An ammonia-induced universal synthesis approach for manganese based layered oxides

Domain: physics
Published: 2025-12-02
Editorial outcome: major_revision

Abstract
Manganese-rich layered transition-metal oxides are of great interests as positive electrodes for sodium-ion batteries considering their high specific capacity and operating voltage. However, it remains challenging to synthesize layered transition-metal oxides with high stability and rapid ion-migration kinetics. Herein, we develop a straightforward and universal synthesis approach for P2-type layered transition-metal oxides by simply updating the conventional solid-state reaction with only additional dicyandiamide introduced. We reveal the ammonia released successfully restrains the irreversible oxygen redox for stabilizing cationic migration and decreases the crystallization temperature for high-quality structure formation. The effectively suppressed Jahn-Teller distortion and facilitated Na+ transport kinetics therefore endow Na1/2MnO2 with overall performance enhancement, including a 6.7-fold improvement in rate capability and enhanced cycling stability. Similar performance enhancements are also found in other manganese-based layered transition-metal oxides, presenting great strategy universality. This work marks a key step forward in the synthesis-by-design of high-performance sodium-ion batteries electrode materials.

Introduction
Recently, sodium-ion batteries (SIBs) are emerging as one of the most promising alternatives to its lithium counterparts for next-generation energy storage devices due to their advantages of sodium availability and reliable performance advantages for fast charging1,2. The design of positive electrodes with excellent redox reversibility and high Na+ mobility is the key focus for promoting the commercialization of SIBs3,4. P2-type Mn-based layered transition-metal oxides (LTMOs) attract broad attentions because of their high theoretical capacity, compositional flexibility and structure adaptability5. Their structural similarity to Li-LTMOs also enable the great potential for feasible transformation of current manufacturing lines for lithium ion batteries6,7. However, most materials still suffer from unstable cyclic performances and low reversible capacity8,9, while the development of high-performance positive electrodes for SIBs still remains challenging.

It is always accepted for Mn-rich LTMOs that undistorted [MnO6] octahedral unit is formed by the presence of Mn4+ for stabilizing structure, while high-spin Mn3+ always results in a Jahn-Teller distortion that causes the irreversible multiphase transition and thus rapid capacity decay10,11. Meanwhile, the synthesis of high-quality LTMOs with stable ion tunnels can not only stabilize the structure but also accelerate Na+ desodiation with high reaction kinetics12,13. To date, the solid state reaction is viewed as the primary synthesis strategy for LTMOs in practical use14,15. However, its limited ion diffusion rate and the difficulty for crystal structure control in solid phase retards its use for high-performance positive electrode synthesis. Currently, the development of Mn based LTMOs with high stability and rapid ion migration kinetics faces the critical challenges and it is highly imperative to propose effective synthesis strategies that offer new possibilities for SIB positive electrodes.

In this work, we develop a universal synthesis strategy for P2-type LTMOs positive electrodes for SIBs. Na1/2MnO2 (NM) that is one of the most commonly studied LTMO benchmarks was selected as the modeled material. By introducing dicyandiamide (DA) into the conventional solid-state processing without additional treatment, we synthesize the optimized Na1/2MnO2 (NMD) with all-around performance elevation compared with its pristine counterpart NM. Ammonia released from thermal decomposition of DA was found to facilitate the solid-state reaction kinetics by creating oxygen vacancies to activate intermediate bondings and decrease thermodynamic barriers of ion migration. Meanwhile, the presence of NH3 gas phase can also decrease the interaction force between atoms/ions in solid reactants and lower the diffusion resistance, which reduces the crystallization temperature and promotes the formation of high-quality layered structure. In addition to the different working mechanism of gas-phase assisted solid state reaction revealed in this strategy, its broad universality for the synthesis of highly active P2-type LTMOs positive electrodes with various composition attracts our interests. Furthermore, the simple and direct processing with only DA additive introduced into the current solid-state synthesis route avoids complicated facility and hardware updates. The aforementioned merits of different mechanism, universal approach and simple synthesis routes enable this strategy with great potential for large-scale production of P2-type LTMOs. A detailed design principle for this strategy is illustrated in Fig. 1a.

Results
Synthesis and structural characterizations of NM and NMD
NM and NMD were synthesized by the classic solid-state reaction with the only difference in the addition of dicyandiamide as the ammonia release source. The X-ray diffraction (XRD) patterns along with highly fitted Rietveld refinement results (Fig. 1b–c) confirm the successful synthesis of P2-type Na1/2MnO2 with a space group of P63/mmc for both samples16. Inductively coupled plasma optical emission spectrometry (ICP-OES) analysis confirms their similar stoichiometry which is consistent with the designed reactant ratios (Supplementary Table 1). The further analysis of refined crystallographic data (Supplementary Table 2) shows the reduced a/b lattice parameter of NMD from 2.990 Å to 2.920 Å while c-axis increases from 11.127 Å to 11.163 Å, attributed to the weakening of O-O repulsion due to the formation of oxygen vacancies. This crystallographic change results in its enhanced structural stability compared with NM. Meanwhile, the decreased d(O-TM-O) and increased d(O-Na-O) values of NMD compared with NM are beneficial for inhibiting Mn dissolution and MnO6 layer gliding (Fig. 1d). Meanwhile, the primary peak (002) of NMD shows an apparent shift to the lower angle (Supplementary Fig. 1a), indicating its enlarged d(O-Na-O) spacing which can reduce Na+ diffusion resistance during insertion/extraction. Furthermore, NMD show significantly enhanced crystallinity with sharp and high-intensity (002) peak compared with NM (Supplementary Fig. 1a), confirming this proposed strategy can effectively overcome the bottleneck of conventional solid-state reaction for forming the P2-type LTMOs with high crystallinity. Furthermore, the presence of impurity phases is observed in NM, while a pure phase is confirmed for NMD (Supplementary Fig. 1b). The low-magnification scanning electron microscopy (SEM) image shows a higher proportion of hexagonal stacked layer structure for NMD compared with the irregular micron-particles of NM (Supplementary Figs. 2 and 3). Further high-resolution transmission electron microscope (HRTEM) analysis (Fig. 1e, g) reveals their crystalline nature with clear lattice strips and meanwhile enlarged spacing of (002) plane for NMD (0.574 nm) compared with that of NM (0.564 nm)17. The corresponding line profiles are presented in Supplementary Fig. 4. Furthermore, the fast Fourier transformation (FFT) patterns (insets of Fig. 1e, g) indicate the formation of single-crystalline structure that can be indexed to P2 phase within NMD, while NM patterns with diffraction rings demonstrate its localized polycrystalline structure. Geometric phase analysis (GPA) reveals the large stress fluctuations of NM (Fig. 1f) while the relatively homogenous stress distribution of NMD (Fig. 1h) indicates the mild phase transition in a wide temperature range effectively relieves the internal stress that is beneficial for enhancing the stability of crystal structure18,19. Scanning transmission electron microscopy (STEM) equipped with high-angle annular dark field (HAADF) (Fig. 1i) and annular bright field (ABF) (Fig. 1j) is employed to analyze the atomic-level crystal structure of NMD. In the HAADF mode, the light dots signify the transition metal atoms while the O and Na atoms are revealed within the dark columns, whereas the manifestation is reversed in the ABF mode17,20. According to the STEM-HAADF image, the TM-TM distance is 0.273 nm corresponding to the (100) plane of the P2 structure. Selected area electron diffraction (SAED) pattern along the [001] axis substantiates the single-crystalline hexagonal structure of the bulk phase of NMD (Fig. 1k). Correspondingly, energy-dispersive spectroscopy (EDS) analysis demonstrates the homogeneous distribution of Na, Mn, and O within NMD (Fig. 1l).

Structural growth analysis
To investigate the detailed influence of DA addition on the growth mechanism of LTMOs, the structural evolution from precursor to the final crystalline form was analyzed with in-situ XRD technique from room temperature to 500 °C (Fig. 2a–c). The (002) peak of sodium manganite at 15.6° can be detected at a low temperature of 320 °C for NMD (vs. 400 °C for NM), demonstrating this strategy significantly decreases the formation temperature of Na1/2MnO2 structure and allows the fully growth of crystals with more ordered atomic arrangement during wide temperature range. The (001) peak of Na2CO3 at 14.9° progressively disappears while the (222) peak of Mn2O3 at 33° also gradually diminishes with increasing temperatures, indicating the transformation into Na1/2MnO2. Meanwhile, a gradual sharpening of the (002) peak with the increase of the half-peak full width is also observed (Supplementary Fig. 5). The extended phase transition temperature region with the earlier appearance of crystalline structure endows NMD with alleviated internal structure stress, resulting in a more robust structure for stabilized Na+ storage. In addition, since non-uniform strains within the material will lead to variations in lattice spacing, the lattice strains of both samples were also calculated based on the neutron diffraction data (discussed later) using the Williamson-Hall equation (Supplementary Fig. 6)21. The calculated lattice strain is 7.5×10−2 % for NM and 5×10−2 % for NMD, indicating that NMD shows a robust crystal structure with alleviated internal strain.

