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O3- vs P2-type Nax(Ni,Zn,Mn,Ti)O2 layered oxides:
Comparative study on electrode-electrolyte reactivity
and structural stability for cycling performance

Juan Forero-Saboya, Yue Zhou, Stephen Browne, Ivan Moiseev, Chloe Pablos, John Abou-Rjeily, Arame Mboup, Clémence Alphen, Leiting Zhang, Biao Li, et al.

► To cite this version:

Juan Forero-Saboya, Yue Zhou, Stephen Browne, Ivan Moiseev, Chloe Pablos, et al. O3- vs P2-type Nax(Ni,Zn,Mn,Ti)O2 layered oxides: Comparative study on electrode-electrolyte reactivity and structural stability for cycling performance. Energy Storage Materials, 2025, 80, pp.104423. 10.1016/j ensm.2025.104423. hal-05308175

HAL Id: hal-05308175

https://hal.science/hal-05308175v1

Submitted on 13 Oct 2025

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Distributed under a Creative Commons Attribution 4.0 International License
O3- vs P2-type Naₓ(Ni,Zn,Mn,Ti)O₂ layered oxides: Comparative study on
electrode-electrolyte reactivity and structural stability for cycling
performance.

Juan Forero-Saboya,<fnref n="1" />,<fnref n="2" />,# Yue Zhou,<fnref n="3" />,<fnref n="4" />,# Stephen Browne,<fnref n="1" />,<fnref n="2" /> Ivan A. Moiseev,<fnref n="5" /> Chloe Pablos,<fnref n="6" />
John Abou-Rjeily,<fnref n="6" /> Arame Mboup,<fnref n="6" /> Clemence Alphen,<fnref n="1" />,<fnref n="2" />,<fnref n="7" /> Leiting Zhang,<fnref n="8" /> Biao Li,<fnref n="3" /> Artem M.
Abakumov,<fnref n="5" /> Jean-Marie Tarascon<fnref n="1" />,<fnref n="2" />,<fnref n="7" /> and Sathiya Mariyappan<fnref n="1" />,<fnref n="2" />,<fnref n="7" />,*

**Abstract.** Sodium layered oxides, having either O3, P2 or P3 stacking, are extensively studied as low-cost cathode materials for high energy Na-ion batteries (NIBs). Previous efforts from our group focused on the optimization of layered oxide compositions, identifying O3-Na₀.85Ni₀.38Zn₀.04Mn₀.48Ti₀.₁O₂ and P2-Na₀.67Ni₀.3Zn₀.03Mn₀.52Ti₀.₁O₂ phases as potential candidates to establish prototype cylindrical 18650 cells with 120-150 Wh/kg specific cell energy. In this study, we focus particularly on the electrode-electrolyte reactivity of these phases, especially at high state of charge (~70% or more) and at high temperatures. Our results indicate that the end-of-charge phase, O1 and O2 formed during complete de-sodiation of O3 and P2, respectively, plays a major role in determining their reactivity. The O1 phase is particularly prone to transition metal migration and oxygen oxidation, having increased reactivity with electrolyte. On the other hand, the P2 layered oxide, while having lower capacity than O3, offers better cycling stability (90% retention after 1000 cycles at 25 °C) due to the greater stability of the O2 end-of-charge structure. These results once again underline the fact that specific capacity should not be the sole metric for determining the most suitable electrode materials for Na-ion or other battery chemistries.

**Keywords:** Sodium-ion batteries, layered oxides, O3 vs P2 phase, cylindrical 18650, interfacial reactivity, interphase, electrolyte degradation

<footnote n="1">¹Chimie du Solide-Energie, UMR 8260, Collège de France, 75231 Paris Cedex 05, France</footnote>

<footnote n="2">² Réseau sur le Stockage Electrochimique de l’Energie (RS2E), FR CNRS 3459, France</footnote>

<footnote n="3">³ Beijing Key Laboratory of Theory and Technology for Advanced Batteries Materials, School of Materials Science and Engineering, Peking University, Beijing 100871, P. R. China</footnote>

<footnote n="4">⁴ The Hong Kong University of Science and Technology (Guangzhou) Sustainable Energy and Environment Thrust and Guangzhou Municipal, Key Laboratory of Materials Informatics, Nansha, Guangzhou, Guangdong 511400, P. R. China</footnote>

<footnote n="5">⁵ Center for Energy Science and Technology, Skolkovo Institute of Science and Technology, Bolshoy Boulevard 30, bld. 1, Moscow, 121205 Russia</footnote>

<footnote n="6">⁶ TIAMAT Energy, Hub de l'énergie, rue Baudelocque, 80000 Amiens, France</footnote>

<footnote n="7">⁷ Sorbonne Université, 4 Place Jussieu, 75005, Paris, France</footnote>

<footnote n="8">⁸ Department of Chemistry – Ångström Laboratory, Uppsala University, Box 538, Uppsala, SE-751 21 Sweden</footnote>

<footnote># These authors contribute equally to this study</footnote>

<footnote>*Correspondence: sathiya.mariyappan@college-de-france.fr</footnote>
# 1. Introduction

Sodium-ion batteries have entered the market with diverse prototypes being developed globally for applications ranging from large-scale energy storage to high-power tools and automotive systems. The specific demands of these applications are better met by different cathode materials, with the most extensively studied being Prussian blue analogues, polyanionic compounds, and layered oxides. Among these, layered oxides display higher capacity, lower cost, and superior volumetric energy density.

