Layered P2-NaxMn3/4Ni1/4O2 Cathode Materials For Sodium-Ion Batteries: Synthesis, Electrochemistry and Influence of Ambient Storage

Sodium-ion batteries promise efficient, affordable and sustainable electrical energy storage that avoids critical raw materials such as lithium, cobalt and copper. In this work, a manganese-based, cobalt-free, layered Na_xMn_3/4Ni_1/4O₂ cathode active material for sodium-ion batteries is developed. A synthesis phase diagram was developed by varying the sodium content x and the calcination temperature. The calcination process towards a phase pure P2-Na_2/3Mn_3/4Ni_1/4O₂ material was investigated in detail using in-situ XRD and TGA-DSC-MS. The resulting material was characterized with ICP-OES, XRD and SEM. A stacking fault model to account for anisotropic broadening of (10l) reflexes in XRD is presented and discussed with respect to the synthesis process. In electrochemical half-cells, P2-Na_2/3Mn_3/4Ni_1/4O₂ delivers an attractive initial specific discharge capacity beyond 200 mAh g⁻¹, when cycled between 4.3 and 1.5 V. The structural transformation during cycling was studied using operando XRD to gain deeper insights into the reaction mechanism. The influence of storage under humid conditions on the crystal structure, particle surface and electrochemistry was investigated using model experiments. Due to the broad scope of this work, raw material questions, fundamental investigations and industrially relevant production processes are addressed.

1 Introduction

To combat man-made climate change, many societies are aiming to reduce their greenhouse gas emissions to net zero (UNFCCC, 2022). To reach this goal, electrical energy generation from renewable, green sources such as solar and wind is being quickly expanded (International Renewable Energy Agency, 2021). The intermittent nature of solar and wind requires great capacities of stationary electrical energy storage to stabilize the grid (Dunn et al., 2011). Additionally, the electrification of cars to reduce carbon dioxide emissions from internal combustion engines requires mobile electrical energy storage solutions. For both applications, electrochemical energy storage in batteries is one of the most promising technologies (Gür, 2018).

The continuous development of lithium-ion batteries (LIBs) since their first commercialization in 1991 by Sony has made LIBs today’s most promising technology for electrochemical energy storage in mobile devices and grid scale applications (Armand et al., 2020). However, lithium-ion batteries are incorporating critical raw materials in view of availability and economic importance, such as cobalt, lithium, copper and graphite (International Energy Agency, 2021). For the near future, the dependency of LIB production on these critical raw materials is expected to result in material shortages and increasing prices (Boer et al., 2021; Boer, 2022).

Sodium-ion batteries (SIBs) utilize a similar working principle as todays LIBs (Sun, 2020), but enable the replacement of the above-mentioned critical raw materials with highly abundant materials, such as sodium for lithium, hard carbon for graphite, and manganese-based, but cobalt-free oxides as host materials for positive (cathode) electrodes, which offer promising performance (Hasa et al., 2021; Tapia-Ruiz et al., 2021). Additionally, the copper current collector on the anode side can be replaced with the cheaper and lighter aluminum foil (Larcher and Tarascon, 2015; Vaalma et al., 2018; Tarascon, 2020). The similar production processes of LIBs and SIBs makes sodium-ion batteries a drop-in technology (Tarascon, 2020). The sum of these aspects is expected to result in reduced prices per energy storage capacity ($/kWh) for SIBs (Vaalma et al., 2018). Due to the higher potential of sodium compared to lithium and the higher ionic radius and mass of sodium ions versus lithium ions, SIBs most likely will not be able to exceed the energy density of the best LIBs (Tapia-Ruiz et al., 2021), but might offer advantages in terms of low temperature performance (Rudola et al., 2021), thermal stability (Kuze et al., 2013) and transportation/storage safety (Bauer et al., 2018). These aspects make SIBs a promising complementary technology to LIBs (Tarascon, 2020).