The decomposition of Na1/2MnO2 was then analyzed through thermogravimetric (TGA) curves to investigate the detailed role of DA (Supplementary Fig. 7). DA undergoes a dramatic weight loss at 210–385 °C for NH3 release and this process is consistent with the weight losses within NMD considering the high content of DA in the precursor (~40 wt%), confirming the primary contribution of DA for ammonia release to the phase formation. The decomposition of DA also takes the primary weight change for the phase transition between 170-470 °C, in which ammonia facilitates the diffusion kinetics of Na+ and Mn3+ in the solid-phase reaction for promoting the early formation of intermediate phases. Meanwhile, the released ammonia can seize lattice O atoms by forming N-containing gases, leading to the formation of oxygen vacancies and accelerated solid-reaction kinetics. This is confirmed by the higher content of oxygen vacancies within NMD as characterized through its sharper electron paramagnetic resonance (EPR) signal than that of NM at g=2.003 (Supplementary Fig. 8)17. X-ray photoelectron spectroscopy (XPS) was then performed and the results (Supplementary Fig. 9, Supplementary Table 3 and Fig. 2c) show the content of Mn4+ within NMD decreases from 9.76 to 5.16 at.% whereas Mn3+ increases from 3.89 to 10.86 at.%, which is attributed to the introduction of oxygen vacancies in NMD22. High resolution O 1s profiles (Fig. 2d) can be deconvoluted into peaks associated with lattice oxygen (Olat) at 529.5 eV, oxygen vacancy (Ovac) peak at 530.9 eV and chemisorbed oxygen (Oads) at 531.5 eV23. The reduced Olat content of NMD from 27.44 to 19.84 at.% (Fig. 2d) confirms its high air stability and great inertness to H2O and CO2, which is highly desirable for stabilized positive electrodes24. Compared to NM, Oads content in NMD decreases from 13.25 to 12.97 at.% due to the partial oxygen seized by the produced ammonia during low-temperature sintering to create oxygen vacancies. The content of Ovac in NMD also increases and this is consistent with EPR results, induced by O capture from ammonia. The introduction of oxygen vacancies can lead to an unbalanced charge distribution through non-saturated coordination sites, creating a local electric field that enhances conductivity and provides extra Na+ storage sites. Furthermore, it can also effectively suppress oxygen release and retard the irreversible oxygen chemistry, which is generally accepted as one of the primary reasons leading to structural instability and voltage fading of LTMO positive electrodes25.

Ex situ X-ray absorption fine structure (XAFS) spectra were collected to reveal the change of valance states for metal elements between samples. The X-ray absorption near edge structure (XANES) at the Mn K-edge and the corresponding Fourier transformed of extended X-ray absorption fine structure (FT-EXAFS) spectra were presented in Fig. 2e–f. Compared to NM, NMD shows a left-shift to the lower energy for its Mn K-edge (Fig. 2e), indicating the oxidation of Mn with the decreased valence state, which coincides with the enlarged interatomic distances within the first Mn-O coordination shell that is presented with the increased Mn-O peak in Fig. 2f21. This observation is also consistent with the XPS result of NMD that presents the increased Mn3+/Mn4+ ratio compared with NM. Meanwhile, the decreased second Mn-Mn coordination shell peak of NMD (Fig. 2f) suggests the reduced interatomic distances of Mn, which demonstrates the contraction of MnO6 unit. The wavelet transform (WT) contour plot of EXAFS spectra of NMD (Fig. 2g) exhibit two distinct scattering peaks of Mn-O and Mn-Mn coordination shells located at (5.6 Å⁻1, 1.4 Å) and (5.6 Å⁻1, 2.3 Å), respectively. Correspondingly, the center of these scattering peaks is shifted with the same trend of Mn-O and Mn-Mn peaks of FT-EXAFS for those in NM (5.4 Å⁻1, 1.4 Å and 5.7 Å⁻1, 2.3 Å) (Fig. 2h)22. The Mn-Mn coordination shell in NMD with low R-value demonstrates the locally disordered Mn-Mn coordination environment that is caused by the introduction of oxygen vacancies. Meanwhile, the well-maintained shape without change from NM to NMD suggests the uniform introduction of oxygen vacancies into NMD bulk. The aforementioned analysis confirms the introduction of oxygen vacancies into NMD along with electronic/structural changes that are the basis for its electrochemical performance enhancement.

Electrochemical performance evaluation
The electrochemical performances of samples were continuously evaluated in half-cell configurations with galvanostatic charge/discharge (GCD) curves recorded (Supplementary Figs. 10 and 11). NMD presented significantly enhanced rate capabilities with 187.2, 178.6, 163.3, 142.9, 123.8 and 102.1 mAh g⁻1 delivered at 0.05, 0.1, 0.2, 0.5, 1.0 and 2.0 A g⁻1 compared with those of NM (144.3, 91.5, 61.9, 29.0 and 18.4 mAh g⁻1 at 0.05, 0.1, 0.2, 0.5 and 1.0 A g⁻1, respectively) (Fig. 3a). The capacity of NMD is even 6.7 times of NM at 1.0 A g⁻1, further highlighting its significantly enhanced rate capability. The addition content of DA was also optimized while NMD samples with different contents of DA always show the enhanced performance compared with NM (Supplementary Fig. 12). Its fully overlapped cyclic voltammetry (CV) curves within the initial 3 cycles (Supplementary Fig. 13a) also confirm the highly reversible redox reactions within NMD, which is not achieved in NM (Supplementary Fig. 13b). The current area for NMD is significantly larger than that of NM, indicating a higher Na+ storage capacity in NMD (Supplementary Fig. 14). Based on the CV, NMD also presents higher diffusion rate than NM (Supplementary Figs. 15–17) and narrow/sharp peaks with high intensities in dQ/dV curves (Fig. 3b), indicating its rapid redox reaction kinetics and great reversibility26. Electrochemical impedance spectroscopy (EIS) analysis (Fig. 3c, Supplementary Table 4) demonstrates that NMD shows the apparently increased Na+ diffusion kinetics with the effectively suppressed internal resistance and robust interfacial stability to deliver a fast ionic response, resulting in rate performance improvement27. Correspondingly, a 2.6-fold increase of derived Na+ diffusion coefficient (DNa+) value for NMD (2.41×10⁻13 cm2 s⁻1) compared with NM (8.46×10⁻14 cm2 s⁻1) (Fig. 3c and Supplementary Fig. 18) also confirms the rapid Na+ diffusion rate within NMD that is consistent well with CV results, facilitating its high discharge capacity under elevated specific currents. Galvanostatic intermittent titration technique (GITT) was also conducted to further probe into the ion diffusion kinetics of positive electrodes (Supplementary Fig. 19). Consistent with EIS results, NMD also demonstrated higher DNa+ values than that of NM, confirming its efficient Na+ transport kinetics. In addition, the in-situ EIS and the distribution of relaxation times (DRT) analysis were performed to reveal the evolution of the internal states of NM and NMD at different potentials. In-situ EIS analysis reveals that the charge transfer resistance initially decreases and then increases (Supplementary Fig. 20), fluctuating with the Na concentration in the electrode during the early charge-discharge process28. Notably, NMD exhibits a lower transfer resistance after charge-discharge cycles (Supplementary Fig. 21), highlighting its great ion transport kinetics. Further DRT analysis (Supplementary Fig. 22) shows lower intensity values across all regions for NMD compared to NM, indicating its reduced electrochemical impedance29,30. The reduced impedance and enhanced ion transport kinetics in NMD can be ascribed to oxygen vacancies that weaken the Na-O electrostatic interaction, thereby expanding interlayer spacing and promoting both Na-ion diffusion and electronic conductivity.

The long-term stability of samples was then evaluated at different specific currents (Supplementary Fig. 23, Fig. 3d–e). NMD presents excellent cycling performances with a capacity retention rate of 67.3% and 74.0% after 300 cycles at 0.2 and 0.5 A g⁻1, far surpassing those of NM (54.5% and 63.7%). Besides, NMD also exhibits 74.1 and 60.3 mAh g⁻1 after 500 and 1000 cycles at 1.0 A g⁻1, with the capacity retention of 80.0% and 64.1%. Correspondingly, the capacity attenuation of NMD is 0.036% per cycle, much lower than that of NM (0.084%) at 1.0 A g⁻1 (Fig. 3f). A direct comparison of overall electrochemical performances between NMD and other state-of-the-art LTMOs is also listed in Fig. 3g and Supplementary Table 5, showing appealing combination of positive electrode features4,8,11,23,31–36. Full cells were then evaluated with commercial hard carbon (HC) as the negative electrode (Fig. 3h). HC||NMD full cell exhibits an initial discharge capacity of 130.7 mAh g⁻1 at 0.05 A g⁻1 and excellent rate capability of 75.3 mAh g⁻1 at 1.0 A g⁻1 with a high specific energy (based on the mass of the active material) of 244.4 Wh kg⁻1 achieved (Fig. 3i and Supplementary Fig. 24). Additionally, it also demonstrates excellent cycling performance, maintaining a discharge capacity of 77.8 mAh g⁻1 with a retention rate of 66.7% after 300 cycles at 0.2 A g⁻1 (Fig. 3j and Supplementary Fig. 25). In contrast, HC||NM showed an overall-weakening performance with a retention rate of only 47.4% (23.7 mAh g⁻1). The HC||NMD full batteries present better performance advantages in capacities and specific energy (based on the mass of the active material) over the recently reported representative full cells (Supplementary Fig. 26 and Supplementary Table 6). Moreover, pouch cells (single layer) were assembled by using presodiated HC and NMD materials to assess their practical cycling stability within 1.5-3.6 V (Supplementary Fig. 27). HC||NMD pouch cell presents a reverse capacity of 140.0 mAh g⁻1 at 0.1 A g⁻1, which is higher than HC||NM pouch cell (86.9 mAh g⁻1). The corresponding GCD curves in Supplementary Fig. 27b further shows the excellent electrochemical reversibility of the NMD electrode. The HC||NMD pouch cell also presents high capacity retention of 82.5% for 200 cycles and high specific energy (based on the mass of the active material) of 210.3 Wh kg⁻1 at 0.5 A g⁻1, verifying its great practical potential as the SIB positive electrodes.