Sodium layered oxides can be synthesized in various structures that can be classified as P-type or O-type.[1] Both families having distinct advantages, but neither emerging as the definitive choice for all applications. P-type materials are often favored for high-power applications due to their high Na⁺ mobility through adjacent prismatic sites, whereas O-type materials offer higher capacity owing to their greater sodium content.[2] Additional differences in moisture stability, cost, and usable potential window have been previously discussed.[3]

Two of the most studied sodium layered oxides are P2-Na<sub>0.67</sub>Ni<sub>0.33</sub>Mn<sub>0.67</sub>O<sub>2</sub> and O3-NaNi<sub>0.5</sub>Mn<sub>0.5</sub>O<sub>2</sub>, although cycling either of them up to their full desodiation (above 4 V) results in severe capacity fading.[4-6] Such degradation is often ascribed to a series of complex phase transitions, Na⁺/vacancy ordering, and MO<sub>2</sub> slab gliding during sodiation/desodiation, particularly at high state-of-charge. A common strategy to reduce the effect of such phase transitions during cycling relies on doping the material with small amounts of Zn and Ti, a strategy that our group has previously explored resulting in the optimized compositions: P2-Na<sub>0.67</sub>Ni<sub>0.3</sub>Zn<sub>0.03</sub>Mn<sub>0.52</sub>Ti<sub>0.15</sub>O<sub>2</sub> (P2-ZNMT) and O3-Na<sub>0.85</sub>Ni<sub>0.38</sub>Zn<sub>0.04</sub>Mn<sub>0.48</sub>Ti<sub>0.1</sub>O<sub>2</sub> (O3-ZNMT), displaying much higher stabilities and rendering them attractive for application in commercial Na-ion cells.[3]

Some far less understood parasitic reactions are rooted at the cathode-electrolyte interphase. At high state-of-charge, evidences of transition metal migration to the Na⁺ layer, lattice oxygen oxidation, oxygen release, and electrolyte oxidation have been observed.[7,8] However, the interplay between these processes and their collective impact on cycling stability requires further attention, particularly when comparing the relative stabilities of P-type vs. O-type cathodes at high state-of-charge.

In this study we explore the use of our previously optimized cathode compositions in coin and 18650 Na-ion full cells, with hard carbon (HC) as a negative electrode. Our studies show that the structure of the end-of-charge phase (O1 and O2 for the O3 and P2 pristine materials, respectively) plays a key role in determining the interfacial reactivity and the long-term cycling stability. The end-of-charge O1 phase, being more reactive with the electrolyte, leads
to poor cycling stability of the O3-ZNMT|HC cells compared to the P2-ZNMT|HC cells. As a result, O3-ZNMT, though exhibits higher capacity than P2-ZNMT, must be cycled within a limited voltage window (to reduce O1 phase formation), which limits its achievable capacity and energy. These results underline the importance of taking interfacial reactivity into account when designing cell chemistry, in addition to structural stability in the bulk.

## 2. Results:

### 2.1. Electrochemical performance of O3 and P2 sodium layered oxides in Na-ion full cells:

The O3-Na<sub>0.85</sub>Ni<sub>0.38</sub>Zn<sub>0.04</sub>Mn<sub>0.48</sub>Ti<sub>0.1</sub>O<sub>2</sub> and P2-Na<sub>0.67</sub>Ni<sub>0.3</sub>Zn<sub>0.03</sub>Mn<sub>0.52</sub>Ti<sub>0.15</sub>O<sub>2</sub> sodium layered oxides, represented hereafter as O3-ZNMT and P2-ZNMT respectively, were synthesized by solid state synthesis as reported elsewhere,[3] and the phase purity of the obtained solids was corroborated by powder XRD (Figure S1a,b). When charging the material to 4.4 V vs. Na<sup>+</sup>/Na<sup>0</sup>, a capacity of 173 mAh g<sup>-1</sup> (0.7 Na) for O3-ZNMT and 131 mAh g<sup>-1</sup> (0.55 Na) for P2-ZNMT can be extracted (Figure S1c). This represents around 82% of the theoretical capacity in both cases (212 and 174 mAh g<sup>-1</sup>, respectively).

The long cycling performance of both cathode materials was tested in full Na-ion cells using hard carbon (HC) as anode active material, 1 M NaPF<sub>6</sub> in EC:PC:DMC (1: 1: 2 by volume ratio) as electrolyte and cycling between 2-4.25 V, at 25 °C. Thus assembled Na-ion full cells O3-ZNMT|HC and P2-ZNMT|HC were used for all our electrochemical analyses discussed hereafter, unless otherwise mentioned.

