Tailoring superstructure units for improved oxygen redox activity in Li-rich layered oxide battery’s positive electrodes

Domain: earth_science
Published: 2024-11-18
Editorial outcome: major_revision

Abstract
The high-voltage oxygen redox activity of Li-rich layered oxides enables additional capacity beyond conventional transition metal (TM) redox contributions and drives the development of positive electrode active materials in secondary Li-based batteries. However, Li-rich layered oxides often face voltage decay during battery operation. In particular, although Li-rich positive electrode active materials with a high nickel content demonstrate improved voltage stability, they suffer from poor discharge capacity. Here, via physicochemical and electrochemical measurements, we investigate the correlation between oxygen redox activity and superstructure units in Li-rich layered oxides, specifically the fractions of LiMn6 and Ni4+-stabilized LiNiMn5 within the TM layer. We prove that an excess of LiNiMn5 hinders the extraction/insertion of lithium ions during Li metal coin cell charging/discharging, resulting in incomplete oxygen redox activity at a cell potential of about 3.3 V. We also demonstrate that lithium content adjustment could be a beneficial approach to tailor the superstructure units. Indeed, we report an improved oxygen redox reversibility for an optimized Li-rich layered oxide with fewer LiNiMn5 units.

Introduction
In contrast to conventional layered positive electrode oxides, such as LiCoO2, relying solely on transition metal (TM) redox activity, Li-rich layered oxides have emerged as promising positive electrode materials due to their utilization of both TM and oxygen redox at high voltage, resulting in an improved discharge capacity1. However, the instability of oxygen redox triggers severe voltage decay, posing a substantial challenge for the implementation of commercial applications2,3. Over the past decades, intensive studies have been proposed to elucidate the origins of voltage decay, including the continuous reduction of TM valence states, irreversible TM ions migration, and phase transitions (from layered to spinel/rock-salt phases) during cycling4–7. Recently, a more profound understanding8 has emerged, highlighting the crucial role of the accumulation of mesoscale lattice strain and lattice displacement within distinct nanoscale domains (LiTMO2 and Li2MnO3) as the driving force behind voltage decay during battery operation. Consequently, conventional post-synthesis treatments such as surface coating have proven ineffective in addressing this issue.

Tailoring the local structure and optimizing chemical composition is emerging as a promising avenue to mitigate voltage decay in Li-rich layered oxides. For instance, a consensus has emerged from relevant studies9–11, highlighting the detrimental role of localized superstructure domains characterized by in-plane Li/Mn order of Li-rich oxides, while delocalized or dispersed domains characterized by in-plane Li/Mn disorder effectively improve the oxygen redox reversibility and voltage stability. In addition, Li-rich oxides with O2-type exhibit improved voltage stability compared to conventional O3-type oxides, attributed to the reversible migration of TM ions12,13. In particular, an O2-type Li-rich oxide with a capped-honeycomb structure demonstrates negligible voltage decay14. However, synthesizing these positive electrodes through an ion exchange method from P2-type sodium positive electrode precursors presents challenges for large-scale production15. Introducing concentration gradients in Li-rich layered oxides, with low manganese and high nickel content on the particle surface, has shown improved voltage retention16,17. Nevertheless, uncertainties in nickel/manganese ions migration during high-temperature calcination lead to deviations from the intended design, compromising materials synthesis reproducibility. In terms of optimizing chemical composition, high-nickel Li-rich layered oxides enhance voltage stability by promoting the formation of Ni3+ in the pristine material. This serves as a redox buffer, suppressing the redox activation of Mn4+/3+ at low potentials18,19. Recently, Li et al.20 demonstrated the capability of nickel over cobalt in slowing the kinetics of ligand-to-metal charge transfer, mitigating TM ions migration, and resulting in reduced oxygen release and lower voltage decay. Moreover, Li-rich layered oxides can be obtained through hydroxide/carbonate co-precipitation, which is currently the preferred synthesis method for large-scale industrial production21. Despite their potential, high-nickel Li-rich layered oxides, like Li1.2Ni0.4Mn0.4O2, have received less attention due to their unsatisfactory electrochemistry, including lower discharge capacity and slower lithium ions diffusion compared to low-nickel Li-rich layered oxides22–25. However, there is still no comprehensive explanation for the unsatisfactory electrochemistry of high-nickel Li-rich layered oxides. Thus, the key to advancing Li-rich layered oxides is achieving voltage stability in high-nickel systems without compromising discharge capacity.

In this work, we present a comprehensive investigation of cobalt-free Li-rich layered oxides with varying nickel contents (Li1.20NixMn0.8-xO2, x=0.28, 0.32, 0.36, 0.40). By employing advanced techniques including synchrotron X-ray diffraction (SXRD), X-ray pair distribution function (PDF) analysis, X-ray absorption spectroscopy (XAS), solid-state nuclear magnetic resonance (NMR) spectroscopy, and density functional theory (DFT) calculations, we identify the types of honeycomb superstructure units within the TM layers and elucidate the influence of the valence state of nickel ions on these units. Furthermore, by integrating electrochemical investigations with soft XAS and resonant inelastic X-ray scattering (RIXS) spectroscopy, we establish a direct correlation between oxygen redox behaviors and the ratio of superstructure units. This correlation provides insights into the underlying reasons for the poor electrochemistry observed in high-nickel Li-rich layered oxides. Finally, we propose a practical solution to achieve improved electrochemical performance with reversible oxygen redox by tailoring honeycomb superstructure units in high-nickel Li-rich layered oxides.

Results
Structural characterization and electrochemistry
The Li1.20NixMn0.8-xO2 materials with x values of 0.28, 0.32, 0.36, and 0.40, denoted as N28, N32, N36, and N40, respectively, were synthesized via a solid-state reaction involving TM hydroxide precursors and Li2CO3 (with a molar ratio of Li/TM=1.5) under ambient air conditions (see “Methods”). Scanning electron microscopy (SEM) images (Supplementary Fig. 1) show that all materials have a typical polycrystalline microstructure with spherical secondary particles, each approximately 6 μm in diameter. High-resolution synchrotron X-ray diffraction (SXRD) measurements were used to obtain the crystal structure information. As shown in Supplementary Fig. 2, the main Bragg reflections for N28, N32, N36, and N40 correspond to a layered rhombohedral (R-3m) phase without impurities. Additionally, the broad diffraction peaks at 2θ values of 2.7–3.1° indicate honeycomb Li/TM ordering (C2/m space group) within the TM layer8. Meanwhile, these superstructure diffraction peaks show similar intensities (see enlarged image on the right), indicating all materials have comparable amounts and faulting of honeycomb domains. To obtain detailed structural parameters, we conducted SXRD refinements (Fig. 1a). For simplicity, we used a single R-3m phase as the initial structural model. The detailed refinement results are shown in Supplementary Table 1. The value of Li/Ni disorder is 2.4(4), 2.3(4), 2.2(4), and 1.8(4)% for N28, N32, N36, and N40, respectively, indicating that the cation arrangements in these materials are similar. To capture local structural arrangement, X-ray pair distribution function (PDF) measurements were conducted. The PDF patterns of LiTMO2 and Li2MnO3 are calculated and shown in Supplementary Fig. 3. The first two peaks below 3.0 Å represent the octahedral TM-O and the nearest TM-TM atom pairs, respectively26. The intensity of the second PDF peak (~2.9 Å) differs significantly between the R-3m model (mostly TM-TM pairs) and the C2/m model (partial TM-TM and partial Li-TM), leading to a distinct difference in the intensity ratio of the first and second PDF peaks. This distinction arises from the unique Li/Mn ordering in Li2MnO3. Consequently, the intensity of the second PDF peak is significantly weakened in the C2/m due to the relatively poor X-ray scattering capability of the Li atoms. As shown in Fig. 1b, the PDF patterns of N28, N32, N36, and N40 show no noticeable differences in the first two PDF peaks, indicating nearly identical atomic arrangements across all materials. High-resolution transmission electron microscopy (HRTEM) measurements were performed to verify local atomic arrangements. The HRTEM images (Supplementary Fig. 4) of N28, N32, N36, and N40 show clear lattice fringes without obvious evidence of Li/Ni mixing, consistent with the SXRD refinements. Additionally, the observed average interplanar spacing of approximately 0.47 nm corresponds to the (003) or (001) plane of the layered rhombohedral (R-3m) or monoclinic (C2/m) structures.

We then evaluated the electrochemical performance of these materials using Li metal coin cells with non-aqueous liquid electrolyte solution at a rate of 20 mA g−1 within the voltage range of 2.0–4.7 V at 25 °C. Figure 1c displays the characteristic voltage profiles of the Li metal cells with Li-rich layered oxides at the positive electrode, featuring a sloping part below 4.45 V and a subsequent plateau above 4.45 V8. The former is associated with cationic redox activity in LiTMO2 domains, whereas the latter involves an anionic (oxygen) redox activity in Li2MnO3 domains. Despite comparable total charge capacities, a notable divergence emerges in discharge capacity with increasing nickel content, leading to a gradual decrease. For example, Li||N28 coin cell and Li||N40 coin cell enable discharge capacities of 235.6 and 168.8 mAh g−1, with initial columbic efficiencies (ICE) of 71.4% and 58.1%, respectively. During discharging, all coin cells display nearly overlapping curves above 3.7 V. The primary difference arises from the reduction plateau at approximately applied potential of 3.3 V, which gradually disappears, as shown in the differential capacity (dQ/dV) curves depicted in Fig. 1d (indicated by the green arrow). Typically, the reduction peaks around 4.4 and 3.7 V are associated with the reduction of oxygen and nickel at high and mid potentials, respectively. The peak below 3.7 V, corresponding to low potential, probably relates to the reduction of oxygen or manganese1. Distinguishing these peaks is challenging in low-nickel Li-rich oxides due to their overlap. However, in high-nickel Li-rich oxides, the low-potential peak nearly disappears as nickel content increases. Thus, high-nickel oxides serve as an ideal model for investigating the redox mechanism in this region.