Structural evolution analysis
In-situ XRD analysis was then performed to reveal the phase transition changes during the initial charge/discharge cycles. The extraction of Na+ weakens the attraction between sodium and oxygen layers, expanding interlayer spacing. Both samples therefore show (002) and (004) peaks shift towards lower angles during charging and the opposite trend during discharging (Fig. 4a, d). During the high-voltage charging stage, the disappearance of the P2 phase (002) peak is accompanied by the emergence of a new diffraction peak at 17.4°, which is attributed to the (001) crystal plane of the OP4 phase. The original lattice symmetry and periodicity are disrupted due to the presence of oxygen vacancies, which are more favorable for the transition to the OP4 phase during Na+ extraction. As visualized in contour plot (Fig. 4b), NMD maintained a small peak shift of 0.52° for (002) plane and 1.09° for (004). The calculated c-axis value is also increased by only 2.56% of the pristine cell parameter due to the expansion of Na layers (Fig. 4c), confirming its small volume change during discharging. Furthermore, after subsequent charging and discharging, the diffraction peak returns to its original position with no new phases detected, indicating high reversibility of the P2-OP4 phase transition within NMD34. In contrast, NM exhibits more pronounced structure changes with a larger shift in both (002) (0.71°) and (004) peaks (1.22°) as well as c-axis value (2.87%) in Fig. 4e, f. The incompletely recovered diffraction peaks of NM suggest its poor structural reversibility. A detailed structural change comparison of samples during discharging/charging is illustrated in Fig. 4g. NMD achieves minimum structural changes and robust structural reversibility over NM during Na+ de-intercalation due to oxygen vacancies to buffer interlayer stresses in NMD.

Analysis of Na+ transport mechanism
To further elucidate the impact of ammonia-induced strategy on Na+ transfer kinetics within Na1/2MnO2, neutron powder diffraction (NPD) combined with maximum entropy method (MEM) was performed to visualize the transport path of Na+. Based on the NPD Rietveld refined data (Fig. 5a and Supplementary Table 7), a/c axis of NMD is shrunk that is consistent with XRD results. Furthermore, the corresponding three-dimensional nuclear density distribution calculated by MEM (Fig. 5b) shows Na+ with positive scattering lengths migrate along the Z path through oxygen sites or vacancies in NMD, which provides shorter Na+ diffusion distances and more active sites than NM37. This indicates enhanced Na+ transport kinetics along a-axis with an optimized diffusion pathway and low energy barrier, contributing to improved capacity and rate performance in NMD.

Density-functional theory (DFT) calculations were performed to further investigate the diffusion mechanism of Na+ using a computational model that introduces oxygen vacancies for NMD (Supplementary Fig. 28). Based on the analysis for electronic density of states (DOS) at different desodiation states (Fig. 5c–e and Supplementary Fig. 29), NMD demonstrates the dramatically increased number of electronic states at the Fermi level and significantly reduced bandgap compared to NM, even in different desodiation states38,39, which confirms its improved electronic conductivity. The partial density of states (PDOS) analysis demonstrates that this conductivity enhancement originates from the synergistic contribution of Mn 3d and the O 2p orbitals surrounding the oxygen vacancies, which create localized states that facilitate electron hopping throughout the structure. This finding solidified the critical role of oxygen vacancies that engineered by our synthesis approach in optimizing the electronic structure of Mn-based layered oxides. Furthermore, the calculated Na+ transport barrier for NMD (ΔENMD = 0.258 eV) is also lower than that for NM (ΔENM = 0.369 eV) (Fig. 5f) owing to the slight expansion in the c-axis direction caused by oxygen vacancies, which directly reduces the electrostatic repulsion and the spatial potential resistance of Na⁺ migration in NMD40. A detailed schematic is correspondingly illustrated in Fig. 5g.

General applicability investigation
Considering the aforementioned results, we prepared three sets of manganese based layered oxides (Na1/2Mn6/7Ni1/7O2 (NMN), Na1/2Mn6/7Cu1/7O2 (NMC), Na1/2Mn6/7Mg1/7O2 (NMM)) for investigating the broad applicability of this strategy (Supplementary Figs. 30–35 and Supplementary Tables 8–11). All samples were P2-type structure with the space group P63/mmc. The detailed rate and cycling performances of samples with and without DA addition are listed in Fig. 6a–f with the summary of performance improvements presented in Fig. 6g. All samples synthesized through the proposed strategy showed an apparent performance enhancement compared with their pristine counterparts, even by >300%. Notably, the other two ammonia-releasing sources including melamine and 5-aminotetrazole also presented similar performance-enhancing effect to DA (Supplementary Figs. 36–38), although their improvements were not as pronounced as those of DA. Therefore, this synthesis strategy can be considered as a widely applicable approach for comprehensively improving the electrochemical performances of Mn-rich P2-type LTMOs as the SIB positive electrodes.

Methods
Materials synthesis
Na1/2MnO2 (NM) with dicyandiamide (DA, 99.9%, Sinopharm Chemical Reagent) addition (NMD) was synthesized directly by high temperature solid-state reaction. Specifically, Na2CO3 (99.9%, Macklin), Mn2O3 (99.9%, Aladdin) and DA with the fixed stoichiometric proportions were mixed, following with the ball-milling. A planetary ball milling was performed with a 100 mL agate jar and balls (4–12 mm diameter) at a 20:1 ball-to-powder mass ratio. The milling process consisted of 30 minutes of forward rotation, a 5-min pause, and then 30 min of reverse rotation.

The obtained precursor after drying was pressed into the pellet, following with heating at 570 °C for 2 h and then 970 °C for 10 h to achieve complete phase formation. The obtained powder is denoted as NMD. The same procedure but without the addition of DA was used to synthesize Na1/2MnO2 (NM). All other samples Na1/2Mn6/7Ni1/7O2 (NMN), Na1/2Mn6/7Cu1/7O2 (NMC), Na1/2Mn6/7Mg1/7O2 (NMM) were synthesized with the same procedure but different added oxides while their counterparts with the addition of DA were denoted as NMND, NMCD, NMMD, etc. In addition, as a universal validation, NM with the addition of melamine (99.9%, Macklin) (NMME) and 5-aminotetrazole (99.9%, Macklin) (NMAT) were synthesized by the same procedure, respectively. All the prepared samples were quickly transferred to an argon-filled glovebox (H2O, O2 < 0.01 ppm, 25 °C) for storage.

Materials characterizations
The relative proportions of various elements were analyzed by inductively coupled plasma optical emission spectrometry (ICP-OES, prodigy XP). X-ray diffraction (XRD) patterns were collected through Bruker D8-Advance diffractometer (λ=1.54 Å, Cu Kα radiation) for identifying phase composition. GSAS II software was used to refine the result for the corresponding Rietveld data. High-resolution transmission electron microscopy (HRTEM, JEM-ARM200, JEOL) was used to check the detailed structure of materials at 300 kV. Scanning transmission electron microscopy (STEM) equipped with high-angle annular dark field (HAADF) and annular bright field (ABF) were performed on JEM-ARM200. Scanning electron microscope (SEM, Hitachi S-4800) equipped with an energy dispersive spectrometer (EDS, Oxford) was used to determine the surface morphology of samples. In-situ high-temperature XRD patterns were employed to observed the phase formation of precursors through Smartlab diffractometer (λ=1.54 Å, Cu Kα radiation) at a heating rate of 5 °C/min from room temperature to 500 °C and a scanning rate of 20°/min with 10°~50°. A simultaneous thermal analyzer (TGA, DSC I/1600LF) was used to monitor the pyrolysis process of samples in the temperature range of 30-1000 °C with a temperature increase rate of 10 °C min⁻1. Electron paramagnetic resonance spectrometry (EPR, Bruker EMX Plus) with the spins of 1.111×1012 was used to analyze the oxygen vacancies of samples at room temperature. Raman spectral analysis was performed using a confocal Raman microspectrometer (Raman, XPLORA) with an excitation wavelength of 532 nm in the range of 400–2000 cm⁻1. X-ray photoelectron spectrometer (XPS, Thermo Scientific K-Alpha) with a monochromatic Al Kα X-ray source was used to analyze the elemental valence states on the material surface. The X-ray absorption fine structure (XAFS) spectra of Mn K-edge were recorded on a commercial Laboratory-Based XAFS spectrometer (Table XAFS-500A, Specreation Instruments Co., Ltd.) with the voltage and current of 20 kV and 15 mA. The Ge (531) spherically bent crystal analyzers and the R250 mm Rowland circle were used to provide monochromatized X-ray beam. To clarify the phase transition during charging/discharging, in-situ XRD tests were performed in a specially designed cell with a Be window at 0.05 A g⁻1. For further analysis of the Na+ diffusion pathway, time-of-flight neutron powder diffraction data (NPD) were recorded by the Spallation Neutron Source Science Center in Dongguan, China (CSNS, Dongguan, China). The original data were then investigated through Rietveld analysis and maximum entropy (MEM) calculations using Z-Rietveld software. Three-dimensional nuclear density distributions are visualized by the VESTA program.