When cycled at room temperature, the O3-ZNMT|HC and P2-ZNMT|HC full cells retain ~80% capacity after 200 cycles (Figure S2), similar to reported in previous work.[3] To accelerate cell degradation and exacerbate secondary reactions at the electrode-electrolyte interface, we used a high temperature (55 °C) cycling protocol, the performance of which is summarized in **Figure 1**. Under these conditions, full cells lose 20% of their initial capacity after only 30 cycles. Since we did not use any electrolyte additive at this stage, the capacity fading could be attributed to a poor SEI on the HC electrode, a situation that is well known in full Na-ion cells.[9] We have therefore included a self-discharge step at cycle 10, where the cells were rested at 100% state-of-charge for one week and the capacity loss during this calendar ageing is recorded (see Figure S3, S4 and Supporting Note S1 in supporting information for more details). Both O3-ZNMT|HC and P2-ZNMT|HC cells exhibit only 7-8% capacity loss during self-discharge, out of which nearly 96% is recovered during the following cycles, performing significantly better than our previous studies on polyanionic
$Na_3V_2(PO_4)_2F_3$ (NVPF)|HC full cells which presented 15.4% self-discharge loss and 90.7% recovery.[10]

<figure><img src="image_3.png" /><figcaption>**Figure 1.** Cycling performance of O3-ZNMT|HC and P2-ZNMT|HC full cells in 2032 coin type cell configuration at 55°C. (a,d) Voltage profiles in the first 10 cycles, (b,e) Capacity retention upon cycling, including a self-discharge (noted as ‘SD’ in the figure) step at cycle 10 (inset: differential capacity plots of the cells at cycle 10, 12, and 100), and (c,f) Average charging and discharging voltages of the cells (inset: voltage profiles as function of normalized capacity).</figcaption></figure>

The greater capacity retention of layered oxide cells during calendar ageing is the result of the different voltage profiles of the different materials. While the P2-ZNMT|HC and O3-ZNMT|HC full cells relax rapidly to a voltage close to 4 V during the self-discharge period, the NVPF|HC cells remain above 4.1 V due to the high voltage plateau of the NVPF. This difference may be sufficient to account for continuous electrolyte oxidation in the NVPF cells, and thus higher self-discharge. This self-discharge step does not produce obvious changes in the structures of O3-ZNMT and P2-ZNMT, as shown by the comparison of the differential capacity plot of cycles 10 and 12 (inset **Figure 1b and 1e**).

However, during continuous cycling, the high-voltage plateau in both P2-ZNMT and O3-ZNMT materials is suppressed due to a significant increase in the polarization (defined as $\Delta V = \bar{V}_{charge} - \bar{V}_{discharge}$), as shown by the plots of average voltage versus number of cycles illustrated in **Figure 1c and 1f**. If we focus on the first 10 cycles, however, a
simultaneous decrease in both charge and discharge voltage is observed for both materials, and more severely for O3-ZNMT. Similar voltage decay has been previously observed in lithium and sodium layered oxide cathodes and has been ascribed to permanent transition metal migration, lattice oxygen loss, and more generally to a permanent damage of the structure of the material.[11] This led us to investigate more deeply the degradation occurring in these two materials during the first few cycles, and particularly upon holding at high potential for extended periods of time.

## 2.2. Investigating the underlying degradation mechanisms

We start by analyzing the gas evolution in full O3-ZNMT|HC and P2-ZNMT|HC cells by Online Electrochemical Mass Spectrometry (OEMS). The production of C₂H₄, CO₂, O₂ and H₂ from the full cells is shown in **Figure 2a, and 2b**. An important generation of ethylene and hydrogen gases is observed in both cells at low voltage (< 3 V) corresponding to the reduction of EC and protic species at the HC electrode as previously observed in Li-ion batteries using similar electrolytes.[12,13] At high voltage (> 3.8 V), a significant amount of CO₂ was observed in both O3-ZNMT|HC and P2-ZNMT|HC cells, being nearly double for the former, 83 µmol g⁻¹ vs 44 µmol g⁻¹. Though no O₂ evolution was observed, the CO₂ release clearly indicates electrolyte oxidation by the layered oxide material, being more prominent for O3 than for P2. Although some amount of the produced CO₂ could potentially be ascribed to residual Na₂CO₃ in the material, we do not think this is the case here. Sodium carbonate and hydroxide on the surface of sodium layered oxides are very common hydrolysis products when the materials are exposed to moisture (Figure S5a,b), but in this study we were careful to always treat the material in a dry atmosphere. Additionally, to corroborate that the material does not contain excessive amounts of carbonate, we analyzed the first charge derivative plot of all the cells. The oxidation of Na₂CO₃, if present, would appear as a clear sharp peak near 4 V vs Na⁺/Na⁰ (Figure S5c). The material used in this study does not show this feature in any of our experiments, as evidenced in Figure S6, thus we believe that the CO₂(g) observed in the OEMS experiment does not have significant contribution from Na₂CO₃ but is related mostly to carbonate solvent chemical/electrochemical oxidation.
<figure><img src="image_4.png" /><figcaption>**Figure 2.** Gas evolution of the (a) O3-ZNMT|HC and (b) P2-ZNMT|HC during the first charge at room-temperature. (c) Cycling protocol of the cells to study cell degradation at high state-of-charge, consisting of one full cycle and one charge at C/10, followed by a constant voltage step for 60 hours. (d) Recorded current during the constant voltage step, normalized by the current employed during charge: 100% = C/10. (e) Quantified Ni dissolved from the cathode material at different holding potentials as displayed in panel (c). (f) ¹H NMR spectra of the pristine electrolyte and the electrolytes recovered from the cells after the 60 hours hold experiment, the peaks marked with * denote electrolyte degradation products.</figcaption></figure>

To understand better the interfacial reactions between the material and the electrolyte, we designed our next experiments with a CC-CV (constant current- constant voltage) protocol in which the O3-ZNMT|HC and P2-ZNMT|HC cells were charged to different state-of-charge (4 V, 4.2 V and 4.4 V) at constant current, followed by a constant voltage step for 60h. A representative experimental protocol is shown in **Figure 2c** with the observed residual current during the CV step shown in **Figure 2d**. The current values are normalized with respect to the current used for the charging step (100% = C/10). During the voltage hold, the current quickly decreases for the first 20 hours, as the electrodes reach the equilibrium Na⁺ composition at the given potential. Beyond that time, we assume that the current recorded corresponds only to
parasitic redox reactions occurring at the electrodes, namely oxidation, reduction, or redox shuttling. This parasitic current remains relatively low if the voltage is kept below 4.2 V for both materials, but it increases importantly when the cells are charged at 4.4 V, with residual current at all voltages always slightly higher for the O3-ZNMT|HC cell than for P2-ZNMT|HC.

