XPS was conducted on the Thermo ESCALAB Xi+ X-ray photoelectron spectrometer and peak splitting was performed using Advantage software. The C−C peak of all data was set to 284.8 eV for calibration.

ToF-SIMS was conducted on IONTOF M6 hybrid SIMS and the data were analyzed on Surfacelab software. The ToF-SIMS tests were performed in negative-ion mode using a thermally ionized Cs ion source. The test area was a square with a side length of 100 μm and the etching duration was 200 s. The etching rate was approximately 0.5 nm s−1.

Preparation of all samples for all ex-situ tests and cycled tests: the cells were disassembled to get the electrodes within 1 h after charging and discharging. The electrodes were soaked in DMC (99.9% purity, purchased from Suzhou DodoChem Technology Co., Ltd) for 5 min to wash glass fibers and residual electrolyte away, and then stored in a glove box under argon atmosphere. All samples for ex-situ tests were transferred from the glove box to the testing devices sealed within argon-filled centrifuge tubes, with final loading completed inside the glove box.

Theoretical calculation method
Periodic spin-polarized DFT calculations were performed using the Vienna ab-initio simulation package (VASP) in a plane wave basis set54,55. The Perdew-Burke-Ernzerhof (PBE) functional within the generalized gradient approximation (GGA) was applied to evaluate the exchange and correlation energies56. The cutoff energy of 400 eV and Monkhorst-Pack K point grid spacing of 3×3×1 were chosen for all structures57. The convergence criteria of energy and force calculations were set to 10−5 eV and 0.01 eV Å−1, respectively.

The AIMD simulations were performed using CP2K/Quickstep package with explicit solvent for different adsorption configurations55. The simulations involve adding 1 M NaClO4 (or NaPF6) in PC with 5% FEC on P2-type layered NM positive electrodes, The structures were equilibrated at 300 K in an NVT ensemble for 20~25 ps, trajectories were sampled over the last 5 ps.

The adsorption energy (Eads) between adsorbed species (FEC, PC, ClO4− and PF6−) and the electrode surface (*) was defined respectively as follows58
Eads(FEC) = E*FEC − E FEC − E* (4)
Eads(PC) = E*pc − E pc − E* (5)
Eads(ClO4−) = E*ClO4 + 1/2E H2 − E* − E HClO4 − eU vs:RHE (6)
Eads(PF6−) = E*PF6 + 1/2E H2 − E* − E HPF6 − eU vs:RHE (7)