Electrochemical measurements
Coin cells (CR2025, 15.4×1.1 mm shrapnel, 16.1×0.8 mm washer) were assembled in an argon-filled glovebox (H2O, O2 < 0.01 ppm, 25 °C) to evaluate the electrochemical performances of samples. The active material, polyvinylidene fluoride (PVDF, Guangdong Canrd New Energy Technology Co., Ltd., >99.5%, 50,000 molecular weight) and conductive charcoal black (Super P, TIMCAL Graphite & Carbon, >99%) were weighed in a mass ratio of 8:1:1, and then were mixed with N-methylpyrrolidone (NMP, Shanghai Macklin Biochemical Technology Co., Ltd., >99%) for a slurry that was coated using a 150 μm squeegee on a carbon coated aluminum foil (16 µm thick with 1 µm carbon layers on both sides, 99.65% purity) and heated at 75 °C for 12 h. The mass loadings of active materials were 2–3 mg cm⁻2. Before battery assembly, sodium metal electrodes (Sinopharm Chemical Reagent Co., Ltd., chemical purity) were prepared directly in an argon-filled glovebox (H2O, O2 < 0.01 ppm, 25 °C), cutting with a 14 mm diameter cutter at a thickness of approximately 0.5–0.6 mm. For half-cell testing, Na metal foil was used as the counter electrode, glass fiber (GF-A2916, Olegeeino (Chongqing) New Materials Co., Ltd., aperture is 1.63, thickness is 0.29, porosity is >90%) as the separator and 1.0 M NaPF6 dissolved in 1,2-dimethoxyethane (DME) (80 µL, Nanjing Modges Energy Technology Co., Ltd.) as the electrolyte. For full-cell assembly, hard carbon after activation was used as the counter electrode. The hard carbon negative electrode was pre-sodiated first at 20 mA g⁻1 within 0.01–3.0 V for 3 cycles to form a strong solid electrolyte interphase (SEI) film in the half-cell for further testing. This coin cell was disassembled in an argon-filled glovebox to obtain the activated negative electrode, and then assembled with the positive electrodes to form the full cell. The theoretical capacity ratio between negative and positive electrodes was set to be 1.2. The galvanostatic charge-discharge (GCD) tests were carried out on a Land CT2001A battery test system (Wuhan LAND Electronic Co., Ltd) with a voltage range of 1.8-4.3 V at 300 K. Cyclic voltammetry (CV) curves were recorded with a CHI 760E electrochemical workstation (Chenhua Shanghai, China) at 1.8–4.3 V at the scan rates of 0.1, 0.08, 0.05, 0.03, 0.01 mV s−1 and electrochemical impedance spectroscopy (EIS) was measured in constant-potential mode with a frequency of 10⁻2~105 Hz after leaving the battery in an open-circuit state for 2 h with a total of >80 data points recorded. The full battery and pouch cells tests were also conducted with Land CT2001A battery test system for charge/discharge testing during 1.9–4.2 V. Pouch cells (single layer) were assembled by using NM/NMD and presodiated HC as the positive/negative electrodes with a N/P capacity ratio in the range of 1.02~1.15. The mass loadings of the positive and negative electrodes are 23~25 mg cm−2 and 13~15 mg cm−2, respectively, with corresponding sizes of 6.0×7.0 cm2 and 6.5×7.5 cm2. Electrolyte injection (3 g/Ah) was performed in a glove box (H2O, O2 < 0.01 ppm, 25 °C) and a vacuum sealing machine (MSK-115-III, Hefei Kejing Material Technology Co., Ltd) was used to remove gas and seal the pouch cells. The pouch cells were cycled in the potential windows within 1.5–3.6 V.

Theoretical calculations
The computational methodology in this work was based on spin-polarized density functional theory (DFT)41,42 as implemented in the Vienna ab initio simulation package (VASP)43,44. The projector augmented-wave (PAW) method was used with a plane-wave cutoff energy of 450 eV, and the Perdew-Burke-Ernzerhof (PBE)45 functional was selected for the generalized gradient approximation (GGA). A vacuum layer of 15 Å was applied to suppress inter-slab interactions. Structural optimizations for NM and NMD were conducted with a 3×3×2 k-point mesh, with convergence reached when the energy difference was less than 10⁻5 eV and the maximum force on any atom was below 0.02 eV Å⁻¹. The computational structural information files for NM and NMD are provided in Supplementary Data 1 and 2, respectively.