The transition metal dissolution during this voltage-hold step was analyzed by ICP-MS.

**Figure 2d** shows dissolved Ni as a function of the cycling voltage (or state of charge) for both P2 and O3 materials. At low voltage, we observe minimal TM dissolution, but it increases rapidly at higher potential, particularly for the O3 material. At 4.4 V, the O3-ZNMT displays more than 10 times more Ni dissolution than the P2-ZNMT, a difference that is not accounted from the initial different Ni content and highlighting the fact that the O1 stacking at the end of charge is promoting this transition metal dissolution.

Finally, the electrolyte recovered after 60h voltage-hold step (using the same protocol as in **Figure 2c**) was analyzed by NMR, where we evidenced the presence of numerous electrolyte degradation products as new peaks in the ¹H NMR spectrum (**Figure 2f**). We assigned the strong signal at 4.3 ppm to sodium ethylene dicarbonate (NaEDC) by comparison with the reported lithium compound observed in lithium-ion batteries originating from EC reduction.[14,15] The high intensity of this peak would confirm the hypothesis that SEI components on Na-ion batteries are more soluble than the equivalent Li ones. Other peaks remain unassigned, and a thorough ¹H NMR study would be needed to fully determine the decomposition products. In any case, no significant differences were observed between the cells using P2 or O3 material, suggesting that the electrolyte decomposition follows similar pathways regardless of composition and structure of the layered oxide cathode.

Nevertheless, the extent of the parasitic reactions does differ between the two systems, with O3-ZMNT|HC full cells displaying higher production of gaseous byproducts and higher transition metal dissolution. We confirm that these parasitic reactions of O3-ZNMT|HC cells are associated with O3-positive electrode material by assembling cells using desodiated Na₃V₂(PO₄)₃ (Na1VP) as the negative electrode, which operates at a higher potential of ~3.2V in comparison to hard carbon (~0.03V vs Na⁺/Na⁰). The results are summarized in Figure S7.

The O3-ZNMT|Na1VP cells degrade even faster than the ones using HC as counter electrode, in comparison to P2-ZNMT|Na1VP cells, demonstrating that the degradation of the cathode material is the main phenomenon behind capacity fade.
## 2.3. Investigating the structural changes of O3 and P2 materials during harsh cycling conditions:

We hypothesize that the differences in electrolyte degradation between the two systems observed in the previous section are due to differences in the structure of the material itself, as the electrolyte used is same in all experiments for both materials. To study the evolution of the structure of both materials upon cycling, we performed *operando* XRD analyses in a home-made cell with Be window and using O3 or P2 materials against sodium metal. The *operando* XRD analyses were carried out either for full charge-discharge cycles between different voltage cut-offs (Figure S8 and S9), or by introducing a constant voltage step for 12h at the end of charge (Figure S10 and S11). Note that the constant voltage step at the end of charge leads to further material oxidation and electrolyte oxidation depending on the oxidation cut-off voltage (Figure S12) and the XRD evolution for the CV step at 4.4V, especially in the first 00l peak is shown in **Figure 3** with the complete data set provided in Figures S10 and S11. The structural evolutions of O3 → P3 → P3-O1 intergrowth and P2 → P2-O2 intergrowth are observed for O3 and P2 materials, respectively, during the CC step, as reported before. However, during the voltage-hold step, new peaks start to appear (mentioned (ii) and (iii) in **Figure 3a** right) at the expense of the initial phase (i), indicating some additional structural modifications of the P3-O1 intergrowth. In turn, the P2-O2 end-phase from the desodiated P2 phase exhibits a small shift in the 002 peak during the CV step, possibly related to the increase in O2 domain formation on continuous charging.
<figure><img src="image_5.png" /></figure>

<figure><img src="image_6.png" /><figcaption>**Figure 3.** *Operando* XRD studies of both materials during the first charge (at C/5 rate), and during a subsequent voltage hold (12h). (a) O3-ZNMT, and (b) P2-ZNMT</figcaption></figure>

To get further insight into these structural changes, we analyzed the O3 and P2 electrodes from the O3-ZNMT|HC and P2-ZNMT|HC cells after a long high-voltage hold (60h at 4.4 V), by EDX and Transmission Electron Microscopy (TEM). The Na content in the materials estimated by EDX deviates significantly from the Na content in the material extracted from the electrochemical data (compare Figure S13, Figure S14 and Table ST1 in the supporting information). This demonstrates that a big proportion of the current recorded during the high-voltage hold relates to parasitic oxidation of the electrolyte or degradation of the material without Na⁺-deintercalation, particularly for O3-ZNMT, from which there is virtually no Na⁺ deintercalation happened during the voltage hold. This material degradation causes severe cracking in both materials, observed in the low magnification HAADF-STEM imaging (Figure S15).
<figure><img src="image_7.png" /><figcaption>(b)</figcaption></figure>