Identification of incomplete oxygen redox activity
To verify the redox nature of the dQ/dV peak observed in the region below 3.7 V, we conducted soft X-ray absorption spectroscopy (SXAS) measurements, using N36 as an example. Ni and O were analyzed in the fluorescence-yield mode (FY), while Mn was examined in the inverse partial-fluorescence-yield mode (iFY) at different charge/discharge states (as marked in Fig. 2a). FY and iFY measurements probe to a depth of approximately 100 nm, thus providing information about the bulk of the material. Charging to 4.45 V led to Ni oxidation, evidenced by a decrease in peak A intensity (~852.9 eV) and an increase in peak B (~854.8 eV) in the Ni L3-edge spectra (Fig. 2b)27. At the end of the 4.7 V charge, the spectra exhibited slightly inverted changes, possibly related to charge transfer from ligand to metal20. Upon discharging to 2.0 V, reversed changes in the Ni L3-edge spectra indicated Ni reduction, with peak A showing higher intensity and peak B lower intensity compared to the pristine state, suggesting a reduction of nickel species to a lower valence state. The Mn L-edge spectra remained constant, indicating Mn stability in an inactive Mn4+ valence state throughout the entire process (Fig. 2c).

In the O K-edge spectra (Fig. 2d), the pre-edge region (below 535 eV) primarily reflects the electron transition from O 1s level to the unoccupied TM 3d-O 2p hybridized states28. Charging to 4.45 V induces increased intensity at around 528.5 eV (marked as a gray arrow), attributed to Ni oxidation, consistent with prior analysis. Additionally, a new peak emerges at approximately 530.8 eV (dashed lines), becoming more pronounced upon charging to 4.7 V. This peak is recognized as characteristic of oxidized oxygen species, as observed in the Li-rich layered oxides29,30. Notably, this peak associated with oxidized oxygen retains some intensity after discharging of 2.0 V, failing to fully revert to its pristine state, suggesting incomplete oxygen redox reaction within the material. To further elucidate the nature of oxidized oxygen, we conducted O K-edge resonant inelastic X-ray scattering (RIXS) measurements. In Fig. 2e, at 4.7 V in the charged state, N36 exhibits a distinct characteristic peak C (around 7.8 eV, denoted by the dashed line) and vibrational peak D around the elastic region (0 eV). In addition, the enlarged view of the vibrational peak is presented in Fig. 2f, revealing vibrational frequencies of 1498 cm−1, similar to the O-O bond length of molecular O2, aligning with recent research findings31,32. More importantly, after discharging to 2.0 V, the distinctive feature of molecular O2 persists (indicated by the gray arrow), suggesting an incomplete reduction of molecular O2 to O2- species. These findings are consistent with the SXAS results. Thus, we identify that the incomplete oxygen redox occurs at approximately applied potential of 3.3 V in Li-rich layered oxides for a high-nickel system.

Superstructure characterization and theoretical calculation
Typically, molecular O2, in the charged state, is trapped in vacancy clusters within the TM layer of Li-rich layered oxides31. Our experiments have demonstrated the presence of molecular O2 at the end of discharge in high-nickel Li-rich layered oxides. This suggests that lithium ions cannot smoothly re-enter the TM layer to coordinate with oxygen during discharge, leaving some oxygen species as molecular O2. Therefore, the issue may lie in the TM layer, specifically at the local honeycomb superstructure. To detect the honeycomb superstructure distribution, we conducted 6Li solid-state nuclear magnetic resonance (NMR) spectroscopy measurements. Figure 3a and Supplementary Table 2 present the 6Li NMR patterns with fitting results. The sharp peak at 0 ppm represents diamagnetic species such as LiOH, Li2CO3, and organic lithium salts33. Two prominent groups of peaks are observed at 500–1000 ppm and 1300–1500 ppm, corresponding to Li in the Li layer (LiLi) and Li in the TM layer (LiTM), respectively34. The NMR patterns reveal two distinct peaks at approximately 1300 and 1500 ppm, indicating two different Li environments within the TM layer. Consistent with previous literature35, these peaks are assigned to honeycomb superstructure units for LiMn6 (purple, Li surrounded by six Mn) and LiNiMn5 (blue, Li surrounded by one Ni and five Mn), respectively. The corresponding atomic structures are plotted in Fig. 3b. Interestingly, although all materials show similar total lithium content in the TM layer, the local lithium environments differ. Specifically, NMR patterns reveal a gradual increase in the intensity of peaks corresponding to LiNiMn5 units, while the intensity of peaks for LiMn6 units decreases from N28 to N40. Fitting results show that the ratio of LiNiMn5/(LiNiMn5 + LiMn6) increases progressively: 0.183 for N28, 0.373 for N32, 0.540 for N36, and 0.613 for N40 (Fig. 3c). Thus, the varying Ni/Mn ratios in Li-rich layered oxides affect the type of superstructure units within the TM layer. As nickel content increases, the Li-rich oxide exhibits more LiNiMn5 units, replacing LiMn6 units. Since it is difficult to detect LiNiMn5 superstructure units in conventional Li-rich layered oxides (e.g. Li1.2Ni0.2Mn0.6O2), the potential impact of this unit on electrochemistry may have been overlooked. To understand the formation mechanism of LiNiMn5, we conducted density functional theory (DFT) calculations. We chose Li44Mn18Ni10O72 (Li1.22Mn0.50Ni0.28O2, LiNiMn5 within the TM layer) as the structural model (Supplementary Fig. 5a). By calculating the spin density of all atoms, we confirmed their valence states. The calculations show Mn with +4 valence state, Ni with +2, +3, and +4 valence states, and O with -2 valence state. Importantly, Ni occurs exclusively as Ni4+ within the highly symmetric LiNiMn5 units, whereas Ni2+ or Ni3+ are identified in more Ni-rich local configurations. Note that we have calculated the oxidation state of Ni in the LiNiMn5 configuration without a priori assumptions on the Ni oxidation state. The relaxed structure shows that Ni possesses the oxidation state +4. In addition, we have calculated models with the same overall composition but without the LiNiMn5 configuration (always without assuming the oxidation state a priori), in which we find that Ni has an oxidation state lower than +4. These findings are supported by the density of states (DOS) analysis (Supplementary Fig. 5b). The detailed calculation processes are described in the “Methods” section. To confirm the Ni valence state changes in these materials, we conducted XAS measurements. As shown in Fig. 3d, the Ni K-edge spectra shift to higher energy from N28 to N40, indicating an increase in the average Ni oxidation state. In contrast, the Mn K-edge spectra position remains constant for all materials, indicating an Mn4+ oxidation state (Fig. 3e). These valence state changes are further supported by the Fourier transform extended X-ray absorption fine structures (FT-EXAFS) results, shown in Supplementary Fig. 6. While Ni4+ is atypical in conventional layered oxides, its presence in Li-rich layered oxide systems is plausible due to the lithium excess structural configuration36. Consequently, forming LiNiMn5 superstructure units is easier in N40 than in N28, as the higher Ni valence state in N40 allows for more Ni4+ to be present.

In addition, Pauling’s electrostatic valence principle37 and semi-quantitative equation38 can be used to evaluate the possibility of a coordination configuration. According to principle of local electrical neutrality, a structure may become unstable if 4Z>0. The most common LiMn6 superstructure units, equivalent to Li3-O-LiMn6, exhibit a stable state with 4Z=0. Similarly, LiNiMn5 units (Li3-O-LiNiMn5) also demonstrate 4Z=0, attributed to the presence of Ni4+ rather than Ni2+ (4Z=1=3) and Ni3+ (4Z=1=6). This aspect elucidates why prior investigations failed to detect nickel within Li2MnO3-like domains, given that the possibility of the presence of Ni4+ in LiNiMn5 superstructure units has been previously overlooked38,39. Previous theoretical calculations39 indicate that the Li4MnNi-coordinated O, referred to as LiNiMn5 in our work, exhibits lower cationic and anionic redox activity compared to LiMn6 units, indicating limited electrochemical activity. Moreover, lithium ions move from one octahedral site to another, passing through an intermediate tetrahedral site where they encounter strong repulsion from nearby TM ions. The associated activation barriers for lithium movement with Ni4+, Ni3+ and Ni2+ are 490, 310, and 210 meV, respectively40. These theoretical findings imply that LiNiMn5 may impede the extraction/insertion of lithium ions. During charging/discharging, Li-rich oxides experience a cationic–anionic redox inversion, leading to electrochemical asymmetry27. As a consequence, at low discharge potentials, some lithium ions fail to return to the TM layer to coordinate with oxygen, resulting in incomplete oxygen redox. This analysis is consistent with the findings from O K-edge SXAS and RIXS experiments.

Combining all the analysis described above, a clear correlation between the incomplete oxygen redox reaction and honeycomb superstructure units can be established. In Li-rich layered oxides, as the nickel valence state increases (reflecting increased nickel content from N28 to N40), positive electrodes exhibit a rise in LiNiMn5 at the expense of LiMn6 units. Positive electrodes with higher LiNiMn5 content exhibit limited electrochemical activity, resulting in incomplete oxygen redox reactions (retention of molecular O2) at a low potential of 3.3 V, thereby leading to decreased discharge capacity. In contrast, Li-rich oxides with less LiNiMn5, such as N28, demonstrate enhanced oxygen redox reversibility, showing higher discharge capacity.