Peer review rounds
[{"round": 1, "reviewer_comments": {"reviewer_1": "All previous comments have been addressed, and the article is ready for publication", "reviewer_2": "1. Although the authors have provided HRTEM images and corresponding GPA analysis, conventional HRTEM has limitations in resolving atomic-scale ordering and light-element contrast. To more clearly visualize the layered structure and possible atomic-level defects, it is strongly recommended to perform scanning transmission electron microscopy (STEM) using both high-angle annular dark-field (HAADF) and annular bright-field (ABF) detectors. These techniques will enable direct imaging of atomic columns and light elements (e.g., Na), offering more definitive structural insights to support the current conclusions.\n2. While the electrochemical performance has been thoroughly evaluated, additional insight into reaction kinetics and interfacial behavior could further strengthen the work. In situ charge/discharge EIS would enable direct correlation between electrochemical behavior and structural or compositional changes during cycling. Complementary DRT analysis could help clarify ion transport and interfacial dynamics. Inclusion of such analyses would enhance the mechanistic interpretation of the observed improvements.\n3. The proposed synthesis strategy shows good potential for extension to other layered oxide cathodes. To better highlight the practical relevance of the developed material, it is recommended to include a brief comparison with typical sodium-ion battery cathodes in terms of specific capacity, average voltage, and energy density. This would help position the material’s performance within the broader landscape and clarify its advantages for practical applications.", "reviewer_3": "The manuscript proposes an innovative ammonia-assisted solid-state synthesis strategy for P2-type layered transition-metal oxides, demonstrating enhanced performance in sodium-ion battery cathodes. The concept of leveraging gas-phase chemistry to modulate crystallization kinetics and suppress structural degradation is an innovative perspective for structural control in the field. However, the current study requires substantial strengthening in mechanistic validation, methodological rigor, and structural analysis to fully support its claims. Below, I outline major concerns that necessitate rejection in the present form, though the work holds potential for high impact if these issues are adequately addressed.\n1. The proposed link between ordered TM layer arrangement and improved structural stability/kinetics requires deeper mechanistic justification. In P2-type materials, TM ordering may induce ordered Na-layer vacancies, which could restrict Na+ mobility rather than enhance it. To strengthen this premise, quantify vacancy distribution via neutron diffraction or DFT. Discuss how specific ordering types impact stability.\n2. While the DA-modified method appears efficient, the critique of established routes e.g., co−precipitation, sol−gel as universally suffering \"poor phase purity and sluggish kinetics\" overgeneralizes. These methods are industrially viable with optimized protocols. To highlight novelty: Directly compare phase purity/yield with conventional solid-state or hydrothermal synthesis. Emphasize scalability or cost benefits unique to this approach.\n3. Lattice fringe analysis alone is insufficient. Aberration-corrected STEM e.g., along [010]/[001] zone axes is essential to resolve atomic-scale ordering/vacancies. SAED patterns from multiple zone axes should supplement XRD to confirm hexagonal symmetry. High-temperature in situ XRD is critical to correlate ammonia release with crystallization dynamics and structural ordering.\n4. The claim that ammonia suppresses irreversible oxygen redox lacks direct proof. In situ mass spectrometry or isotopic labeling could trace ammonia evolution and oxygen activity. Enhanced XRD peak intensity e.g.,002 is suggestive but not conclusive. In situ XRD peak sharpening/FWHM evolution during synthesis would provide stronger evidence. Mn valence/oxygen vacancy analysis via XPS is semi-quantitative. Supplement with EELS or Mn K-edge XANES for robust quantification.\n5. EIS: Provide equivalent circuit models, fitting errors, and Nyquist plots in 1:1 aspect ratio.\n6. Coin-cell data are preliminary. Soft-pouch cells with explicit N/P ratios, electrolyte volume, and cycling protocols are needed for practical relevance.\n7. The P2-to-P´2 phase transition during low-voltage discharge linked to Jahn−Teller distortion is a key degradation pathway for Mn-based LTMOs. Its absence in in situ XRD data is puzzling.\n8. The \"Analysis of Na⁺ transport mechanism\" section duplicates earlier EIS/CV/GITT discussions.\n9. Oxygen Vacancy Modeling: Computational models must incorporate oxygen vacancies noted in XPS to validate energy barriers.\n10. Electronic Structure: Desodiation DOS analysis is essential to explain enhanced kinetics.\nThe manuscript presents a creative synthesis strategy with potential to advance LTMO cathode design. However, the current evidence is insufficient to support the proposed mechanisms. The current presentation does not meet the high standards of Nature Communications for novelty, rigor, and impact. I recommend rejection without further review."}, "author_response": "Reviewer #1:\n“All previous comments have been addressed, and the article is ready for\npublication.”\nR (Response): We sincerely thank you for your careful review and valuable\ncomments. We greatly appreciate your recognition of our work.\nReviewer #2:\n“1. Although the authors have provided HRTEM images and corresponding GPA\nanalysis, conventional HRTEM has limitations in resolving atomic-scale ordering and\nlight-element contrast. To more clearly visualize the layered structure and possible\natomic-level defects, it is strongly recommended to perform scanning transmission\nelectron microscopy (STEM) using both high-angle annular dark-field (HAADF) and\nannular bright-field (ABF) detectors. These techniques will enable direct imaging of\natomic columns and light elements (e.g., Na), offering more definitive structural\ninsights to support the current conclusions. ”\nR: Thanks for the comments. We appreciate your input to improve the quality of this\nmanuscript. We have followed the reviewer’s suggestion and made the corresponding\nrevisions by supplementing the required evidences. The data are added in Fig. 1i\n(STEM-HAADF) and 1j (ABF), and both images confirm that NMD has a highly\ncrystalline lamellar structure (P2 type), which is consistent with XRD/NPD\nrefinement results. To make this point clear, the following lines and data are updated:\nPage 7:\nScanning transmission electron microscopy (STEM) equipped with high-angle\nannular dark field (HAADF) (Fig. 1i) and annular bright field (ABF) (Fig. 1j) is\nemployed to analyze the atomic-level crystal structure of NMD. In the HAADF mode,\nthe light dots signify the transition metal atoms while the O and Na atoms are\nrevealed within the dark columns, whereas the manifestation is reversed in the ABF\nmode17,20. According to the STEM-HAADF image, the TM-TM distance is 0.273 nm\nR 1\ncorresponding to the (100) plane of the P2 structure. Selected area electron diffraction\n(SAED) pattern along the [001] axis substantiates the single-crystalline hexagonal\nstructure of the bulk phase of NMD (Fig. 1k).\nPage 21:\nScanning transmission electron microscopy (STEM) equipped with high-angle\nannular dark field (HAADF) and annular bright field (ABF) were performed on\nJEM-ARM200.\nPage 5:\nFig. 1│Structural and morphological characterizations. e, g High-resolution TEM\nimages of (e) NM and (g) NMD with the insets showing the corresponding FFT\npatterns. f, h The corresponding strain distributions of (e) and (g). i, j Scanning\ntransmission electron microscopy equipped with (i) high-angle annular dark field\n(STEM-HAADF) and (j) annular bright field (ABF) for NMD. k The selected area\nelectron diffraction (SAED) images for NMD. l EDS mappings of NMD: Na, Mn, O.\n“2.While the electrochemical performance has been thoroughly evaluated, additional\ninsight into reaction kinetics and interfacial behavior could further strengthen the\nwork. In situ charge/discharge EIS would enable direct correlation between\nelectrochemical behavior and structural or compositional changes during cycling.\nComplementary DRT analysis could help clarify ion transport and interfacial\ndynamics. Inclusion of such analyses would enhance the mechanistic interpretation of\nthe observed improvements.”\nR 2\nR: Thanks for the comments. We have followed the reviewer’s suggestion and added\nthe corresponding data. In-situ EIS and DRT techniques were performed to further\ncomprehend the reaction kinetics and interfacial behavior of the materials. To make\nthis point clear, the following lines and data are updated:\nPage 12:\nIn addition, the in-situ EIS and the distribution of relaxation times (DRT) analysis\nwere performed to reveal the evolution of the internal states of NM and NMD at\ndifferent potentials. In-situ EIS analysis reveals that the charge transfer resistance\ninitially decreases and then increases (Supplementary Fig. 20), fluctuating with the Na\nconcentration in the electrode during the early charge–discharge process28. Notably,\nNMD exhibits a lower transfer resistance after charge-discharge cycles\n(Supplementary Fig. 21), highlighting its superior ion transport kinetics. Further DRT\nanalysis (Supplementary Fig. 22) shows lower intensity values across all regions for\nNMD compared to NM, indicating its reduced electrochemical impedance29,30. The\nreduced impedance and enhanced ion transport kinetics in NMD can be ascribed to\noxygen vacancies that weaken the Na-O electrostatic interaction, thereby expanding\ninterlayer spacing and promoting both Na-ion diffusion and electronic conductivity.\nPage S25:\nSupplementary Fig. 20 In-situ EIS measurements of NM and NMD half-cells during\ncharging/discharging within 1.8–4.3V.\nPage S26:\nR 3\nSupplementary Fig. 21 The Nyquist plots of (a) initial, (b) charging to 4.3V and (c)\ndischarging to 1.8 V during in-situ EIS measurements.\nPage S27:\nSupplementary Fig. 22 Corresponding contour plots of the distribution of relaxation\ntimes (DRT) of in-situ EIS.