<figure><img src="image_8.png" /><figcaption>(c)</figcaption></figure>

<figure><img src="image_9.png" /><figcaption>(d)</figcaption></figure>

<figure><img src="image_10.png" /></figure>

**Figure 4.** Transmission electron microscopy observations of both materials after 60 hours constant voltage hold at high potential. (a,b) [100] HAADF-STEM images of the ZNMT-O3 and ZNMT-P2 samples with inserted magnified images. Dashed rectangles indicate the position of HAADF intensity profiles in panel d. (c) Scheme of the O2-type, O2-type with a stacking fault and O3-type structures with corresponding close-packed layer sequences. Dashed lines and indicated angles illustrate the tilt of the TM column rows relative to the (00l) plane. (d) HAADF intensity profile across the TM layers in the ZNMT-O3 and ZNMT-P2 samples. Interlayer distances for the ZNMT-O3 and ZNMT-P2 samples are measured. Asterix indicate the HAADF signal from TMs in the sodium layers.

The SAED patterns (Figure S16) indicate predominantly an O1 stacking for O3-ZNMT and O2 stacking for P2-ZNMT after nearly complete oxidation of the materials by the used CV step for 60h at 4.4V. Further analysis by the [100] HAADF-STEM image in **Figure 4a** confirms that the crystal structure of the oxidized O3 sample is O1, as no lateral shift of the dots of the TM columns is observed (**Figure 4a**, inserted magnified image). The case of P2-ZNMT is less trivial, as the [100] HAADF-STEM image of the charged material revealed the O2-type structure but with many stacking faults. The dots of the TM columns are laterally shifted by 1/2 of the interdot distance (**Figure 4b**, inserted magnified image), but the *hchc* (ABCBA) close-packed sequence of the O2 structure is accidentally violated by switching a single “hexagonal” layer to the “cubic” one resulting in the *hcch* (ABCAC) local sequence (**Figure 4c**). This stacking sequence requires the angle between the row of the TM columns
(dashed line in **Figure 4c**) and (00l) plane of 80°, in contrast to the O3-type structure, in which the angle is 70°, or the perfect O2-type structure in which the angle is 90°. The angle measurement from HAADF-STEM image of the P2-ZNMT provides the value of 80.2(2) degrees, supporting the proposed model.

The HAADF-STEM images of both materials show significant TM cation migration to the Na⁺ layer extending deep into the bulk of the crystallite. This can be clearly evidenced in the intensity profiles of the HAADF signal (**Figure 4d**), displaying additional peaks, tagged with *, between the main TM layer peaks. This phenomenon seems to be more pronounced in the O3-ZNMT material than in P2-ZNMT.

In summary, the microscopy analyses show that O3-ZNMT and P2-ZNMT transform into O1 and faulted O2 under extreme conditions, respectively, but that the extent of the migration of transition metal to the Na-layer is relatively greater in O1 than in O2 end-of-charge phases.

## 2.4. DFT calculations

To further investigate the origin of the stability difference between O1-type and O2-type structures, we conducted DFT calculation for these end-phases of the materials under study. To simplify the DFT calculation, we used the supercells with the chemical composition of Na<sub>0.67</sub>Ni<sub>0.33</sub>Mn<sub>0.67</sub>O<sub>2</sub> and Na<sub>0.83</sub>Ni<sub>0.417</sub>Mn<sub>0.583</sub>O<sub>2</sub> for the P2 and O3 initial structures, respectively. The end member O1 structure of Na<sub>x</sub>Ni<sub>0.417</sub>Mn<sub>0.583</sub>O<sub>2</sub> is constructed by removing all the sodium ions in the structure to mimic the end of charge for O3 structure, as illustrated in **Figure 5a**. For the O2 structure we followed a similar procedure with the adequate composition of the P2 material.

Our calculations show that transition metal migration will significantly lower the energy of the O1 phase (**Figure 5a and 5b**), whereas it is not energetically favored in the O2 structure. This aligns with the observed lower extent of transition metal migration to the Na layer in P2-ZNMT. Note that we also considered various possibilities of TM migration, including the kinds of TM migration (Ni or Mn) and different migration sites, with all showing the same trend (Figure S17). This demonstrates that the TM migration is indeed thermodynamically more favored in O1 structures than in O2 ones. This is in line with previous reports that the stronger M-O bond in the O2 phase helps reducing bond breaking and cation migrations and reactivity of the O-ligand with the electrolyte.[16]

We then explored further how the TM migration affects the charge compensation and corresponding gas release behavior. Notably, TM migration in the O1 phase leaves one of its bonded oxygen ions to form a “vacancy-O-vacancy” environment alike previously proposed
“Li-O-Li” configuration and thus creating oxygen 2p non-bonding states along this direction (encircled oxygen ion in **Figure 5a**). The Density of States (DOS) of the “isolated” oxygen was calculated to evaluate how the electronic structure is affected by TM migration. By integrating DOS, we observed a larger area above the Fermi level after TM migration (**Figure 5c**), suggesting more O 2p hole states are created during charging after TM migration. Besides, Bader charge analysis also reveals more charge loss from oxygen when the O3 phase is charged to O1 phase with TM migration, compared with charging of the P2 phase to O2 phase (Supporting Information Table ST2).