Reversible oxygen redox activity through tailoring superstructure units
To address these intrinsic properties in high-nickel Li-rich layered oxides, we propose a practical solution: reducing the lithium content to lower the Ni valence state and thus tailor the LiNiMn5 superstructural units. During the synthesis, we varied the lithium content and the ratio of TM precursors. Specifically, we used Li/TM ratios of 1.2, 1.3, 1.4, and 1.5 based on the N36 precursor Ni0.45Mn0.55(OH)2. The resulting chemical formulas are Li1.09Ni0.41Mn0.50O2 (LM12), Li1.13Ni0.39Mn0.48O2 (LM13), Li1.17Ni0.37Mn0.46O2 (LM14), and Li1.20Ni0.36Mn0.44O2 (LM15). It is important to note that LM15 is equivalent to the N36 mentioned earlier. Inductively coupled plasma-optical emission spectroscopy (ICP‒OES) results confirm that the Li/TM molar ratios are consistent with the design values, except for some lithium evaporation at high temperatures (Supplementary Table 3). As shown in Supplementary Fig. 7, LM12, LM13, LM14, and LM15 maintain a layered phase with space group R-3m without impurities. In the enlarged view on the right, the diffraction intensity related to the superstructure decreases from LM15 to LM12, indicating a reduction in superstructure content as lithium content decreases. The results of the SXRD refinement and corresponding data are presented in Supplementary Fig. 8 and Supplementary Table 4. The Li/Ni mixing reveals an increasing level of disorder from LM15 to LM12, with LM15 displaying a minimal disorder value of 1.8(4)%, while LM12 shows a significantly higher disorder value of 6.3(3)%. Considering both discharge capacity and cycle stability, LM13 is selected as the optimized positive electrode for further discussion (Supplementary Fig. 9).

The PDF patterns of both cathodes are shown in Supplementary Fig. 10. LM13 shows a higher intensity of the second PDF peak correspond to TM-TM pairs compared to LM15 (indicated by the arrow). It suggests LM13 exhibits more R-3m than C2/m domains, which agrees with SXRD analysis. Additionally, both materials exhibit a uniform distribution of Ni, Mn, and O, confirmed by transmission electron microscopy (TEM) energy dispersive X-ray spectroscopy (EDS) mapping images (Supplementary Fig. 11). To investigate the distribution of honeycomb superstructure units, we conducted 6Li NMR measurements of LM13. Compared to LM15, LM13 exhibits a lower NMR peak intensity at approximately 1300 ppm, indicating fewer LiNiMn5 units (Fig. 4a). The NMR fitting results (Supplementary Table 5) show that the ratio of LiNiMn5/(LiNiMn5 + LiMn6) is 0.427 for LM13, lower than 0.540 for LM15 (Fig. 3c). Meanwhile, XAS results show a lower Ni oxidation state in LM13, while Mn maintains the same oxidation state (+4) for both materials (Fig. 4b and Supplementary Fig. 12). The XAS analysis is further supported by the corresponding FT-EXAFS results, shown in Supplementary Fig. 13. Overall, through our comprehensive analyzes (including SXRD, PDF, XAS, and NMR), we have successfully tuned the composition and content of the honeycomb superstructure units simultaneously. In Fig. 4c, Li||LM13 coin cell exhibits a longer sloping part (below 4.45 V) and a shorter oxygen redox plateau than Li||LM15 coin cell, reflecting differences in superstructure contents. Despite a slight trade-off in charge capacity due to lower lithium content, Li metal coin cell with the LM13-based positive electrode demonstrates a notable capacity of 231.1 mAh g−1 with an ICE of 78.4%. This surpasses the 199.0 mAh g−1 of Li||LM15 coin cell (ICE of 69.8%). The dQ/dV curves in Fig. 4d highlight the significant difference in discharge capacity, attributed to the co-contribution of nickel redox and oxygen redox (especially at a low potential of 3.3 V, indicated by the green arrow). This modulation of cationic and anionic redox aligns precisely with the impact of superstructure content and units in Li-rich layered oxides. To confirm the reversibility of oxygen redox in the LM13, we conducted O K-edge RIXS measurements. Figure 4e shows a clear molecular O2 signature during charging, indicating oxygen oxidation. In contrast to LM15 (Fig. 2e), no signal is detected at the end of discharge for LM13. The enlarged view of the elastic region (0–1 eV) in Fig. 4f further supports these observations, indicating complete oxygen redox after tailoring the superstructure units. Note that the relatively lower oxygen redox signal in LM13 is attributed to the reduced superstructure content, consistent with above structural analysis.

Structuralevolutionanalysis
To track the evolution of lithium local environments during charge/discharge, ex-situ 7Li NMR spectroscopy was performed. In Fig. 5a-b, the 7Li NMR results in the pristine state, in agreement with the 6Li NMR measurements described above, confirm fewer LiNiMn5 units in LM13. In the charged state at 4.35 V for LM15, the intensity of the peak corresponding to LiNiMn5 remains unchanged (see green dashed line), indicating minimal changes in the oxidation state of TM ions within LiNiMn5, confirming the presence of Ni4+-stabilized LiNiMn5 units rather than Ni2+ or Ni3+. Beyond 4.5 V, the NMR peak broadens due to structural disorder induced by TM ions migration in Li environments, making it challenging to identify specific superstructure units in both positive electrodes. In the charged state at 4.7 V, some NMR signals persist in the 1300–1500 ppm region in LM15, indicating lithium ions in the TM layers. In sharp contrast, LM13 shows no signal in this region, implying complete lithium-ion removal from the TM layer. Moreover, the lower LiLi (500–1000 ppm) intensity suggests fewer lithium ions in the lattice for LM13. During discharge, the peak intensity increases as lithium ions are inserted, with LM13 exhibiting improved capability for extraction/insertion of lithium ions, as indicated by the higher total peak intensity in both LiLi and LiTM.

To capture the structural evolution during the first cycle, in-situ XRD measurements were conducted, and lattice parameters along with volume changes were obtained through Rietveld refinement (Fig. 5c-d). Below 4.45 V (indicated by the dashed line), both positive electrodes show an increase in lattice parameter c, with the 003 Bragg reflection shifting to lower angles. This shift is attributed to increased electrostatic repulsion accompanying the removal of lithium ions from the Li layer8. LM15 exhibits minimal changes in lattice parameter c at the oxygen redox plateau, whereas LM13 undergoes significant lattice shrinkage, indicating lithium ions being removed from the TM layer, consistent with previous reports on Li-rich layered oxides41. The minor changes in LM15, supported by NMR analysis, suggest some lithium ions still reside in the lattice, offsetting the lithium removal effect at the end of the charge. During discharge, both positive electrodes show similar changes, initially increasing and then decreasing in lattice c, implying the insertion of lithium ions into the TM and Li layers, respectively. The lattice parameter c, however, cannot revert to the initial position, showing a difference of 0.065 Å. In stark contrast, LM13 demonstrates minor changes in parameter c, with a change of only 0.015 Å after one complete cycle. The lattice parameter a, sensitive to the TM-O bond length in the TMO2 slab, serves as indirect evidence of cationic redox20. In both positive electrodes, a linear decrease in the slope region and minor changes at the oxygen redox plateau are observed, corresponding to nickel oxidation and oxygen oxidation, respectively. During discharge in LM13, lithium ions are inserted into the TM and Li layers, corresponding to oxygen, nickel, and oxygen reduction at high/middle/low potential, respectively (indicated by the colored arrows in Fig. 5d). This redox mechanism evolution aligns with other reports on Li-rich layered oxides1,42. However, in LM15, the absence of oxygen reduction at the low potential of 3.3 V, confirmed by O K-edge XAS and RIXS results, means that changes in the lattice parameter a correspond to oxygen reduction followed by nickel reduction during discharge. Additionally, both positive electrodes experience volume shrinkage after the first cycle, with LM13 exhibiting a smaller volume change of 0.382 Å3 compared to LM15 (0.675 Å3). This, combined with the previous evolution of the lattice parameter c, confirms the superior structural stability of LM13.

Electrochemistry, chemical and structural stability analysis
Benefiting from the tailored superstructure content and units, as shown in Fig. 6a, Li||LM13 coin cell displays improved capacity retention at 84.8% (187.7 mAh g−1) after 200 cycles at current of 67 mA g−1. This contrasts sharply with Li||LM15 coin cell, which exhibits a capacity retention of 71.2% (134.9 mAh g−1) based on the maximum capacity. When evaluating Li-rich layered oxides, it is more reasonable to assess the average cell discharge voltage due to the significant challenge of voltage decay. After 200 cycles, Li||LM13 and Li||LM15 coin cells show average cell discharge voltages of 3.47 and 3.28 V, corresponding to a decrease of 1.36 and 2.43 mV per cycle (Fig. 6b). The selected charge/discharge curves show the distinction in discharge voltage evolution (Supplementary Fig. 14). Furthermore, LM13 demonstrates improved electrochemical energy storage performances at a specific current of 20 mA g−1 (Supplementary Fig. 15). Additionally, Fig. 6c highlights the good rate capability of Li||LM13 coin cell, delivering a discharge capacity of 140.9 mAh g−1 even at the high rate of 1000 mA g−1.