\n“3. The proposed synthesis strategy shows good potential for extension to other\nR 4\nlayered oxide cathodes. To better highlight the practical relevance of the developed\nmaterial, it is recommended to include a brief comparison with typical sodium-ion\nbattery cathodes in terms of specific capacity, average voltage, and energy density.\nThis would help position the material’s performance within the broader landscape\nand clarify its advantages for practical applications.”\nR: Thanks for the comments. We have followed reviewer’s suggestion and made the\ncorresponding revisions. A brief comparison with typical sodium-ion battery cathodes\nin terms of specific capacity, average voltage and energy density have been\nsupplemented in Supplementary Table 6. To make this point clear, the following lines\nand data are updated:\nPage S49:\nSupplementary Table 6 Comparisons of discharge and energy density of the\nNMD||HC full cell in this work with recently published full cells.\nDischarge Average\nCathode Voltage Energy density\ncapacity (mAh voltage Refs.\nmaterials range (V) (Wh kg⁻1)\ng⁻1) (V)\nCoin-full cells: 244.4;\nNa MnO with 130.7, 50 mA 1.9-4.2; 450 (based on the\n1/2 2\n2.7 This work\ndicyandiamide g⁻1 2.6 mass of cathode);\nPouch cells: 210.3\nNa Li Fe Nat. Energy,\n0.67 0.1 0.37\n-- -- 2.0-4.3 230\nMn0.5 O 202418\n3 2\nNa Ni Mn Nat. Comm.,\n0.78 0.31 0.\n60.4, 37 mA g⁻1 3.4 2.3-4.14 202.0\nNb O 20225\n67 0.02 2\nNa Mg Mn 155.0, 70 mA Angew. Chem.\n0.75 0.25\n2.39 1.4-3.9 233.0\nO g⁻1 Int. Ed., 20237\n0.75 2\nNa Ni Fe M\n0.8 0.3 0.2\nAdv. Energy\nn Li Mg C 114.8, 14 mA g⁻1 -- 1.0-4.2 208.0\n0.3 0.1 0.02 Mater., 202319\na Sb O\n0.05 0.03 2\nNa Li Mn Angew. Chem.\n0.72 0.24 0.\n233.8, 20 mA ⁻1 -- 1.6-4.3 211.0\nO Int. Ed., 202420\n76 2\nNa Cu Li J. Am. Chem.\n0.8 0.22 0.08\n117.9, 26 mA g⁻1 2.78 2.0-4.3 228.4\nMn O Soc., 20238\n0.67 2\nNa Ni Mn Adv. Mater.,\n2/3 1/3 1/3\n99.0, 18 mA g⁻1 3.33 2.0-4.3 215.9\nTi O 202021\n1/3 2\nAngew. Chem.\nNa V (PO ) /C 110, 1/3 C 2.9 2.5-3.9 --\n3 2 4 3 Int. Ed., 202422\nHigh 101.9, 100 mA 2.5 1.75-4.25 81.4 Adv. Energy\nR 5\nentropy-Na Fe ( g-1 Mater., 202523\n4 3\nPO ) P O\n4 2 2 7\nHigh\nAdv. Funct.\nentropy-Na V ( 133, 14.3 mA g-1 4.0 1.4-4.4 342\n2 2 Mater., 202524\nPO )\n4 3\nNa Ni[Fe(CN\n1.34 54.5, 10 mA g-1 3.32 2.0-4.4 -- Small, 202325\n) ]\n6 0.92\nFeVO-Prussian J. Am. Chem.\n130.0, 50 mA g-1 2.64 1.6-4.0 259.7\nblue analogue Soc., 202526\nPrussian blue 370 (based on the ACS Energy\n112.3, 0.2 C -- 1.5-3.7\nanalogs mass of cathode) Letters, 202527\nReviewer #3:\n“The manuscript proposes an innovative ammonia-assisted solid-state synthesis\nstrategy for P2-type layered transition-metal oxides, demonstrating enhanced\nperformance in sodium-ion battery cathodes. The concept of leveraging gas-phase\nchemistry to modulate crystallization kinetics and suppress structural degradation is\nan innovative perspective for structural control in the field. However, the current\nstudy requires substantial strengthening in mechanistic validation, methodological\nrigor, and structural analysis to fully support its claims. Below, I outline major\nconcerns that necessitate rejection in the present form, though the work holds\npotential for high impact if these issues are adequately addressed.”\nR: Thanks for the comments. We appreciate your input to improve the quality of this\nmanuscript. We have followed your suggestion and made the corresponding revisions\nby supplementing the required evidences. Please check the details in the following\nsection for point-to-point response.\n“1. The proposed link between ordered TM layer arrangement and improved\nstructural stability/kinetics requires deeper mechanistic justification. In P2-type\nmaterials, TM ordering may induce ordered Na-layer vacancies, which could restrict\nNa+ mobility rather than enhance it. To strengthen this premise, quantify vacancy\ndistribution via neutron diffraction or DFT. Discuss how specific ordering types\nimpact stability.”\nR 6\nR: Thanks for the comments. We have followed the reviewer’s suggestion and made\nthe corresponding revisions. We remove the statement “ordered TM layer\narrangement” and focus on the oxygen vacancy-mediated stabilization mechanism.\nThe selected area electron diffraction (SAED) shows diffraction spots of the basic\nlattice without additional ordered signals. Hence, the intrinsic disorder combined with\noxygen vacancies avoids detrimental Na vacancy ordering while maintaining\nstructural flexibility. Geometric phase analysis (GPA) analysis (Fig. 1f and h) shows\nthat NMD has uniform lattice strain, which is significantly better than that of NM and\neffectively inhibits crack extension. DFT-calculated Na+ migration pathways (Fig. 5b,\ng) demonstrate oxygen vacancies create low-energy interstitial channels (barrier:\n0.258 eV). Moreover, the analysis of DOS (Fig. 5c, d and Supplementary Fig. 29)\nreveals that NMD exhibits a significant increase in peak broadening and a continuous\nand higher DOS near the Fermi energy level compared to NM, demonstrating electron\ndelocalization around Mn3+ due to oxygen vacancies accompanied by an increase in\nthe electrical conductivity. The disordered distribution of oxygen vacancies at the\nadequate concentration allows the material to obtain enhanced stability and kinetics.\nTo make this point clear, the following lines and data are updated:\nPage 3:\nMeanwhile, the synthesis of high-quality LTMOs with stable ion tunnels can not only\nstabilize the structure but also accelerate Na+ desodiation with high reaction\nkinetics12,13.\nPage 4:\nIn addition to the new working mechanism of gas-phase assisted solid state reaction\nrevealed in this strategy, its broad universality for the synthesis of highly active\nP2-type LTMOs cathodes with various composition attracts our interests.\nPage 7:\nScanning transmission electron microscopy (STEM) equipped with high-angle\nannular dark field (HAADF) (Fig. 1i) and annular bright field (ABF) (Fig. 1j) is\nemployed to analyze the atomic-level crystal structure of NMD. In the HAADF mode,\nthe light dots signify the transition metal atoms while the O and Na atoms are\nR 7\nrevealed within the dark columns, whereas the manifestation is reversed in the ABF\nmode17,20. According to the STEM-HAADF image, the TM-TM distance is 0.273 nm\ncorresponding to the (100) plane of the P2 structure. Selected area electron diffraction\n(SAED) pattern along the [001] axis substantiates the single-crystalline hexagonal\nstructure of the bulk phase of NMD (Fig. 1k).\nPage 17:\nDensity-functional theory (DFT) calculations were performed to further investigate the\ndiffusion mechanism of Na+ using a computational model that introduces oxygen\nvacancies for NMD (Supplementary Fig. 28). Based on the analysis for electronic\ndensity of states (DOS) at different desodiation states (Fig. 5c-e and Supplementary\nFig. 29), NMD demonstrates the dramatically increased number of electronic states at\nthe Fermi level and significantly reduced band gap compared to NM, even in different\ndesodiation states38, 39, which confirms its improved electronic conductivity. The\npartial density of states (PDOS) analysis demonstrates that this conductivity\nenhancement originates from the synergistic contribution of Mn 3d and the O 2p\norbitals surrounding the oxygen vacancies, which create localized states that facilitate\nelectron hopping throughout the structure. This finding solidifies the critical role of\noxygen vacancies that engineered by our synthesis approach in optimizing the\nelectronic structure of Mn-based layered oxides. Furthermore, the calculated Na+\ntransport barrier for NMD (ΔE = 0.258 eV) is also lower than that for NM (ΔE =\nNMD NM\n0.369 eV) (Fig. 5f) owing to the slight expansion in the c-axis direction caused by\noxygen vacancies, which directly reduces the electrostatic repulsion and the spatial\npotential resistance of Na⁺ migration in NMD40.\nPage 21:\nScanning transmission electron microscopy (STEM) equipped with high-angle\nannular dark field (HAADF) and annular bright field (ABF) were performed on\nJEM-ARM200.\nPage 5:\nR 8\nFig. 2│Structural and morphological characterizations. e, g High-resolution TEM\nimages of (e) NM and (g) NMD with the insets showing the corresponding FFT\npatterns. f, h The corresponding strain distributions of (e) and (g). i, j Scanning\ntransmission electron microscopy equipped with (i) high-angle annular dark field\n(STEM-HAADF) and (j) annular bright field (ABF) for NMD. k The selected area\nelectron diffraction (SAED) images for NMD. l EDS mappings of NMD: Na, Mn, O.\nPage S34:\nSupplementary Fig. 29 The calculated total density of states (TDOS) and partial\ndensity of states (PDOS) of the Mn 3d and O 2p orbitals under different desodiation\nstates.\nR 9\n“2.While the DA-modified method appears efficient, the critique of established routes\ne.g., co−precipitation, sol−gel as universally suffering \"poor phase purity and sluggish kinetics\" overgeneralizes. These methods are industrially viable with optimized protocols. To highlight novelty: Directly compare phase purity/yield with conventional solid-state or hydrothermal synthesis. Emphasize scalability or cost benefits unique to this approach.”\nR: Thanks for the comments. We have followed reviewer’s suggestion and made the\ncorresponding revisions. Here, we concur with the reviewer's suggestion and the\nmanuscript emphasize the improvement of the a conventional solid-phase method. To\nhighlight the novelty, we compare the present method with conventional solid-state or\nhydrothermal synthesis in terms of phase purity/yield, emphasizing the scalability or\ncost advantages specific to this method. To make this point clear, the following lines\nand data are updated:\nPage 18:\nConsidering the aforementioned results, we prepared three sets of manganese based\nlayered oxides (Na Mn Ni O (NMN), Na Mn Cu O (NMC),\n1/2 6/7 1/7 2 1/2 6/7 1/7 2\nNa Mn Mg O (NMM)) for investigating the broad applicability of this strategy\n1/2 6/7 1/7 2\n(Supplementary Figs. 30-35 and Supplementary Tables 8-11).\nPage S54:\nSupplementary Table 11 Comparison of this method with traditional solid-state\nmethod and hydrothermal method.