These results suggest that TM migration can generate structures similar to cation disorder, raising the oxygen 2p non-bonding states and activating more oxygen redox. Consequently, the calculated energy for oxygen vacancy formation is lower in the O1 phase than in O2 phase (detailed calculation method is shown in Supporting Note S2), implying more favored oxygen de-coordination or oxygen release. This would explain our experimental observation of more CO2 release observed in the O1 phase by OEMS, which is the result of chemical oxidation of the electrolyte by reactive oxygen.[17] Further studies are required to identify quantitatively the extent of oxygen anion participation in the redox process for O3 vs P2 materials as previous studies have reported the occurrence of anionic redox at high voltages in many different sodium layered oxides. Altogether, our experimental and theoretical studies confirm that O3-ZNMT is less stable at high states of charge due to the O1 phase formation that is associated with transition metal migration and anionic redox, hence leading to higher reactivity with the electrolytes. Additionally, some suggested electrolyte decomposition pathways result in the formation of HF, leading to the observed transition metal dissolution.[10,18]
<figure><img src="image_11.png" /><figcaption>**Figure 5.** (a) Structure of desodiated O3-Na<sub>x</sub>MO<sub>2</sub> phase (O1 structure) without TM migration (up) and with TM migration (down). (b) Free energy comparison before and after TM migration of O1 and O2 structures, respectively. Note that the energy values shown in this figure were taken in the sense of demonstrating the largest thermodynamic driving force for TM migration, whereas other values of possible scenarios are shown in Figure S17. (c) Density of states (DOS) of the isolated oxygen in O1 phase before and after TM migration. (d) Average charge of oxygen and oxygen vacancy formation energy of the O3, O1, P2, and O2 phase, respectively.</figcaption></figure>

**2.5. Development and application of the O3 and P2 materials in 18650 cells**

Full Na-ion cells based on O3-ZNMT show poor structural stability and reduced cycle life in comparison to the ones employing P2-ZNMT material, while it is the most attractive for developing high-energy cells. Thus, in this section we investigate the possibility of improving the cycle life of layered oxide|HC cells by first utilizing the electrolyte formulation containing electrolyte additives that help to produce stable interphase and second by limiting the high-voltage cutoff.
<figure><img src="image_12.png" /><figcaption>**Figure 6.** Cycling performance of O3-ZNMT|HC and P2-ZNMT|HC full cells in 18650 format. (left) Capacity retention plots (both charge and discharge capacities that nearly superimpose on each other with coulombic efficiency of ~99.5 % for O3|HC and ~99.9 % for P2|HC) when the cells were cycled to different cut-off potentials. (right) Derivative plots of cycle 1 (C1) after formation cycle and cycle 200 (C200). The formation cycle consist of a charge-discharge cycle at 25 °C between 1-4.2V for O3-ZNMT|HC and 2-4.2 V for P2-ZNMT|HC at C/10 (1C= 212 mAh/g for O3-ZNMT|HC and 175 mAh/g for P2-ZNMT|HC).</figcaption></figure>

Towards this purpose, we assembled cylindrical 18650 cells using layered oxide and HC as positive and negative electrodes, respectively. The slurry preparation and electrode coating were done in ambient air due to practical limitations. It should be noted that the O3- and P2-ZNMT materials used in this study are relatively more stable in humid air than the Zn,Ti-unsubstituted materials, facilitating electrode processing in ambient air with limited exposure, as no slurry agglomeration was observed during slurry preparation.[3]

After formation, the cells were cycled between 2.0 to 4.1, 4.15, or 4.2V at C/3 (Figure 6). The O3-ZNMT|HC cells showed relatively poor capacity retention in all studied cycling windows, though a comparatively better cycling stability is observed by limiting the oxidation potential to 4.1V (~80% retention after 500 cycles). The derivative plots (dQ/dV) in the right panels indicate the loss of the high voltage plateau as the major contributor to this capacity fade. The formation of the O1 phase (O1 domains in O1-P3 intergrowths) and associated parasitic reactions with electrolyte are the main source of cycle life deterioration, as described in our previous experiments. Although the included electrolyte additives seem to be beneficial, they
are not preventing the interface deterioration completely, and thus further electrolyte optimization is needed. In turn, P2-ZNMT|HC cells showed better cycling stability, especially when the cycling window was limited to 2.0-4.1V and showed ~95% capacity retention even after 1000 cycles (>200 days of cycling).

### 3. Discussion

The O3 and P2 layered oxides are often compared because of the former's better capacity (due to the high sodium content in the starting phase) and the latter's better rate (due to facile Na-Na migration path; refer Figure S18). In addition, the Na-poor P2 starting phase also helps to improve the moisture stability of the pristine material compared to the high-Na O3 phases. In addition to these known advantages of P2 over O3, we have presented here evidence of lower interfacial reactivity of the P2 phase upon desodiation, particularly at high state-of-charge (> 4V; ~80% SOC). Though our experiments and analyses were carried out in the extremely harsh conditions of charging to 4.4 V followed by a constant voltage step for 60 hours, similar structural and interfacial degradations are expected upon repeated charge-discharge cycles as is confirmed by the cycling experiments in 18650 cells.