To verify the oxygen redox stability during long-term cycling, we conducted O K-edge RIXS measurements after 100 cycles at 20 mA g−1 (approximately 92 days) in the charged state at 4.7 V. As shown in Fig. 6d-e, characteristic features of molecular O2 are clearly observed around an excitation energy of 531 eV and an emission energy of 523.7 eV in LM13 (indicated by the white arrow), while they are barely visible in LM15. Additionally, the O K-edge RIXS spectra collected at an excitation energy of 531 eV (Supplementary Fig. 16) show that LM13 still exhibits a distinct characteristic peak A (around 7.8 eV) and a vibrational peak B around the elastic region (0 eV). In sharp contrast, no such signals are detected in LM15. These results indicate that LM13 has higher oxygen redox stability and reversibility, which explains its improved electrochemical performance. To demonstrate the structural stability of LM13 and LM15, we conducted electron energy loss spectroscopy (EELS) measurements. In the pristine state, as shown in Supplementary Fig. 17, the Mn and Ni L-edge spectra of both electrodes remain at the same energy loss position from the surface to the bulk. However, after 200 cycles, the Mn L-edge spectra of LM15 exhibit a continuous shift to lower energy within a depth of 20 nm (indicated by arrows in Fig. 6g). This shift suggests a gradual reduction of Mn at the surface43, which is further supported by the fading of the pre-edge peak in the O K-edge spectra (highlighted by the rectangle in Fig. 6h). In contrast, Fig. 6j-k shows that the spectral shifts in LM13 mainly occur up to a depth of 10 nm from the surface. The reproducibility of these results was confirmed by acquiring EELS spectra of the same particle along different directions (Supplementary Fig. 18). Additionally, as shown in Fig. 6f, the TEM images LM15 after 200 cycles clearly reveal extensive voiding in LM15, characterized by a high density of voids distributed throughout the particle. In contrast, LM13 exhibits minimal voids, confirmed by the bright particle images (Fig. 6i). These findings indicate that LM13 demonstrates robust surface stability. To validate the versatility of our strategy for modulating the honeycomb superstructure in high-nickel Li-rich oxides, we applied it to Li1.2Ni0.4Mn0.4O2, maintaining the same Li/TM ratio of 1.3 during synthesis (Supplementary Fig. 19). Similarly, improved electrochemical performance is also observed for the modified oxide (Li/TM=1.3). Overall, our results confirm that adjusting lithium content is an effective and straightforward approach to enhancing the discharge capacity of high-nickel Li-rich layered oxides.

Methods
Materials Synthesis
Li1.2NixMn0.8-xO2 (x=0.28, 0.32, 0.36, 0.40) materials were synthesized using a hydroxide co-precipitation method followed by a solid-state reaction. First, a 2.0 M aqueous solution containing NiSO4·6H2O (≥98%, Sigma‒Aldrich) and MnSO4·H2O (≥98%, Sigma‒Aldrich) (molar ratios of 7:13, 2:3, 9:11 and 1:1 for N28, N32, N36 and N40, respectively), along with a 4.0 M NaOH solution (≥98%, Sigma‒Aldrich) and an aqueous 0.5 M NH3·H2O solution (28%–30%, Sigma‒Aldrich), were separately fed into a batch reactor with a volume of 1 L. The aqueous TM solution and NH3·H2O solution were fed at the same rate while maintaining the pH value at 11.0±0.2 by adjusting the feeding rate of the NaOH solution. The reaction process was carried out at a temperature of 50 °C and under a N2 atmosphere. After a reaction duration of 20 h, followed by a 2-h aging process, the precursor was washed three times with deionized water and subsequently dried in an oven at 80 °C for 12 h under an air atmosphere. Finally, the target positive electrodes were obtained by mixing the precursor and Li2CO3 (≥99%, Sigma‒Aldrich) in a molar ratio of 1:1.5, followed by a heat treatment at 500 °C for 5 h and subsequently at 850 °C for 12 h under an air atmosphere. Similarly, other positive electrodes were prepared utilizing identical procedures except for a change in the ratio of Li2CO3 and precursor (1.2:1, 1.3:1, 1.4:1, and 1.5:1 for LM12, LM13, LM14, and LM15, respectively).

Electrochemical measurements
The electrodes were prepared by thoroughly mixing the active material, Super C65 conductive carbon black (MTI Co., Ltd.), and polyvinylidene fluoride (PVDF, Sigma-Aldrich) in a mass ratio of 8:1:1 using a mortar. N-methyl-2-pyrrolidone (NMP) with a moisture content of less than 0.1% (VMR) was used as the solvent. The resulting mixture was then processed in a planetary mixer (THINKY ARV-310) with 2000 rpm for 10 min under an air atmosphere to ensure homogeneity. The resulting slurry was applied to aluminum foil (15 μm thickness, häberle LABORTECHNIK GmbH & Co. KG) at a speed of 25.0 mm s−1 using a ZUA 2000 Universal applicator and the thickness was 150 nm. The coated aluminum foil was subsequently dried in an oven at 80 °C under ambient conditions to evaporate the NMP solvent. After a drying period of 6 h, the electrode material was punched into disks with a diameter of 12 mm (with an active material loading of approximately 3 mg cm−2) using a handheld punch (NOGAMI GIKEN Co., Ltd.). A subsequent drying step was conducted using a Büchi glass oven (B-585) under vacuum conditions at 120 °C for 12 h. To evaluate the electrochemical performance, 2025-type coin cells were assembled. These cells comprised a Li metal anode (14 mm diameter, 0.25 mm thickness, PI-KEM), a single-layer Celgard 2500 membrane (25 μm thickness, 55% porosity, DODOCo., Ltd.) as the separator, and 80 μL of electrolyte consisting of 1 M lithium hexafluorophosphate (LiPF6) dissolved in a solvent mixture of ethylene carbonate (EC), ethyl methyl carbonate (EMC), and dimethyl carbonate (DMC) in a volume ratio of 1:1:1, with moisture content maintained below 10 ppm (DODO Co., Ltd.). All assembly procedures were conducted within an argon-filled glovebox (O2 and H2O < 0.1 ppm). The cells were cycled at a constant temperature of 25 °C using a Biologic VMP3 multichannel battery test system within a voltage range of 2.0–4.7 V. The charging protocol included a constant voltage step at 4.7 V for 10 min, followed by a 10-min resting period after each charge-discharge cycle. To ensure reproducibility, all electrochemical experiments were performed using a minimum of two-coin cells. Additionally, the current density for each cell was calculated based on the mass of the active material in the electrode. The coulombic efficiency was determined as the percentage ratio of discharging capacity to charging capacity, multiplied by 100. The average cell discharge voltage was defined as the voltage when the discharge capacity reached half of its maximum. For all ex-situ electrode measurements conducted in this work, which included SXAS, NMR, and RIXS, the coin cells were subjected to specific voltage applications before being disassembled in the glovebox. The obtained electrodes were then thoroughly washed with DMC solvent three times and sealed under vacuum conditions within the glovebox.

Materials characterizations
X-ray diffraction (XRD) measurements were measured by a Mo source powder diffractometer (STOE STADI P) in transmission geometry, with a wavelength of 0.7093 Å. The samples were loaded into capillaries with an outer diameter of 0.5 mm. The acquisition time for each pattern was 150 min. Synchrotron X-ray diffraction (SXRD) measurements were conducted at beamline P02.1, PETRA III (wavelength of ~0.207 Å) at DESY in Hamburg44,45. The samples were loaded into capillaries with an outer diameter of 0.5 mm. Measured intensities were collected using a VAREX CT4343 detector (2880×2880 pixels, 150×150 μm2 each) and the acquisition time of 120 s for each pattern. NIST SRM 660c (LaB6) was used for geometry calibration performed with the software DAWN46 followed by image integration including geometry, solid-angle, and polarization corrections. In-situ XRD measurements were conducted using an Ag source powder diffractometer (STOE STADI P, wavelength of 0.5594 Å) in transmission geometry. In-situ 2032-type coin cells with a Kapton window were employed. All cells were cycled at 20 mA g−1 in a 2.0–4.7 V range for one cycle at 25 °C using a battery system (NOVA). The collection time for each pattern was set at 20 min. Rietveld refinements were carried out using the Fullprof software package47.

Pair distribution function (PDF) measurements were performed at beamline ID31 at the European Synchrotron Radiation Facility (ESRF). The sample powders were loaded into cylindrical slots (approx. 1 mm thickness) held between Kapton windows in a high-throughput sample holder. Each sample was measured in a transmission with an incident X-ray energy of 75.00 keV (wavelength of 0.1653 Å). Measured intensities were collected using a Pilatus CdTe 2M detector (1679×1475 pixels, 172×172 μm2 each) positioned with the incident beam in the corner of the detector. The sample-to-detector distance was approximately 0.3 m for the total scattering measurement. Background measurements for the empty windows were measured and subtracted. NIST SRM 660b (LaB6) was used for geometry calibration performed with the software pyFAI followed by image integration including a flat-field, geometry, solid-angle, and polarization corrections. The calculated PDF data with R-3m and C2/m was performed by PDFgui package48.

The morphology of all samples was conducted by a scanning electron microscopy (SEM, Zeiss Merlin) with an acceleration voltage of 10 kV. The compositions of the samples were quantitatively determined by inductively coupled plasma-optical emission spectroscopy (ICP‒OES) using a Thermo Fischer Scientific iCAP 7600 DUO. Scanning Transmission electron microscopy (STEM) energy dispersive X-ray spectroscopy (EDS) mapping and electron energy loss spectroscopy (EELS) experiments were conducted using a double aberration corrected Thermo-Fisher Themis-Z (operated at 300 kV equipped with a Super-X EDS detector and a Gaten GIF Continuum 970 High Res EELS spectrometer). The structural information of samples was studied by high-resolution transmission electron microscopy (HRTEM). Each specimen was examined using Thermofisher Talos F200X TEM operated at 200 kV.