\nTraditional solid-state Hydrothermal\nThis work\nsynthesis synthesis\nHigh (Possible\nPhase\nHigh presence of High\npurity\nimpurities)\nYield High High Low28\nIt is readily scalable and highly There are certain\nIt is readily scalable\nScalabil suitable for industrial production, challenges with the\nand highly suitable for\nity which is universal without requiring relatively high\nindustrial production.\nadditional equipment. equipment costs.\nR1 0\nThere is no need to introduce any It requires high\nadditional metal elements. In terms of pressure reactor and\nUsually introduces\nthe cost of introducing additional reprocessing, and\nCost additional metal-like\nsubstances, the introduction of some experiments\nbenefits elements (Mostly two\ndicyandiamide is less expensive than still require\nor more)19.\nthe introduction of metals, which are high-temperature\nmostly more than two. sintering29.\n“3.Lattice fringe analysis alone is insufficient. Aberration-corrected STEM e.g., along\n[010]/[001] zone axes is essential to resolve atomic-scale ordering/vacancies. SAED\npatterns from multiple zone axes should supplement XRD to confirm hexagonal\nsymmetry. High-temperature in situ XRD is critical to correlate ammonia release with\ncrystallization dynamics and structural ordering.”\nR: Thanks for the comments. We have followed reviewer’s suggestion and made the\ncorresponding revisions. Scanning transmission electron microscopy (STEM)\nequipped with high-angle annular dark field (HAADF) (Fig. 1i) and annular bright\nfield (ABF) (Fig. 1j) is employed to analyze the atomic-level crystal structure of\nNMD. STEM-HAADF, ABF and SAED images confirm that NMD has a highly\ncrystalline lamellar structure (P2 type), consistent with XRD/NPD refinement results.\nIn-situ XRD technique from room temperature to 500 ℃ was also used to investigate\nthe detailed influence of DA addition on the growth mechanism of LTMOs, which is\nthe main temperature range for DA decomposition known from the TGA curves\n(Supplementary Fig. 7). The in situ XRD peak sharpening/full width at half maxima\n(FWHM) evolution during synthesis was further analyzed to complement the\nappearance of the (002) peak (Supplementary Fig. 5). DA gradually releases ammonia\nwith increasing temperature while seizing lattice oxygen atoms to obtain oxygen\nvacancies. The generation of oxygen vacancies reduces the diffusion resistance of Na+\nand Mn3+ in the solid-phase reaction, promoting the early formation of the\nintermediate phase. From the Fig. 2a-c, the simultaneous observation of the\nweakening of the Na CO and Mn O diffraction peaks was observed until the\n2 3 2 3\npresence of the (002) lattice diffraction peaks was detected in the NMD at 320 °C,\nwhich is earlier than 400 °C in the NM. The above evidence proves that ammonia\nR1 1\nrelease generates oxygen vacancies, which broadens the phase transition temperature\ndomain and enhances the crystal structure kinetics, ultimately leading to a\nhigh-quality P2 phase structure. To make this point clear, the following lines and data\nare updated:\nPage 7:\nScanning transmission electron microscopy (STEM) equipped with high-angle\nannular dark field (HAADF) (Fig. 1i) and annular bright field (ABF) (Fig. 1j) is\nemployed to analyze the atomic-level crystal structure of NMD. In the HAADF mode,\nthe light dots signify the transition metal atoms while the O and Na atoms are\nrevealed within the dark columns, whereas the manifestation is reversed in the ABF\nmode17,20. According to the STEM-HAADF image, the TM-TM distance is 0.273 nm\ncorresponding to the (100) plane of the P2 structure. Selected area electron diffraction\n(SAED) pattern along the [001] axis substantiates the single-crystalline hexagonal\nstructure of the bulk phase of NMD (Fig. 1k).\nPage 7:\nTo investigate the detailed influence of DA addition on the growth mechanism of\nLTMOs, the structural evolution from precursor to the final crystalline form was\nanalyzed with in-situ XRD technique from room temperature to 500 ℃ (Fig. 2a-c). The\n(002) peak of sodium manganite at 15.6° can be detected at a low temperature of 320 ℃\nfor NMD (vs. 400 ℃ for NM), demonstrating this strategy significantly decreases the\nformation temperature of Na MnO structure and allows the fully growth of crystals\n1/2 2\nwith more ordered atomic arrangement during wide temperature range. The (001) peak\nof Na CO at 14.9° progressively disappears while the (222) peak of Mn O at 33° also\n2 3 2 3\ngradually diminishes with increasing temperatures, indicating its transformation into\nNa MnO . Meanwhile, a gradual sharpening of the (002) peak with the increase of the\n1/2 2\nhalf-peak full width is also observed (Supplementary Fig. 5).\nPage 21:\nScanning transmission electron microscopy (STEM) equipped with high-angle\nannular dark field (HAADF) and annular bright field (ABF) were performed on\nJEM-ARM200.\nR1 2\nPage 21:\nIn-situ high-temperature XRD patterns were employed to observed the phase\nformation of precursors through Smart lab diffractometer (λ = 1.54 Å, Cu Kα\nradiation) at a heating rate of 5 °C/min from room temperature to 500 °C and a\nscanning rate of 20 °/min with 10°~50°.\nPage 5:\nFig. 3│Structural and morphological characterizations. e, g High-resolution TEM\nimages of (e) NM and (g) NMD with the insets showing the corresponding FFT\npatterns. f, h The corresponding strain distributions of (e) and (g). i, j Scanning\ntransmission electron microscopy equipped with (i) high-angle annular dark field\n(STEM-HAADF) and (j) annular bright field (ABF) for NMD. k The selected area\nelectron diffraction (SAED) images for NMD. l EDS mappings of NMD: Na, Mn, O.\nPage 9:\nR1 3\nFig. 2│Structural formation analysis. a Y-shifted stacked line plots of NMD/NM at\ndifferent temperatures for in-situ XRD data. b Magnified XRD data of the selected\nregion in (a) with the highlighting (002) peaks. c Corresponding contour plots.\nPage S10:\nSupplementary Fig. 5 The (002) peak full width at half maxima (FWHM) evolution\nof in-situ XRD data at different temperatures.\n“4. The claim that ammonia suppresses irreversible oxygen redox lacks direct proof.\nIn situ mass spectrometry or isotopic labeling could trace ammonia evolution and\noxygen activity. Enhanced XRD peak intensity e.g.,002 is suggestive but not\nconclusive. In situ XRD peak sharpening/FWHM evolution during synthesis would\nR1 4\nprovide stronger evidence. Mn valence/oxygen vacancy analysis via XPS is\nsemi-quantitative. Supplement with EELS or Mn K-edge XANES for robust\nquantification.”\nR: Thanks for the comments. We have followed reviewer’s suggestion and made the\ncorresponding revisions. We agree with the reviewers' observation that in situ mass\nspectrometry or isotope labeling techniques can effectively track ammonia evolution\nand oxygen activity. However, it has to be emphasized that the focus of this study is\nthe additional oxygen vacancies arise due to the release of ammonia, thereby\neffectively suppressing irreversible oxygen oxidation reactions to stabilize cation\nmigration. In addition to EPR and XPS, we have supplemented this with Mn-K-edge\nX-ray absorption near edge structure (XANES) spectroscopy to confirm this. And the\nin situ XRD peak sharpening/full width at half maxima (FWHM) evolution during\nsynthesis was further analyzed to complement the appearance of the (002) peak\n(Supplementary Fig. 5). To make this point clear, the following lines and data are\nupdated:\nPage 8:\nMeanwhile, a gradual sharpening of the (002) peak with the increase of the half-peak\nfull width is also observed (Supplementary Fig. 5).\nPage 10:\nEx situ X-ray absorption fine structure (XAFS) spectra were collected to reveal the\nchange of valance states for metal elements between samples. The X-ray absorption\nnear edge structure (XANES) at the Mn K-edge and the corresponding Fourier\ntransformed of extended X-ray absorption fine structure (FT-EXAFS) spectra were\npresented in Fig. 2e and f. Compared to NM, NMD shows a left-shift to the lower\nenergy for its Mn K-edge (Fig. 2e), indicating the oxidation of Mn with the decreased\nvalence state, which coincides with the enlarged interatomic distances within the first\nMn-O coordination shell that is presented with the increased Mn-O peak in Fig. 2f21.\nThis observation is also consistent with the XPS result of NMD that presents the\nincreased Mn3+/Mn4+ ratio compared with NM. Meanwhile, the decreased second\nMn-Mn coordination shell peak of NMD (Fig. 2f) suggests the reduced interatomic\nR1 5\ndistances of Mn, which demonstrates the contraction of MnO unit.\ntransform (WT) contour plot of EXAFS spectra of NMD (Fig. 2g) exhibits two\ndistinct scattering peaks of Mn-O and Mn-Mn coordination shells located at (5.6 Å⁻1,\n1.4 Å) and (5.6 Å⁻1, 2.3Å), respectively. Correspondingly, the center of these\nscattering peaks is shifted with the same trend of Mn-O and Mn-Mn peaks of\nFT-EXAFS for those in NM (5.4 Å⁻1, 1.4 Å and 5.7 Å⁻1, 2.3 Å) (Fig. 2h)22. The\nMn-Mn coordination shell in NMD with low R-value demonstrates the locally\ndisordered Mn-Mn coordination environment that is caused by the introduction of\noxygen vacancies. Meanwhile, the well-maintained shape without change from NM to\nNMD suggests the uniform introduction of oxygen vacancies into NMD bulk.\nThe\naforementioned analysis confirms the introduction of oxygen vacancies into NMD\nalong with electronic/structural changes that are the basis for its electrochemical\nperformance enhancement.\nFig. 2│Structural formation analysis. d Mn 2p spectra. e O 1s spectra of NM and\nNMD. f Mn K-edge XANES spectra and (g) the corresponding Fourier-transformed\nEXAFS spectra. h, i Wavelet transform (WT) contour plots of NMD and NM.\nPage S10:\nR1 6\nSupplementary Fig. 5 The (002) peak full width at half maxima (FWHM) evolution\nof in-situ XRD data at different temperatures.\n“5. EIS: Provide equivalent circuit models, fitting errors, and Nyquist plots in 1:1\naspect ratio.”\nR: Thanks for the comments. We have followed reviewer’s suggestion and made the\ncorresponding revisions. We performed an equivalent circuit fit to the EIS data and\nupdated Fig. 3c, Supplementary Table 4, and related descriptions. To make this point\nclear, the following lines and data are updated:\nPage 12:\nCorrespondingly, a 2.6-fold increase of derived Na+ diffusion coefficient (D ) value\nNa+\nfor NMD (9.438 × 10⁻13 cm2 s⁻1) compared with NM (3.658× 10⁻13 cm2 s⁻1) (Fig. 3c\nand Supplementary Fig. 18) also confirms the rapid Na+ diffusion rate within NMD\nthat is consistent well with CV results, facilitating its high discharge capacity under\nelevated current densities.