On the cathode side, layered sodium transition metal oxides and polyanionic compounds such as phosphates and Prussian blue analogues are currently considered in literature (Hasa et al., 2021). Among these materials, layered sodium transition metal oxides are considered among the most promising cathode active materials (CAMs) (Ortiz-Vitoriano et al., 2017; Goikolea et al., 2020; Gonzalo et al., 2021; Tapia-Ruiz et al., 2021) because of similar production processes and processability to LIB CAMs (Drop-In technology), high specific energy and fast solid Na⁺ diffusion (Gonzalo et al., 2021). Following Delmas’ classification for alkali metal layered oxides (Delmas et al., 1980), sodium transition metal oxides are denoted as either O3, P3 or P2 (see Figure 1). In this classification the alphabetic character specifies the alkali metal environment (here O for octahedral, P for prismatic) while the number refers to the MO₂ slabs in the hexagonal unit cell (here 3 or 2). Distorted phases are denoted with a prime (e.g., O’3, P’2). The oxygen stacking in O3-type materials follows an AB CA BC sequence. In P2-type materials, the oxygen stacking is AB BA (Delmas et al., 1980). Note that at room temperature metastable phases such as O2 are accessible via electrochemical (de) sodiation due to gliding of MO₂ slabs. Therefore, electrochemically induced transitions between O- and P-type structures can occur at room temperature (kinetic control). On the contrary, transitions between 2-type (e.g., P2) and 3-type (e.g., P3/O3) structures would require the breaking of metal oxygen bonds, which calls for elevated temperatures (synthesis conditions, thermodynamic control). Therefore, no electrochemically induced transitions between P2 and P3/O3 are possible at room temperature (Raphaële J. Clément et al., 2015). For intercalation type cathode materials such as layered alkali transition metal oxides (A_xMO₂), alkali metal ions (A) are reversibly (de)intercalated from/into the layered host structure, while charge neutrality is maintained by oxidation/reduction of the transition metal. The Fermi energy of the transition metal redox center is directly proportional to the potential of the cathode active material. Therefore, the operating potential of intercalation materials is highly influenced by the active transition metal redox center (Manthiram, 2017). Additionally, structural considerations (Padhi et al., 1997b) and inductive effects (Padhi et al., 1997a) can highly influence the Fermi energy and therefore the cathode potential.

FIGURE 1

FIGURE 1. Crystal structures obtained during synthesis of sodium layered transition metal oxides and their description in Delmas notation (O3, P3 and P2).

O3-type sodium layered oxides typically undergo various phase transitions during electrochemical (de)sodiation (e.g., O3—O’3—P3—P’3) and exhibit slow insertion kinetics at a high degree of sodiation (Gupta et al., 2022). Additionally, these materials are often prone to poor stability in humid atmospheres (Wang et al., 2018; Sun, 2020). P2-type layered oxides on the other hand, typically operate within a broad solid solution range with high sodium-ion (de)insertion kinetics due to the wide diffusion channels of the prismatic sodium environments (Tarascon, 2020). Additionally, in O3-type structures such as nickel-rich NCM materials, a structural degradation from layered to spinel-type to rock salt structures is reported to be a common aging effect (Schipper et al., 2017; Bianchini et al., 2019). For lithium-ion exchanged P2-Na_2/3Li_1/6Mn_5/6O₂, Paulsen and Dahn reported stability against structural conversion to spinel, referring to its specific oxygen stacking sequence (J. M. Paulsen et al., 1999). For P2-type sodium layered oxides, a structural degradation pathway via the spinel structure therefore seems unlikely, due to the similarity of the oxygen stacking sequence. The main drawbacks of P2-type layered oxides is their sodium deficiency resulting in reduced full cell capacities (typically x ∼ 2/3 in Na_xMO₂) (Tapia-Ruiz et al., 2021) and detrimental phase transitions at high degree of desodiation (e.g., P2—O2) (Kubota et al., 2018). In P2-type materials, phase transitions at high degree of desodiation can be mitigated by doping (Gonzalo et al., 2021). Enhanced cycle stability was reported with doping of various elements such as Mg (Singh et al., 2016; Wang et al., 2016; Tapia-Ruiz et al., 2018), Al (Hasa et al., 2017), Ti (Yoshida et al., 2014; Bao et al., 2019), Fe (Hasa et al., 2014; Hasa et al., 2015; Feng et al., 2022), Co (Bao et al., 2017; Bao et al., 2018a; Bao et al., 2018b), Cu (Kubota et al., 2017; Zheng et al., 2017; Yang et al., 2021), Zn (Wu et al., 2016).