Spectroscopy characterization
Magic-angle spinning (MAS) nuclear magnetic resonance (NMR) spectroscopy was performed on a Bruker Avance neo 200 MHz spectrometer with a magnetic field strength of 4.7 T. MAS spinning was carried out using 1.3 mm rotors at a frequency of 55 kHz. For the 6Li NMR experiments, 6LiOH·H2O (95 atom%, Sigma‒Aldrich) was utilized as the Li source for synthesis. The Larmor frequencies for 6Li and 7Li NMR were 29.5 MHz and 77.8 MHz, respectively. Spectra were acquired employing a rotor-synchronized Hahn-echo pulse sequence, with a 90° pulse length of 0.85 μs for 7Li and 1.6 μs for 6Li. All spectral shifts were referenced to an aqueous solution of LiCl (6LiCl for 6Li) at 0 ppm. The spectral intensities were normalized based on the sample mass and the number of scans. The NMR spectra were fitted using DMFIT program49. The atomic structures were plotted using VESTA software50.

Hard X-ray absorption spectroscopy (XAS) experiments were performed at the XAS beamline of the KIT synchrotron in Karlsruhe. These measurements were conducted at 25 °C, employing the transmission mode for data acquisition. The obtained XAS data were processed using ATHENA software51. For soft X-ray absorption spectroscopy (SXAS), experiments were carried out at the WERA beamline at the KARA synchrotron in Karlsruhe. The Ni L-edge and O K-edge spectra were recorded using the fluorescence yield (FY) detection mode and the inverse partial fluorescence yield (iFY) for the Mn L-edge spectra. Resonant inelastic X-ray scattering (RIXS) experiments were conducted at beamline U41-PEAXIS at BESSY II, located at Helmholtz-Zentrum Berlin (HZB)52. A vacuum suitcase was used to transfer the sample from an N2-filled glovebox to the test chamber. The spectrometer was positioned at specular conditions relative to 60° scattering angle and was optimized to a combined resolution of 90 meV using a carbon tape. The O K-edge RIXS spectra of the samples were collected at an excitation energy of 531.0 eV. The acquisition time for each pattern was 30 min. RIXS mapping was measured in 0.5 eV energy steps from 526 eV to 534 eV. Data were processed using the Adler-4.0 software package.

DFT calculations
The pre-optimization of the structure model was done by using the Vienna ab-initio simulation package (VASP) within the projector augmented-wave approach using the Perdew-Burke-Ernzerhof (PBE) with r2SCAN functional53–57. The correlation effects of the 3d-Ni and Mn orbitals (corresponding parameter UMn = 1.8 eV, UNi = 2.1 eV) and long-range dispersion with Grimme’s D4 correction were taken into account58,59. A plane-wave energy cut-off of 450 eV and 2×2×2 gamma centered k-point grid were used for the total energy.

To obtain reasonable electronic properties, the obtained structure model was reoptimized with CRYSTAL17 program version 1.0.2 using the PW1PW hybrid functional which has been successfully used, e.g., for the determination of transition metal oxide nitride phase stability and orthophosphates before60–62. Long-range London dispersion was taken into account with Grimme’s D3 correction with Becke-Johnson damping63–65. The empirical s8 parameter was re-adjusted in preliminary calculations of the relative stability of α-VOPO4 and α-NbOPO4, in which a modified value of s8 = 1.5363 was obtained66. Several convergence parameters were changed to increase the numerical precision of the calculations. The truncation criteria for bielectronic integrals (TOLINTEG) were set to 7 7 7 7 14. To accelerate the SCF convergence, FMIXING was increased from default 30% to 85%. The Monkhorst-Pack shrinking factor is 4×4. The pob-TZVP-rev2 basis sets were used in all calculations67,68. Since Ni and Mn have magnetic moments, some pre-settings needed to be done. The calculation of the spins was done in an antiferromagnetic configuration with a maximum total spin of 66. Additionally, SPINLOCK was set to 6630. The densities of states are calculated according to the Fourier-Legendre technique69. For a reliable result, the number of k-points must be set to 44.

The development of different structure models for Li44Mn18Ni10O72 (Li1.22Mn0.50Ni0.28O2) was done by the program supercell70. The initial structure model is based on LiMnO2 (S.G. C2/m, a=4.9370(10) Å, b=8.5320(10) Å, c=5.030(2) Å, β=109.46(3)°). For the calculations the primitive cell was used (a=4.928 Å, b=4.928 Å, c=5.030 Å, α=80.35°, β=99.60°, γ=60.11°) with adjusted occupancies for the metal cations Li, Mn and Ni (Supplementary Table 6). For the given supercell input file with the partial S.O.F in the transition metal positions with a 3×2×2 cell ca. 3:1 ≈ 108 different combinations for Li44Mn18Ni10O72 were found. The structure model for Li44Mn18Ni10O72 was chosen based on low coulombic energy and the honeycomb superstructure LiNiMn5. Supplementary Fig. 5a shows the exemplary crystal structure of the model. Atomic coordinates are given in Supplementary Table 7. The atomic structures were plotted using VESTA software.