Peer review rounds
[{"round": 1, "reviewer_comments": {"reviewer_1": "The manuscript combines both bulk phase and interfacial environment to explain the electrochemical property of P2-type NMCFT cathode in Na-ion batteries. Impressively, the authors provide a competitive adsorption model of anions and molecules in the IHP to elucidate the charge transfer process on the electrode interface. The anion-specific adsorption promotes charge transfer at the interface as the potential changes and consequently facilitates the fast-charging performance of the P2-type cathodes. Moreover, the authors correlate the anion-specific adsorption to the generated CEI after cycling and further assemble pouch cells to embody the application potential of P2-type cathode materials. This manuscript provides an insightful perspective for Na+ diffusion and charge transfer in intercalated materials for energy storage. Therefore, I recommend its publication in Nature Communications. The following are some suggested comments to improve the manuscript and some doubts that need to be further explained.\n1. The authors quantitatively compare the fast-charging performance of NM and NMCFT using 4.0 V as the cut-off in Fig. 1, and conclude that the portion above 4.0 V dominates the degradation of cycling and rate performance. The rate performance tests between 2.2 V and 4.0 V are recommended for full validation.\n2. Although P2-type Ni/Mn-based cathodes have been widely studied and performed ex-situ XRD tests in Supplementary Figure 9, it is still recommended to supplement a comparison of dO-Na-O evolution between NM and NMCFT cathodes.\n3. In Fig. 4, the authors investigate in depth the phenomenon “inductive loops” phenomenon and attribute it to the anion-specific adsorption. Note that only NaClO4-based electrolytes were used in this work, and whether the phenomenon differs for electrolytes with other solutes. And, whether the “inductive loops” phenomenon occurs only in P2-type cathodes. Please give an expanded explanation.\n4. The authors mentioned the finite boundary diffusion impedance (ZT), unlike the common Warburg impedance (Zw) in intercalated electrodes. The reason for the occurrence of boundary layer diffusion in an intercalated electrode such as NMCFT needs to be given.\n5. There is a difference between the authors' statement in Supplementary Figure 25 that the DRT peak between 10–2–10–1 s indicates the formation of CEI and the common viewpoint that the first semicircle of the Nyquist plot indicates CEI or SEI. Please give the scientific basis.\n6. There are still minor spelling and formatting mistakes such as in the third paragraph in the introduction “Most researchers have neglected the face…”, please check and correct them.", "reviewer_2": "Xu et al. reported an interesting work to explore the influence of both interfacial properties and bulk factors on the fast-charging performance of layered cathodes in sodium-ion batteries. Unlike common studies focusing on the cathode-electrolyte interphase (CEI), the authors emphasize the overlooked role of specific adsorption within the inner Helmholtz plane-IHP on the charge transfer kinetics by comprehensive analysis of advanced characterizations and theoretical calculations. This work merits publication in Nature Communications and some questions and feedbacks are listed below:\n1) The diffusion behaviors of the cathodes affect the anion-specific adsorption within the IHP, while Xu's group recently also suggested that the layer spacing of the carbon-based anodes affects the anion distribution in the IHP (Joule 2025, DOI: 10.1016/j.joule.2025.101874). For layered structures, interlayer spacing and diffusion properties are directly linked, does the layer spacing of the cathode similarly influence the anion distribution in IHP?\n2) In Page 4, it is necessary to explain the statement “…the leftward wavevectors (K) of Cu-TM (~0.891 Å1) and Fe-TM (~0.609 Å1) shell peaks at 4.4 V SoC indicate that the mass numbers of the elements coordinated to Cu and Fe become smaller, further verifying the interlayer reversible Cu and Fe migration…”\n3) The “inductive loops” in the EIS plots (Fig. 4a and b) require a complete quantitative analysis. It is suggested that the authors add images or tables of inductance and capacitance values in the mid-frequency region, and analyze the physical significance of the evolutions.\n4) The differences and correlations between the GEIS test and the PEI test are need to expand in terms of electrochemical principles and experiments. Whether the PEIS test makes the battery fulfill the steady-state conditions for performing the EIS test like the GEIS test?\n5) There are more pronounced MnO- signals on the NMCFT surface than NM in the 2D intensity distribution of Tof-SIMS (Supplementary Figure 24), and perhaps more TM elements of NMCFT are exposed. Why the phenomenon is not consistent with Fig. 6b?", "reviewer_3": "This manuscript by Xu et al., reports the investigation of phase evolution of two P2-type cathode materials with Na0.7Ni0.35Mn0.65 as the reference compound. Improved performance is demonstrated when introducing more transition metals, including Cu, Fe, and Ti. The more interesting part is the investigation of charge transfer process and the cathode electrolyte interphase (CEI) on the cathode surface, however, which is separated from the first structural analysis. I could not recommend its publication considering other issues outlined in detail below.\nThe Ni/Mn based P2 cathode has been intensively investigated, and the phase evolution and P2 to O2 transition involving the intergrowth has been somehow reported in higher quality and correctly in previous publications. The current work couldn´t reveal the synergy between the bulk phase and the interfacial optimization.\nThe compositional design is based on Na0.7Ni0.35Mn0.65 to prepare the Na0.7Ni0.27Mn0.53Cu0.04Fe0.08Ti0.08O2, which I understand in a way that the Fe is to replace the Ni, and the Cu and Ti is to replace the Mn. Therefore, the Fe has the oxidation state of III, Cu II and Ti IV. However, there is no compositional analysis on the oxidation number of the transition metals. The structural refinement has constrained the occupancies to the designed nominal composition, while lack of proof (such as ICP) for the molar content of the constitutional elements. Regarding the Rietveld refinement, what´s the reason to fix the Uiso of transition metals to the same value? And what´s the reason to have more Naf in the structure when Cu, Fe and Ti are incorporated into the structure? The Na content is quite determining for the initial charge capacity, one suspects that the Na content in the NM composition is less, as the initial Coulombic efficiency is lower. How is the Na content in these two compositions determined?\nThe identification of Z phase is not correct, the phase involving intergrowth of P-stacking and O-stacking phase, and the OP4 is the intermediate ordered phase. Apparently, the simulated xrd patterns (Fig. 3bIII) don´t match with the in-situ xrd patterns indicated as Z phase region (Fig. 2 a). The broaden peak shifted to around 2theta of 17.5 degree, while the Z phase identified in the simulated xrd patterns ended at around 2theta of 16.5 degree. References can be found in Energy Environ. Sci., 2019,12, 2223-2232 and Nature Energy 9, 184–196 (2024).\nThe core idea in this manuscript is the accumulation of positive charge on the cathode surface, however this remains loosely defined and no proof. The reviewer couldn’t find Ref. 23 and Ref.24 could support this explanation. What is the thermodynamic polarization? Thermodynamic is not about kinetics and polarization is. The Ref. 27 looks randomly cited in such case. Kinetics hysteresis and the polarization is the same thing. Terminology seems to be a big issue in this manuscript; it just raises confusion.\nIt´s interesting to see that ` Interfacial charge redistribution alters the coordination environment of superficial Na, ….`, while not proof is provided. This claim can be verified through X ray absorption spectroscopy, the TEY and TFY will help to differentiate the bulk and surface coordination environment.\nNo discussion on Fig.6f could be found.\nThe active material loading is too low in the coin cell testing.\nWhat´s the thickness of the cathode used in the pouch cells?"}, "author_response": "Thank you for valuable comments to help us improving the manuscript. Your suggestions and comments have been carefully considered and addressed, and all the responses are listed in the following section, point by point. In addition, all changes in the revised Manuscript and Supplementary Information have been highlighted in yellow for your convenience.\n[... Author response text for Round 1 ...]"}, {"round": 2, "reviewer_comments": {"reviewer_1": "The revised manuscript is ready for acceptance of publication.", "reviewer_2": "The manuscript has been modified accordingly, and this work can be accepted for publication in the journal.", "reviewer_3": "I recommend its publication as the questions have been well addressed by the authors."}, "author_response": "We would like to thank the editor and the reviewers for their time and effort in reviewing our manuscript. We are pleased to learn that our previous revisions have satisfactorily addressed all concerns. Below is our response to the reviewers’ final remarks.\nFor Reviewer #1: The revised manuscript is ready for acceptance of publication. Response: We thank the reviewer for their positive assessment and recommendation for acceptance.\nFor Reviewer #2: The manuscript has been modified accordingly, and this work can be accepted for publication in the journal. Response: We appreciate the reviewer’s confirmation that the manuscript has been modified appropriately and is now suitable for publication.\nFor Reviewer #3: I recommend its publication as the questions have been well addressed by the authors. Response: We are grateful to the reviewer for acknowledging that our revisions have addressed the questions raised. We thank them for recommending publication."}]

Decision letter