Such observations, though new in sodium layered oxides, have been the subject of discussion in the context of lithium layered oxides, where a good cycling stability is reported often with the O2 phases. O3 layered oxides are known to readily convert to spinel and disordered rock-salt phases due to their similar cubic close packed oxygen sub-lattice. O2 materials on the other hand, exhibits greater stability over time and in cycling as this transformation presents a higher energy barrier, as we also observed in this study.[19] Previous studies also suggest that the O2-phase stability relates to its lower tendency to evolve oxygen, even in materials displaying a significant degree of anionic redox.[20] The widespread application of O2-type layered oxides in Li-ion technology remains limited, since they can only be prepared by anion exchange from the equivalent P2-type sodium material, a process that entails higher production costs, and can lead to significant stacking faults if not carried out properly.[21] In contrast, Na-based layered oxides can be prepared directly in the O3 or P2 stacking, offering an advantage for optimizing materials for commercial Na-ion batteries.

### 4. Conclusion and perspectives

High sodium content O3-type oxides are widely regarded as desirable cathode materials over Na-deficient P2-type. Our observations indicate, however, that O3-type materials exhibit poor cycling stability arising from a structural instability at high state-of-charge, resulting in a
higher degree of transition metal migration, and uncontrolled electrolyte oxidation. We observe that these two phenomena (transition metal migration and electrolyte oxidation) are interconnected and are set in motion particularly when O1 starts forming, when the oxygen atoms in the material are more readily accepted by solvent molecules in the electrolyte (chemical oxidation). Nevertheless, the exact underlying mechanism deserves further attention. Future research efforts can be targeted to expand our understanding of this matter by leveraging the capabilities of advanced characterization techniques, for example.

Given the superior structural stability of P2 materials, focusing on increasing their capacity—such as through the introduction of suitable sacrificial salts or formulating O3-P2 intergrowths with preferably P2 on the surface—could be the path forward. Alternatively, the O3-type material could be protected with inert surface coatings, a strategy widely employed currently for high-voltage lithium-ion cathode materials.

One could also envision the development of electrolyte formulations less prone to nucleophilic attack by the oxygen in the material. Novel electrolyte formulations, such as high-concentration or locally high-concentration electrolytes (HCE, LHCE) can be an attractive option, although their price tag limits their application in mass production.

While the path toward the ideal layered oxide cathode material for sodium batteries remains long, we are optimistic that the combined study of material properties, interfacial reactivity, and electrolyte optimization will provide a clearer vision for designing more suitable materials.

# 5. Materials and Methods

## 5.1. Layered oxides synthesis

P2-type Na0.67Ni0.3Zn0.03Mn0.52Ti0.15O2 (P2-ZNMT) and O3-type Na0.85Ni0.38Zn0.04Mn0.48Ti0.1O2 (O3-ZNMT) were prepared by traditional solid-state synthesis. Na2CO3, NiO, Mn2O3, ZnO, and TiO2 (all purchased from Sigma-Aldrich) were mixed in stoichiometric ratios and ball-milled for 1 hour using a SPEX ball mill, with a powder-to-ball ratio of 1:10. The mixed powders were calcined at 900 °C for 12 hours, followed by intermediate grinding and reheating at 1000 °C for 12 hours. The heating and cooling rates were 3 °C/min and 1 °C/min, respectively. The prepared materials were stored in an Ar-filled glovebox to avoid possible reactions with air or moisture.

## 5.2. One-sided coated electrodes preparation
The one-side coated electrodes were prepared in a dry room having a dew point below -50 °C. The active material, either P2-ZMNT, O3-ZNMT, or hard carbon (HC) was mixed with a C-45 conductive carbon, polyvinylidene fluoride (PVDF) binder and N-methylpyrrolidone as solvent. After mixing for several hours, the slurry was cast on aluminum foils using a doctor blade, let dry in air, calendered, and punched into disks of 13 mm in diameter, and finally vacuum-dried at 80 °C overnight to completely remove the solvent. The obtained P2-ZNMT, O3-ZMNT, and HC electrodes have active material loading of around 17.2 mg/cm², 12.4 mg/cm², and 6.8 mg/cm², respectively.

### 5.3. Electrolyte preparation and cell assembly

NaPF₆ was purchased from Stella Chemifa and used as received. The solvents, ethylene carbonate (EC), propylene carbonate (PC), and dimethyl carbonate (DMC) were obtained from DoDo Chem and dried over activated 3 Å molecular sieves for at least one week before electrolyte preparation. The water content in all the used solvents are <10 ppm, measured by Karlfisher titration. Vinylene carbonate (VC) and tris(trimethylsilyl phosphate) (TMSPi) were purchased from TCI Chemicals, while succinonitrile (SN) was purchased from Sigma-Aldrich, and the three were used as electrolyte additives without further purification. Sodium oxalatodifluoro borate (NaODFB) was prepared as previously described.[10,22]

Coin cells (type 2032) were assembled in an Ar-filed glovebox using the previously dried 13 mm (cathode) and 14 mm (anode) diameter electrodes and two glass fiber disks as separators soaked with ~150 µL of electrolyte.

Cylindrical 18650 cells were assembled using double-side coated positive and negative electrodes with an N/P balancing of 1.1 (by capacity). Celgard is used as a separator and the electrolyte was a 1 M NaPF₆ solution in EC:PC:DMC (1: 1: 2 by volume) containing four electrolyte additives, VC (3 wt.%), TMSPi (0.5 wt.%), SN (3 wt.%), and NaODFB (0.5 wt.%), following our previously optimized electrolyte composition.[23,24] A formation cycle was done for all cells at 25 °C between 1-4.2V for O3-ZNMT|HC and 2-4.2 V for P2-ZNMT|HC at C/10 (1C= 212 mAh/g for O3-ZNMT|HC and 175 mAh/g for P2-ZNMT|HC). Following the formation, the cells were cycled continuously at C/3 between 2 to 4.1, 4.15, or 4.2V. All electrochemical analyses were carried out using Biologic (France) battery cyclers or potentiostat/galvanostat.