Peer review rounds
[{"round": 1, "reviewer_comments": {"reviewer_1": "The authors revealed a direct correlation between oxygen redox activity and honeycomb superstructure units in Li-rich layered oxides, specifically the fractions of LiMn6 and Ni4+ stabilized LiNiMn5 within the TM layer and an excess of LiNiMn5 impedes the extraction/insertion of lithium ions during charging/discharging, resulting in incomplete oxygen redox activity. It is generally agreed upon that the increase in Ni content will reduce the specific capacity and the first cycle efficiency and result in an increase in Ni oxidation state from the reports from Qi et al (Electrochem 2023, 4, 21–30) and Chong et al (RSC Adv. 2016, 6, 53662). Here the authors show the similar phenomenon but with a reasonable explanation or mechanism for these conclusions, therefore I am unsure of the overall novelty - this should be addressed.\nOverall, this manuscript could be published after a minor revision.\nFor Figure 2d and 2f, the signals are very weak, and the noise is higher, how much confidence you have for the data? Figure 2c, 2e and 2f, it seems that the authors only showed three points, where is the fourth point data/curve?\nTo address these intrinsic properties in high-nickel Li-rich layered oxides, the authors proposed to reduce the Li content during synthesis, please explain the rational design, why Li content determine the amount of the LINiMn5? Also, why not just to optimize the Ni/Mn ratios instead?", "reviewer_2": "Li-rich layered oxides have high capacity because of high-voltage oxygen redox activity. Li-rich cathodes with a high nickel content exhibit improved voltage stability but suffer from poor discharge capacity. This work tailors the LiNiMn5 superstructure content by modification of Ni and Li contents. The study finds that an excess of LiNiMn5 impedes the extraction/insertion of lithium ions during charging/discharging, resulting in incomplete oxygen redox activity (retention of molecular O2) at low potential (~3.3 V, versus Li+/Li). If the evidence is sufficient, the work is more interesting to open a new idea for enhanced oxygen redox of Li-rich layered oxides. There are several puzzling problems, so it is not recommended for publication:\n1, In Li1.20NixMn0.8-xO2, x = 0.28, 0.32, 0.36, 0.40 materials, the content of Ni with +4 valence increases with the increase of x or the decrease of Li content. The work finds that the LiNiMn5 superstructure content increases with the increase of x or Li content. This is contradictory. The LiNiMn5 superstructure content is determined by the Li content or the ratio of Li/TM? Please give a reasonable interpretation.\n2. In conventional LLOs such as Li1.2Ni0.2Mn0.6O2, generally, based on Pauling´s electrostatic valence rule, LiMn6 configurations are preferably aggregated in Li2MnO3 due to their special stability. So it is difficult to permeate Ni to Li2MnO3. Many researches found that the breakdown of aggregated Li2MnO3, such as the permeation of TM with strong TM-O bonds, benefited the enhanced oxygen redox. Though the work gives evidence of LiNiMn5 configurations, why can the LiNiMn5 configurations exist in the work? Why is a reason that LiNiMn5 configurations result in incomplete oxygen redox activity?\n3. The work does not provide TEM images, thus the other existence of Ni is unclear, and for example, the occupation of less Ni at Li-layer sites reduced oxygen release and lowered voltage decay.\n4. The DOS of Li44Mn18Ni10O72 (Li1.22Mn0.50Ni0.28O2) is employed using DFT calculations. The state of Ni4 is +4, as expected in the work’s model. Other models without the LiNiMn5 configurations are stable?\n5. In addition, 231.1 mAh g-1 at C/10 and a capacity retention of 84.8% after 200 cycles at C/3 is not high, such as 276.5 mAh g−1 at C/10 and 85.4% after 300 cycles at 1 C of Li1.2Ni0.2Mn0.6O2(Li J, ACS nano, 2023, 17, 16827-16839)."}, "author_response": "Point-to-point Response to Reviewers’ Comments\nTailoring Superstructure Units for Enhanced Oxygen Redox in Li-Rich Layered\nOxides\n(Manuscript ID NCOMMS-24-16929)\nWe sincerely thank the editor and all reviewers for their valuable review comments that we\nhave used to improve the quality of our manuscript. The reviewer comments are laid out below\nand specific concern have been numbered, and all the responses are listed in the following\nsection, point by point.\nReviewer #1 (Remarks to the Author):\nThe authors revealed a direct correlation between oxygen redox activity and honeycomb\nsuperstructure units in Li-rich layered oxides, specifically the fractions of LiMn and Ni4+\n6\nstabilized LiNiMn within the TM layer and an excess of LiNiMn impedes the\n5 5\nextraction/insertion of lithium ions during charging/discharging, resulting in incomplete oxygen\nredox activity. It is generally agreed upon that the increase in Ni content will reduce the specific\ncapacity and the first cycle efficiency and result in an increase in Ni oxidation state from the\nreports from Qi et al (Electrochem 2023, 4, 21–30) and Chong et al (RSC Adv. 2016, 6, 53662).\nHere the authors show the similar phenomenon but with a reasonable explanation or\nmechanism for these conclusions, therefore I am unsure of the overall novelty - this should be\naddressed.\nOverall, this manuscript could be published after a minor revision.\nResponse: Thank you for the high evaluation of our work. In the following, we will address all\nthe comments sequentially.\n1. It is generally agreed upon that the increase in Ni content will reduce the specific capacity\nand the first cycle efficiency and result in an increase in Ni oxidation state from the reports from\nQi et al (Electrochem 2023, 4, 21–30) and Chong et al (RSC Adv. 2016, 6, 53662). Here the\nauthors show the similar phenomenon but with a reasonable explanation or mechanism for\nthese conclusions, therefore I am unsure of the overall novelty - this should be addressed.\nResponse: We appreciate the reviewer’s comment. As indicated by previous publications\n(Electrochem 2023, 4, 21–30; RSC Adv. 2016, 6, 53662; J. Mater. Chem. A. 2015, 3, 22199),\nthere has been considerable interest in elucidating the role of nickel content in the\nelectrochemistry of Li-rich layered oxides. Unfortunately, these studies primarily focused on\nelectrochemical characterization and did not provide a satisfactory explanation for the poor\nelectrochemical performance of high-nickel Li-rich layered oxides. Moreover, no universal\nsolutions to address this issue have been identified. Our research fills this gap by offering a\nreasonable explanation and potential solutions to improve the electrochemistry of high-nickel\nLi-rich layered oxides.\nBased on earlier studies, our work innovatively employed a series of advanced techniques,\nincluding SXRD, PDF, SXAS, RIXS, TEM (addressing Question 3 from Reviewer #2), NMR,\nDFT calculations, and electrochemical investigations. This comprehensive approach provides\ndefinitive answers to critical questions. Firstly, we identified the types of superstructure units\nwithin the transition metal layer and, for the first time, elucidate the influence of the nickel ions'\nvalence state on these units. More importantly, we successfully revealed the intrinsic\ncorrelation between oxygen redox activity and superstructure units, elucidating the mechanism\nbehind the poor discharge capacity observed in high-nickel Li-rich cathodes. Additionally,\nbeyond elucidating the mechanism, we proposed a potential solution to tailor superstructure\nunits to enhance battery performance with reversible oxygen redox.\nTherefore, we believe that our conclusive evidence on tailoring superstructure units for\nenhanced oxygen redox in Li-rich layered oxides will be of significant interest to both the battery\ncommunity and the broader readership of Nature Communications.\nChanges made:\nIn the revision manuscript, we inserted additional references in the introduction on page 3.\n20. Qi, G. et al. Impact of Ni Content on the Electrochemical Performance of the Co-Free, Li\nand Mn-Rich Layered Cathode Materials. Electrochem 4, 21–30 (2023).\n21. Knight, J. C. & Manthiram, A. Effect of nickel oxidation state on the structural and\nelectrochemical characteristics of lithium-rich layered oxide cathodes. J. Mater. Chem. A 3,\n22199–22207 (2015).\n2. For Figure 2d and 2f, the signals are very weak, and the noise is higher, how much\nconfidence you have for the data? Figure 2c, 2e and 2f, it seems that the authors only showed\nthree points, where is the fourth point data/curve?\nResponse:\nWe appreciate the reviewer’s comments and fully understand their concerns regarding the\ndata quality in the manuscript. In Fig. 2, our primary aim is to identify the redox nature of the\ndQ/dV peak observed below 3.7 V and elucidate the mechanism contributing to the poor\ncapacity in high-nickel Li-rich cathodes. As discussed in the manuscript, this peak is typically\nassociated with the reduction of oxygen (O) or manganese (Mn). Therefore, understanding the\nredox evolution of O and Mn, represented by changes in the O K-edge and Mn L-edge spectra,\nis crucial during charge and discharge.\nIn our study, we selected four points—pristine, charged to 4.45 V, charged to 4.7 V, and\ndischarged to 2.0 V—during the first cycle to collect the Ni, Mn L-edge, and O K-edge spectra\nwith soft-XAS spectroscopy (Fig. 2a-d). The spectra of the Mn L-edge significantly overlap in\nFig. 2c. To enhance clarity, we present the stack of Mn L-edge spectra in Fig. R2. Combining\nthese with the O K-edge SXAS results, we identify that the incomplete oxygen redox occurs at\na low potential (3.3 V vs. Li+/Li) in Li-rich layered oxides for a high-nickel system.\nGiven this abnormal phenomenon, we additionally conducted O K-edge RIXS measurements\n(Fig. 2e,f) to further validate our findings. As depicted in Fig. 2e-f, a distinct characteristic peak\nC (around 7.8 eV) and a vibrational peak D around the elastic region (0 eV) are still observed\nat discharge to 2.0 V. This finding is consistent with previous O K-edge SXAS results, further\nsupporting incomplete oxygen redox. Our data and findings are reasonable and well-justified.\nAdditionally, similar O K-edge RIXS signals are documented in previous studies (Nat Commun.\n2021, 12, 3071; Nat Commun. 2022, 13, 1123; J. Am. Chem. Soc. 2023, 18, 10208-10219;\nNat Energy. 2024, 9, 184–196).\nThe absence of data at 4.45 V for the RIXS measurements does not influence the conclusion\nof this work. Therefore, we selected three of the most critical states for RIXS measurements\n(excluding the point at 4.45 V) to address our research questions. We hope for the reviewer’s\nunderstanding.\nFig. R2. SXAS results of the Mn L-edge collected at pristine, charge to 4.45 V, charge to 4.7\nV, and discharge to 2.0 V states.\n3. To address these intrinsic properties in high-nickel Li-rich layered oxides, the authors\nproposed to reduce the Li content during synthesis, please explain the rational design, why Li\ncontent determine the amount of the LINiMn ? Also, why not just to optimize the Ni/Mn ratios\n5\ninstead?