\nPage 14:\nFig. 3│Electrochemical performances of cathodes and full cells. c Impedance\nfitting results with the inset showing the calculated Na+ diffusion coefficients and\nequivalent circuit model.\nR1 7\nFig. 3c inset table:\nSamples NMD Fitting error (%) NM Fitting error (%)\nRa (Ohm) 3.814 1.562 5.563 3.105\ns\nR b(Ohm) 6.341 2.543 23.199 2.406\nCEI\nR c(Ohm) 45.932 1.102 92.487 2.229\nct\nσd 3.416 5.761\n𝐷 e (cm2 s⁻1) 9.438 × 10⁻13 3.658× 10⁻13\n𝑁𝑎+\n“6. Coin-cell data are preliminary. Soft-pouch cells with explicit N/P ratios,\nelectrolyte volume, and cycling protocols are needed for practical relevance.”\nR: Thanks for the comments. We have followed reviewer’s suggestion and made the\ncorresponding revisions. We have assembled and tested single-layer pouch cells using\nthe NMD cathode and HC anode to evaluate its performance under realistic conditions.\nThe detailed cell parameters (N/P: 1.05 and electrolyte volume: 2.0 g Ah⁻1), cycling\nprotocols and results are presented in Supplementary Fig. 27. The key findings from\nthe pouch cell testing unequivocally confirm the practical promise of our material. To\nmake this point clear, the following lines and data are updated:\nPage 13:\nMoreover, pouch cells were assembled by using presodiated HC and NMD to assess\ntheir practical cycling stability within 1.5–3.6 V (Supplementary Fig. 27). NMD||HC\npouch cell presents a reverse capacity of 140.0 mAh g⁻1 at 0.1 A g⁻1, which is higher\nthan NM||HC pouch cell (86.9 mAh g⁻1). The corresponding GCD curves in\nSupplementary Fig. 27b further shows the excellent electrochemical reversibility of\nthe NMD electrode. The NMD||HC pouch cell also presents high capacity retention of\n82.5% for 200 cycles and high energy density of 210.3 Wh kg⁻1 at 0.5 A g⁻1, verifying\nits great practical potential as the SIB cathodes.\nPage 23:\nPouch cells were assembled by usin NM/NMD and presodiated HC as the\ncathode/anode with a N/P capacity ratio in the range of 1.02~1.15. The mass loading\nof the cathode and anode are 23~25 mg cm⁻2 and 13~15 mg cm⁻2, respectively, with\nR1 8\ncorresponding sizes of 6.0 × 7.0 cm2 and 6.5 × 7.5 cm2. The pouch cells were cycled\nin the potential windows within 1.5–3.6 V.\nPage S32:\nSupplementary Fig. 27 Electrochemical performance of NM||HC and NMD||HC\npouch cells. a Long-term cycling performance and Coulombic efficiency at 0.5 A g⁻1.\nb The first 20 cycles GCD curves within 1.5–3.6 V at 0.5 A g⁻1.\n“7. The P2-to-P´2 phase transition during low-voltage discharge linked to\nJahn−Teller distortion is a key degradation pathway for Mn-based LTMOs. Its\nabsence in in situ XRD data is puzzling.”\nR: Thanks for the comments. We have followed reviewer’s suggestion and made the\ncorresponding revisions. Firstly, the appearance of the P'2 phase diffraction peak is\nnot observed during the first cycle of charging and discharging, but rather the\nappearance of the OP4 phase (001) peak at 17.4° (Fig. 4a, b, d and e). The P'2 phase\ntransition strongly depends on sodium vacancy ordering (e.g., sawtooth chains), and a\nlow concentration of oxygen vacancies in NM can disrupt the Na-layer electrostatic\nperiodicity, which blocks the full-domain slip and P'2 nucleation. NMD with adequate\noxygen vacancy concentration provide greater structural reversibility by lowering the\nvalence state of Mn, reducing oxidation state contraction and buffering interlayer\nstresses. To make this point clear, the following lines and data are updated:\nPage 15:\nDuring the high-voltage charging stage, the disappearance of the P2 phase (002) peak\nis accompanied by the emergence of a new diffraction peak at 17.4°, which is\nattributed to the (001) crystal plane of the OP4 phase. The original lattice symmetry\nR1 9\nand periodicity are disrupted due to the presence of oxygen vacancies, which are more\nfavorable for the transition to the OP4 phase during Na+ extraction.\nPage 16:\nNMD achieves minimum structural changes and superior structural reversibility over\nNM during Na+ de-intercalation due to oxygen vacancies to buffer interlayer stresses\nin NMD.\n“8. The \"Analysis of Na⁺ transport mechanism\" section duplicates earlier\nEIS/CV/GITT discussions.”\nR: Thanks for the comments. We have followed reviewer’s suggestion and made the\ncorresponding revisions. To make this point clear, the following lines and data are\nupdated:\nPage 16:\nFurthermore, the corresponding three-dimensional nuclear density distribution\ncalculated by MEM (Fig. 5b) shows Na+ with positive scattering lengths migrate along\nthe “Z” path through oxygen sites or vacancies in NMD, which provides shorter Na+\ndiffusion distances and more active sites than NM43.\nPage 17:\nDensity-functional theory (DFT) calculations were performed to further investigate the\ndiffusion mechanism of Na+ using a computational model that introduces oxygen\nvacancies for NMD (Supplementary Fig. 28). Based on the analysis for electronic\ndensity of states (DOS) at different desodiation states (Fig. 5c-e and Supplementary\nFig. 29), NMD demonstrates the dramatically increased number of electronic states at\nthe Fermi level and significantly reduced band gap compared to NM, even in different\ndesodiation states38, 39, which confirms its improved electronic conductivity. The\npartial density of states (PDOS) analysis demonstrates that this conductivity\nenhancement originates from the synergistic contribution of Mn 3d and the O 2p\norbitals surrounding the oxygen vacancies, which create localized states that facilitate\nelectron hopping throughout the structure. This finding solidifies the critical role of\noxygen vacancies that engineered by our synthesis approach in optimizing the\nR2 0\nelectronic structure of Mn-based layered oxides. Furthermore, the calculated Na+\ntransport barrier for NMD (ΔE = 0.258 eV) is also lower than that for NM (ΔE =\nNMD NM\n0.369 eV) (Fig. 5f) owing to the slight expansion in the c-axis direction caused by\noxygen vacancies, which directly reduces the electrostatic repulsion and the spatial\npotential resistance of Na⁺ migration in NMD40.\nPage S34:\nSupplementary Fig. 29 The calculated total density of states (TDOS) and partial\ndensity of states (PDOS) of the Mn 3d and O 2p orbitals under different desodiation\nstates.\n“9. Oxygen Vacancy Modeling: Computational models must incorporate oxygen\nvacancies noted in XPS to validate energy barriers.”\nR: Thanks for the comments. We have followed reviewer’s suggestion and made the\ncorresponding revisions. In the theoretical calculation part, the NMD is modeled with\noxygen vacancies, where the calculated Na+ transport barrier of NMD is lower than\nthat of the NM by 0.111 eV (Fig. 5f), which is due to the generation of more oxygen\nvacancies. To make this point clear, the following lines and data are updated:\nPage 17:\nR2 1\nDensity-functional theory (DFT) calculations were performed to further investigate the\ndiffusion mechanism of Na+ using a computational model that introduces oxygen\nvacancies for NMD (Supplementary Fig. 28).\nPage S33:\nSupplementary Fig. 28 The model used for first-principles calculations (NMD\ninvolves oxygen vacancies for comparison).\n“10. Electronic Structure: Desodiation DOS analysis is essential to explain enhanced\nkinetics.”\nR: Thanks for the comments. We have followed reviewer’s suggestion and made the\ncorresponding revisions. Complementary calculations of the density of states of the\nmaterial for different charging and discharging states were carried out to account for\nthe enhancement kinetics in the desodiated and sodiated states. To make this point\nclear, the following lines and data are updated:\nPage 17:\nBased on the analysis for electronic density of states (DOS) at different desodiation\nstates (Fig. 5c-e and Supplementary Fig. 29), NMD demonstrates the dramatically\nincreased number of electronic states at the Fermi level and significantly reduced band\ngap compared to NM, even in different desodiation states38, 39, which confirms its\nimproved electronic conductivity. The partial density of states (PDOS) analysis\ndemonstrates that this conductivity enhancement originates from the synergistic\ncontribution of Mn 3d and the O 2p orbitals surrounding the oxygen vacancies, which\ncreate localized states that facilitate electron hopping throughout the structure. This\nfinding solidifies the critical role of oxygen vacancies that engineered by our synthesis\nR2 2\napproach in optimizing the electronic structure of Mn-based layered oxides.\nFurthermore, the calculated Na+ transport barrier for NMD (ΔE = 0.258 eV) is also\nNMD\nlower than that for NM (ΔE = 0.369 eV) (Fig. 5f) owing to the slight expansion in\nNM\nthe c-axis direction caused by oxygen vacancies, which directly reduces the\nelectrostatic repulsion and the spatial potential resistance of Na⁺ migration in NMD40.\nPage S34:\nSupplementary Fig. 29 The calculated total density of states (TDOS) and partial\ndensity of states (PDOS) of the Mn 3d and O 2p orbitals under different desodiation\nstates.\n“The manuscript presents a creative synthesis strategy with potential to advance\nLTMO cathode design. However, the current evidence is insufficient to support the\nproposed mechanisms. The current presentation does not meet the high standards of\nNature Communications for novelty, rigor, and impact. I recommend rejection without\nfurther review.”\nR: Thanks for the comments. We really appreciate the reviewer’s comments. We have\nfollowed the reviewer’s suggestion and addressed all proposed concerns by\ncomplementing the corresponding data and providing the explanations. The quality of\nthis work is significantly improved by following the suggestion from all reviewers.\nR2 3"}, {"round": 2, "reviewer_comments": {"reviewer_2": "The revised manuscript entitled “An ammonia-induced universal synthesis approach for manganese based layered oxides” has addressed all reviewer's comments thoroughly. The revised manuscript is well organized and the discussions are reasonable. I recommend to accept the manuscript immediately.", "reviewer_3": "All previous comments have been addressed, and the article is ready for publication"}, "author_response": ""}]

Decision letter