### 5.4. XRD
X-ray powder diffraction (XRD) patterns were collected using a BRUKER D8 Advance diffractometer equipped with a Cu Kα radiation source (λ Kα1 = 1.54056 Å, λ Kα2 = 1.54439 Å) and Lynxeye XE detector from 10 to 70°. Structural evolutions during electrochemical cycling were followed by operando XRD analysis with a homemade operando XRD cell using a Be window as a current collector, as explained elsewhere.[25]

## 5.5. OEMS

Online electrochemical mass spectrometry (OEMS) experiments were conducted using a custom-made cell as described elsewhere.[26] Full cell O3-ZNMT|HC and P2-ZNMT|HC configurations were used for the studies. Towards this purpose, free-standing electrodes of P2-ZNMT, O3-ZNMT, and HC were prepared with a composition of 80 wt.% active material, 10 wt.% carbon Super P (CSP), and 10 wt.% polytetrafluoroethylene (PTFE). The components were mixed in a mortar with few drops of dimethyl carbonate (DMC) to form a homogeneous paste, which was then spread onto a glass microscope slide, detached from the slide and dried under vacuum in the glovebox antechamber for 15 minutes. The electrode was cut to obtain the desired mass and size for having electrodes of n/p ratio 1.1. The cells were assembled using glassfiber (GF/A) separator and 120 µL of electrolyte. The full-cells were cycled between 2 and 4.3 V, with a 4-hour open-circuit rest in between. A galvanostatic cycling with current density of 6.43 mA/g and 18 mA/g was used for P2-ZNMT|HC and O3-ZNMT|HC cells, respectively. The ion currents of the mass-to-charge (m/z) channels of 2, 26, 32, and 44 were recorded and converted to the evolution rates of H₂, C₂H₄, O₂, and CO₂, respectively. At least two cells were repeated to ensure the reproducibility of the results.

## 5.6. ICP-MS

The transition metal (TM) dissolution from the cathode material was measured from P2-ZNMT|HC or O3-ZNMT|HC coin cells after 60 hours hold at either 4.0, 4.2, or 4.4 V at 55 °C. After the holding period, the cells were disassembled and the HC electrode and the separator touching the HC electrode were recovered. The HC and separator were digested together in 5 mL of 2 wt.% nitric acid overnight (diluted from a commercial 65% HNO₃, Suprapur, Merck). The samples were then filtered with polypropylene 0.2 µm PP syringe filters and 0.5 mL of the filtrated solution were diluted to 10 mL with fresh 2 wt.% nitric acid. The Ni, Mn, Zn and Ti in the samples were quantified from the obtained diluted sample using an inductively coupled plasma mass spectrometer (ICP-MS, Nexion 2000, Perkin Elmer), and normalized by the amount of active material originally used in each coin cell.
As only one of the separators is recovered, the observed values for the TM dissolution is always underestimated, but we observed reproducible results by measuring each data point at least three times and reporting the average and standard deviations (displayed as error bars in Figure 2e).

## 5.7. TEM

Samples for transmission electron microscopy (TEM) were prepared in an Ar-filled glovebox by crushing the crystals with an agate mortar and pestle under dimethyl carbonate (DMC) and depositing a drop of suspension onto a carbon film supported by a copper grid. A vacuum transfer holder was used to transfer the sample from the glovebox to the microscope column without contact with air. High angle annular dark field scanning transmission electron microscopy (HAADF-STEM) images, selected area electron diffraction (SAED) patterns and energy dispersive X-ray spectroscopy (EDX) maps were acquired on an aberration-corrected Titan Themis Z transmission electron microscope at 200 kV equipped with a Super-X system for EDX analysis.

## 5.8. DFT

Density functional theory (DFT) simulations were performed on The Vienna Ab initio Simulation Package (VASP).[27-29] To simplify the model, $Na_{0.67}Ni_{0.33}Mn_{0.67}O_2$ and $Na_{0.83}Ni_{0.417}Mn_{0.583}O_2$ for P2 and O3 were adopted for simulation with 2×2×1 supercell. Γ-centered k mesh (3×5×2) and (5×5×2) were used for P2 and O3 supercell models respectively. A generalized gradient approximation (GGA) with the Perdew-Burke-Ernzerhof (PBE) pseudopotential was employed to describe the exchange-correlation potential. Spin polarization was considered, and Hubbard U correction was considered in view of the strong localized character of 3d orbitals of transition metals. The Hubbard U values for Mn and Ni were set as 4 and 6.4 eV respectively. The cutoff energy of plane wave expansion is 520 eV. All atoms were relaxed until the forces on each atom were less than 0.03 eV·Å⁻¹. Atomic structure illustration figure was plotted by VESTA.[30]

## Acknowledgements

This work benefited from state aid managed by the National Research Agency (ANR) under France 2030 bearing the reference ANR-22-PEBA-0003 (PEPR- HIPOHYBAT) and through the STORE-EX Labex project ANR-10-LABX-76-01. The DFT study in this work was supported by high-performance computing platform of Peking University. I.A.M. and A.M.A.
are grateful to the Russian Science Foundation for financial support (grant 23-73-30003). The
authors thank TIAMAT's production team for the 18650 cell assembly.

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