\nResponse: We appreciate the reviewer for the careful review on our submitted work.\nRegarding why Li content determine the amount of the LINiMn , our findings reveal that the\n5\ncathodes must be decreased. In our work, reducing the lithium content during synthesis\neffectively increases the overall transition metal (TM) content in the TM layer, as demonstrated\nby the increase from 0.80 in LM15 (Li Ni Mn O ) to 0.87 in LM13 (Li Ni Mn O ).\n1.2 0.36 0.44 2 1.13 0.39 0.48 2\nXAS results confirm the decrease in the valence state of Ni, leading to a corresponding change\nin superstructure units. Specifically, the LiNiMn /(LiNiMn + LiMn ) peak ratio decreases to\n5 5 6\n42.7% for LM13, compared to 58.2% for LM15, as evidenced by NMR. As described in the\nresponse to Question 1 from reviewer #2, for a given Ni/Mn ratio, the LiNiMn superstructure\n5\ncontent is determined by the Li content. This is because the Ni valence state changes as the\nLi content changes.\nRegarding, why not just to optimize the Ni/Mn ratios instead, we thank the reviewer for pointing\nout this method for tailoring LiNiMn units. We fully acknowledge the reviewer's suggestion to\n5\ntailor LiNiMn units by further optimizing the Ni/Mn ratios in Li-rich cathodes. As evident from\n5\nthe electrochemical data (Fig. 1c-d) and the changes in superstructure units (Fig. 3a) observed\nin these cathodes with varying nickel contents, there may indeed be an optimal Ni/Mn ratio that\ncan achieve both high capacity and high voltage.\nHowever, our primary aim is to elucidate the underlying mechanism behind the unsatisfactory\nelectrochemistry observed in high-nickel Li-rich layered oxides. Therefore, we have chosen a\npractical and simple approach to validate the accuracy of our theoretical findings. Additionally,\nimplementing the method of optimizing Ni/Mn ratios may require a significant number of\nprecursor synthesis experiments, which falls beyond the scope of the current manuscript.\nMoreover, adjusting the lithium content provides additional benefits, such as modifying the\nsuperstructure content to fine-tune cation/anion redox contributions. This adjustment promotes\nhigher cation redox, thereby enhancing structural stability compared to the instability\nassociated with anion redox at high voltage. Moreover, moderate Li/Ni disordering promotes\nstructural reversibility due to the pillar effect. Finally, this approach represents a practical and\nsimple solution for industry applications when compared to other complex modification\nprocesses. Considering all these factors, the rationale for reducing Li content is reasonable\nand easily justified.\nReviewer #2 (Remarks to the Author):\nLi-rich layered oxides have high capacity because of high-voltage oxygen redox activity. Li-\nrich cathodes with a high nickel content exhibit improved voltage stability but suffer from poor\ndischarge capacity. This work tailors the LiNiMn superstructure content by modification of Ni\n5\nand Li contents. The study finds that an excess of LiNiMn impedes the extraction/insertion of\n5\nlithium ions during charging/discharging, resulting in incomplete oxygen redox activity\n(retention of molecular O ) at low potential (~3.3 V, versus Li+/Li). If the evidence is sufficient,\n2\nthe work is more interesting to open a new idea for enhanced oxygen redox of Li-rich layered\noxides. There are several puzzling problems, so it is not recommended for publication:\nResponse: Thank you for the careful and in-depth review of our manuscript. In the following,\nwe will address all the comments and revise the manuscript accordingly.\n1. In Li Ni Mn O , x = 0.28, 0.32, 0.36, 0.40 materials, the content of Ni with +4 valence\n1.20 x 0.8-x 2\nincreases with the increase of x or the decrease of Li content. The work finds that the LiNiMn\n5\nsuperstructure content increases with the increase of x or Li content. This is contradictory. The\nLiNiMn superstructure content is determined by the Li content or the ratio of Li/TM? Please\n5\ngive a reasonable interpretation.\nResponse: We appreciate the reviewer's comment. Our findings reveal that the LiNiMn\n5\nsuperstructure units are exclusively stabilized by Ni4+, rather than by Ni2+ or Ni3+.\nThis\nstabilization is facilitated by the presence of Ni4+, which promotes the formation of LiNiMn\n5\nunits in Li-rich cathodes.\nIn materials such as Li Ni Mn O (x = 0.28, 0.32, 0.36, 0.40), with a Li/TM ratio of 1.5, an\n1.20 x 0.8-x 2\nincrease in x results in an increase in the valence state of Ni. As a consequence,\nLi Ni Mn O (N40), with a higher Ni valence state, contains a greater proportion of\n1.2 0.40 0.40 2\nLiNiMn units compared to Li Ni Mn O (N28).\n5 1.2 0.28 0.52 2\nFurthermore, for a given Ni/Mn ratio, reducing the lithium content during synthesis effectively\nincreases the overall transition metal (TM) content in the TM layer, as demonstrated by the\nincrease from 0.80 in LM15 (Li Ni Mn O ) to 0.87 in LM13 (Li Ni Mn O ). As a\n1.2 0.36 0.44 2 1.13 0.39 0.48 2\nresult, there are fewer LiNiMn units in the LM13 with lower valence state of Ni.\n5\nthe Li content or the ratio of Li/TM; rather, it changes only when the Ni valence state changes.\n2. In conventional LLOs such as Li Ni Mn O , generally, based on Pauling´s electrostatic\n1.2 0.2 0.6 2\nvalence rule, LiMn configurations are preferably aggregated in Li MnO due to their special\n6 2 3\nstability. So it is difficult to permeate Ni to Li MnO . Many researches found that the breakdown\n2 3\nof aggregated Li MnO , such as the permeation of TM with strong TM-O bonds, benefited the\n2 3\nenhanced oxygen redox. Though the work gives evidence of LiNiMn configurations, why can\n5\nthe LiNiMn configurations exist in the work? Why is a reason that LiNiMn configurations result\n5 5\nin incomplete oxygen redox activity?\nResponse: We are grateful to the reviewer for pointing out Pauling's electrostatic valence\nprinciple as a demonstration of the stability of the superstructure configuration. Indeed, we fully\nagree with the reviewer that the LiMn units are more likely to form in Li MnO compared to\n6 2 3\nother superstructure units. Furthermore, previous studies (Matter. 2022, 5, 3869-3882;\nMaterials Today. 2021, 51, 15-26) also suggest that the Ni is difficult to detect within Li MnO -\n2 3\nlike domains. However, the above conclusion is only based on the case in which Ni has a\nvalence state of 2 or 3. The possibility of the presence of Ni4+ in LiNiMn superstructure units\n5\nhas been previously overlooked. Note that while Ni4+ is atypical in conventional layered\ncathodes, its presence in Li-rich cathode systems is plausible due to the lithium excess\nstructural configuration (Chem. Mater. 2020, 32, 9211–9227).\nBased on Pauling’s electrostatic valence principle (J. Am. Chem. Soc. 1929, 51, 1010-1026)\nwe can apply a semiquantitative equation to evaluate the stability of a coordination\nconfiguration (Materials Today. 2021, 51, 15-26):\n𝛧𝑐\n∆𝛧 = |𝛧 −Σ |\n𝛢 𝑛\nAccording to the principle of local electrical neutrality, a structure may become unstable if ∆𝛧 >\n0. In Li-rich cathodes, three Li ions in the Li layer are consistently coordinated with central\noxygen atom, while the coordination environments of cations in the transition metal layer may\nhave different combinations. The most common LiMn units, equivalent to Li -O-LiMn ,\n6 3 2\nexhibit\nstable state with ∆Ζ = 0. In our work, LiNiMn units (Li -O-LiNiMn) also demonstrate ∆Ζ = 0,\n5 3\nattributed to the presence of Ni4+ rather than Ni2+ (∆Ζ = 1/3) and Ni3+ (∆Ζ = 1/6). Moreover,\nour 6Li NMR results clearly demonstrate the presence of these LiNiMn units in these Li-rich\n5\ncathodes.\nRegarding Why the LiNiMn configurations result in incomplete oxygen redox activity, previous\n5\ntheoretical calculations (Matter. 2022, 5, 3869-3882) indicate that the Li MnNi-coordinated O,\n4\nreferred to as LiNiMn in our work, has a much lower cationic and anionic redox activity\n5\ncompared to LiMn units, suggesting a poor electrochemical activity. Additionally, Li moves\n6\nfrom one octahedral site to another octahedral site by passing through an intermediate\ntetrahedral site where it encounters strong repulsion from a nearby transition metal. The\nassociated activation barriers for Li movement with Ni4+, Ni3+ and Ni2+ are 490, 310, and 210\nmeV, respectively (Science. 2006, 311, 977-980). Therefore, these theoretical findings suggest\nthat Ni4+-stabilized LiNiMn may impede the movement of lithium ions during\n5\ncharging/discharging. Moreover, Li-rich cathodes undergo a cationic–anionic redox inversion\nwith significant chemical and electrochemical asymmetry between charge and discharge\n(Nature Mater. 2022, 21, 1165-1174). Consequently, at low potential during discharge, some\nLi ions fail to return to the transition metal layer to coordinate with O, resulting in incomplete\noxygen redox. This observation is further supported by our experimental results, including\nSXAS, RIXS, ex-situ NMR, in-situ XRD, and electrochemistry.\nChanges made:\nIn the revised manuscript, we have included additional content on page 10: “In addition,\nPauling's electrostatic valence principle33 and semi-quantitative equation34 can be used to\nevaluate the possibility of a coordination configuration. According to principle of local electrical\nneutrality, a structure may become unstable if the local difference in valences of anions/cations\n∆𝛧 > 0. The most common LiMn units, equivalent to Li -O-LiMn , exhibit a stable state with\n6 3 2\n∆Ζ = 0. Similarly, LiNiMn units (Li -O-LiNiMn) also demonstrate ∆Ζ = 0, attributed to the\n5 3\npresence of Ni4+ rather than Ni2+ (∆Ζ = 1/3) and Ni3+ (∆Ζ = 1/6). This elucidates why prior\ninvestigations failed to detect nickel within Li MnO -like domains, given that the possibility of\n2 3\nthe presence of Ni4+ in LiNiMn superstructure units has been previously overlooked34,35.\n5\nto LiMn units, indicating limited electrochemical activity. Moreover, lithium ions move from one\n6\noctahedral site to another one, passing through an intermediate tetrahedral site where they\nencounter strong repulsion from nearby transition metals. The associated activation barriers\nfor lithium movement with Ni4+, Ni3+ and Ni2+ have been calculated to be 490, 310, and 210\nmeV, respectively36. These theoretical findings imply that LiNiMn may impede the\n5\nextraction/insertion of lithium ions. During charging/discharging, Li-rich cathodes experience a\ncationic–anionic redox inversion, leading to electrochemical asymmetry23. As a consequence,\nat low discharge potentials, some lithium ions fail to return to the transition metal layer to\ncoordinate with O, resulting in incomplete oxygen redox. This analysis is consistent with the\nfindings from our O K-edge SXAS and RIXS experiments.”\n23. Li, B. et al. Capturing dynamic ligand-to-metal charge transfer with a long-lived cationic\nintermediate for anionic redox. Nat. Mater. 21, 1370–1379 (2022).\n33. Pauling, L. The principles determining the structure of complex ionic crystals. J. Am. Chem.\nSoc. 51, 1010–1026 (1929).\n34. Yin, C. et al. Structural insights into composition design of Li-rich layered cathode materials\nfor high-energy rechargeable battery. Materials Today 51, 15–26 (2021).\n35. Yang, Y. et al. Cation configuration in transition-metal layered oxides. Matter 5, 3869–3882\n(2022).\n36. Kang, K., Meng, Y. S., Bréger, J., Grey, C. P. & Ceder, G. Electrodes with High Power and\nHigh Capacity for Rechargeable Lithium Batteries. Science 311, 977–980 (2006).\n3. The work does not provide TEM images, thus the other existence of Ni is unclear, and for\nexample, the occupation of less Ni at Li-layer sites reduced oxygen release and lowered\nvoltage decay.\nResponse: We appreciate the reviewer's comment and suggestion. To investigate the\npotential influence of cation mixing (where Ni occupies Li sites) on the electrochemistry of Li-\nrich cathodes with varying nickel contents, we conducted transmission electron microscopy\n(TEM) measurements. The high-resolution TEM (HRTEM) images of N28, N32, N36, and N40,\nas depicted in Fig. R3, reveal clear lattice fringes without obvious evidence of cation mixing.\nMoreover, the observed average interplanar spacing of approximately 0.47 nm corresponds to\nthe (003) or (001) plane of the layered rhombohedral (R-3m) or monoclinic (C2/m) structure.\nR M\nAdditionally, integrating the previously conducted high-resolution synchrotron X-ray diffraction\n(SXRD) measurements, we conclude that cation mixing has little impact on the\nelectrochemistry of these cathodes. This further underscores the close correlation between\nhoneycomb superstructure units and electrochemistry in our work.\nChanges made:\nIn the revised manuscript, we have included additional content on page 6 in the manuscript\nand page 5 in the supporting information “The high-resolution transmission electron\nmicroscopy (HRTEM) images of N28, N32, N36, and N40, as shown in Supplementary Fig. 4,\nreveal clear lattice fringes without obvious evidence of cation mixing. Moreover, the observed\naverage interplanar spacing of approximately 0.47 nm corresponds to the (003) or (001)\nR M\nplane of the layered rhombohedral (R-3m) or monoclinic (C2/m) structure.”\nAdditionally, we have included additional content on the page 20 in the Methods section\n“Structural information about the samples was obtained by high-resolution transmission\nelectron microscopy (HRTEM). Each specimen was examined using Thermofisher Talos\nF200X TEM operated at 200 kV.”\nFig. R3. HRTEM images of N28 (a-b), N32 (c-d), N36 (e-f), and N40 (g-h) powder.\n4. The DOS of Li Mn Ni O (Li Mn Ni O ) is employed using DFT calculations. The\n44 18 10 72 1.22 0.50 0.28 2\nstate of Ni4 is +4, as expected in the work’s model. Other models without the LiNiMn\n5\nconfigurations are stable?\nResponse: We appreciate the reviewer's comment. We also performed calculations on a\nmodel without the LiNiMn unit of Li Mn Ni O (Li Mn Ni O ). The energy difference\n5 44 18 10 72 1.22 0.50 0.28 2\nbetween this model and the one containing the LiNiMn unit was minimal, approximately 0.005\n5\neV/atom. As a result, it can be concluded that their energies are the same, and they can be\nstable.\n5. In addition, 231.1 mAh g-1 at C/10 and a capacity retention of 84.8% after 200 cycles at C/3\nis not high, such as 276.5 mAh g−1 at C/10 and 85.4% after 300 cycles at 1 C of\nLi Ni Mn O (Li J, ACS Nano, 2023, 17, 16827-16839).\n1.2 0.2 0.6 2\nResponse: We appreciate the reviewer for the careful review on our submitted work. We fully\nagree with the reviewer that our cathode (LM13) exhibits lower discharge capacity compared\nto the classical Li-rich cathode (Li Ni Mn O ). Despite achieving high capacity and stability\n1.2 0.2 0.6 2\nby various modification in recent years, these Li-rich cathodes still encounter substantial\nvoltage decay during cycling. As highlighted in the introduction, optimizing the chemical\ncomposition has emerged as a promising approach to mitigate voltage decay in layered Li-rich\ncathodes (Adv. Energy Mater. 2018, 8, 1800606; ACS Appl. Mater. Interfaces. 2016, 8, 20138–\n20146). Additionally, high-nickel Li-rich layered oxides offer several advantages in stabilizing\nvoltage and reducing oxygen release (Nat. Mater. 2023, 22, 1370–1379). However, high-nickel\nLi-rich cathodes, such as Li Ni Mn O , have garnered less attention due to their lower\n1.2 0.4 0.4 2\ndischarge capacity (less than 200 mAh g-1 at C/10) compared to low-nickel Li-rich cathodes\n(Electrochem 2023, 4, 21–30; RSC Adv. 2016, 6, 53662; J. Mater. Chem. A. 2015, 3, 22199).\nThus, the key to advancing Li-rich cathodes is achieving voltage stability in high-nickel systems\nwithout compromising discharge capacity.\nTherefore, the primary aim of this work is to elucidate the underlying mechanism behind the\nunsatisfactory electrochemistry observed in high-nickel Li-rich layered oxides. Building on the\ninsights gained from the mechanism, we propose a practical and straightforward solution to\ntailor superstructure units, thereby enhancing oxygen redox reversibility and improving battery\nperformance (from 199.0 to 231.1 mAh g-1 at C/10). Our strategy has proven to be efficient\ncompared to pristine high-nickel Li-rich cathodes. As described in the discussion in the\nmanuscript, adjusting the lithium content during synthesis, in addition to changes in\nsuperstructure units, affects other parameters such as superstructure content, Li/Ni disorder,\nand stacking faults. Therefore, future research endeavours may focus on exploring these\nparameters and achieving a balance to further enable high-energy-density, high-nickel Li-rich\nlayered oxides."}, {"round": 2, "reviewer_comments": {"reviewer_1": "After carefully reading the revised the manuscript, I have decided to approve the manuscript to proceed with the publication.\nThanks and let me know if you have any questions.\nBest,\nGongshin", "reviewer_2": "The work is carefully modified. But there is one puzzling problem.\n--The DOS of Li44Mn18Ni10O72 (Li1.22Mn0.50Ni0.28O2) is employed using DFT calculations. The state of Ni4 is +4, as expected in the work’s model. So the structural stability with one LiNiMn5 configuration instead of the state of Ni4 with +4 should be explored, because the chemical valence of Ni4 is related to the structure.\nIn addition, to improve the redox reversibility of lattice oxygen and cycling performance, recently two works (ACS Nano 2023, 17, 16827−16839; Nano Energy, 2024, 123, 109390) are done though tuning the Li2MnO3 domain, please cite them."}, "author_response": "Point-to-point Response to Reviewers’ Comments\nTailoring Superstructure Units for Enhanced Oxygen Redox in Li-Rich Layered\nOxides\n(Manuscript ID NCOMMS-24-16929A)\nWe sincerely thank the editor and all reviewers for their valuable review comments that we\nhave used to improve the quality of our manuscript. The reviewer comments are laid out below\nand specific concern have been numbered, and all the responses are listed in the following\nsection, point by point.\nReviewer #1 (Remarks to the Author):\nAfter carefully reading the revised the manuscript, I have decided to approve the manuscript\nto proceed with the publication.\nThanks and let me know if you have any questions.\nBest,\nGongshin\nResponse: We highly appreciate the reviewer’s positive comments.\nReviewer #2 (Remarks to the Author):\nThe work is carefully modified. But there is one puzzling problem.\n--The DOS of Li Mn Ni O (Li Mn Ni O ) is employed using DFT calculations. The\n44 18 10 72 1.22 0.50 0.28 2\nstate of Ni4 is +4, as expected in the work’s model. So the structural stability with one LiNiMn\n5\nconfiguration instead of the state of Ni4 with +4 should be explored, because the chemical\nvalence of Ni4 is related to the structure.\nIn addition, to improve the redox reversibility of lattice oxygen and cycling performance,\nrecently two works (ACS Nano 2023, 17, 16827−16839; Nano Energy, 2024, 123, 109390) are\ndone though tuning the Li MnO domain, please cite them.\n2 3\nResponse: Thank you for the careful and in-depth review of our manuscript. In the following,\nwe will address all the comments and revise the manuscript accordingly.\n1. The DOS of Li Mn Ni O (Li Mn Ni O ) is employed using DFT calculations. The\n44 18 10 72 1.22 0.50 0.28 2\nstate of Ni4 is +4, as expected in the work’s model. So the structural stability with one LiNiMn\n5\nconfiguration instead of the state of Ni4 with +4 should be explored, because the chemical\nvalence of Ni4 is related to the structure.\nResponse: We appreciate the reviewer's comment. We have calculated the oxidation state of\nNi in the LiNiMn configuration without a priori assumptions on the Ni oxidation state. The\n5\nrelaxed structure shows that Ni possesses the oxidation state +4. In addition, we have\ncalculated models with the same overall composition but without the LiNiMn configuration\n5\n(always without assuming the oxidation state a priori), in which we find that Ni has an oxidation\nstate lower than +4. Based on that we conclude the LiNiMn configuration stabilize Ni in the\n5\noxidation state +4 in the structure.\n2. In addition, to improve the redox reversibility of lattice oxygen and cycling performance,\nrecently two works (ACS Nano 2023, 17, 16827−16839; Nano Energy, 2024, 123, 109390) are\ndone though tuning the Li MnO domain, please cite them.\n2 3\nResponse: We are grateful to the reviewer for pointing out the importance of the tuning of\nLi MnO domain size or distribution to the oxygen redox reversibility of Li-rich cathodes. To\n2 3\nbetter reflect this recent progress, we have added the relevant content in the Introduction and\ncited the relevant studies.\nChanges made:\nIn the revised manuscript, we have included additional content on page 3: “For instance, a\nconsensus has emerged form relevant studies, highlighting the detrimental role of localized\nsuperstructure domains characterized by in-plane Li/Mn order of Li-rich cathodes, while\ndelocalized or dispersed domains characterized by in-plane Li/Mn disorder effectively improve\nthe oxygen redox reversibility and voltage stability10–12.”\n10. Li, J. et al. Tuning Li MnO -Like Domain Size and Surface Structure Enables Highly-\n2 3\nstabilized Li-Rich Layered Oxide Cathodes. ACS Nano 17, 16827–16839 (2023).\n11. Guo, X. et al. Weak σ–π–σ interaction stabilizes oxygen redox towards high-performance\nLi-rich layered oxide cathodes. Nano Energy 123, 109390 (2024).\n12. Zhang, M. et al. Formulating Local Environment of Oxygen Mitigates Voltage Hysteresis in\nLi-Rich Materials. Advanced Materials 36, 2311814 (2024)."}, {"round": 3, "reviewer_comments": {"reviewer_2": "it is recommended for publication."}, "author_response": ""}]

Decision letter