In literature, the most investigated P2-type materials are Na_2/3CoO₂, Na_2/3MnO₂ and Na_2/3Ni_1/3Mn_2/3O₂ (Tapia-Ruiz et al., 2021). Among these, P2-Na_2/3Ni_1/3Mn_2/3O₂ incorporates mainly uncritical and low-cost raw materials and offers promising electrochemical performance as cathode active material in sodium batteries. When cycled in sodium half-cells between 2.0—4.5 V, this material offers reversible discharge capacities beyond 160 mAh g⁻¹ (Lu and Dahn, 2001a). Using in-situ XRD, Lu and Dahn observed a solid solution reaction between 2.0—4.0 V and a two-phase reaction at a voltage plateau of approximately 4.2 V. During charge up to 4.0 V within the P2 structure, the a lattice parameter decreased and the c lattice parameter increased. During the two phase reaction at approximately 4.2 V, a second phase, the O2-phase, with a smaller c lattice parameter emerged upon charge and the lattice parameters of the P2-phase remained constant. The authors suggested that the end members of the P2-O2 phase transition are P2-Na_1/3Ni_1/3Mn_2/3O₂ and O2-Na₀Ni_1/3Mn_2/3O₂ (Lu and Dahn, 2001a). During the P2-O2 phase transition, MO₂ slabs glide in the a-b plane in (1/3 2/3 0) or (-1/3 -2/3 0) direction (Kubota and Komaba, 2015), forming octahedral sodium vacancies between the MO₂ slabs. The random choice of gliding direction in the O2 structure causes stacking faults (Lu and Dahn, 2001a). The resulting oxygen stacking in the O2 phase follows the AB CB sequence (Lee et al., 2013). During desodiation of P2-Na_xNi_1/3Mn_2/3O₂, nickel is oxidized to compensate the charge between 2/3 > x > 1/3. For the two phase region at 1/3 > x > 0, anionic redox activity (oxidation of oxygen ions) was reported (Risthaus et al., 2019; Dai et al., 2020; Zhang et al., 2020). A reductive coupling mechanism between nickel and oxygen was proposed as charge compensation mechanism during the P2-O2 phase transition (Dai et al., 2020). At high degree of sodiation x > 2/3 the tetravalent manganese ions are partially reduced to compensate the charge (Risthaus et al., 2019). The voltage profile of P2-Na_xNi_1/3Mn_2/3O₂ exhibits distinct voltage jumps at x = 2/3, 1/2 and 1/3 (Lu and Dahn, 2001a). These voltage jumps were ascribed to sodium vacancy orderings within the alkali metal layers (Lee et al., 2013).

Increasing the manganese content to P2-Na_2/3Mn_3/4Ni_1/4O₂ results in a smoothed voltage profile (Manikandan et al., 2016; Gutierrez et al., 2018), which was related with a significant weakening of sodium/vacancy orderings within the alkali metal layer (Gutierrez et al., 2018). With respect to P2-Na_xNi_1/3Mn_2/3O_2, in-situ XRD on P2-Na_xMn_3/4Ni_1/4O₂ suggested a similar evolution of the crystal structure during electrochemical (de) sodiation at room temperature (Gutierrez et al., 2018). Additionally, a distorted P’2 phase was reported to form at the end of discharge (Gutierrez et al., 2018). To the best of our knowledge, no lattice parameters for the O2 structure have been reported so far for Na_xMn_3/4Ni_1/4O₂.

The specific crystal structure obtained during synthesis of sodium layered oxides depends on synthesis temperature, sodium content and transition metal composition (Paulsen and Dahn, 1999; Lei et al., 2014; Zhao et al., 2020; Xiao et al., 2021). In the Na_2/3Mn_1-yNi_yO₂ system, Paulsen and Dahn investigated the influence of synthesis temperature and Mn/Ni ratio on the crystal structure. They reported P2 structures for high calcination temperatures (> 800°C) and low nickel contents y ≤ 1/3. With decreasing nickel content, higher calcination temperatures are needed to form pure P2 structures (Paulsen and Dahn, 1999). Recently, Xiao et al. published the dependency of the crystal structure on the Mn/Ni ratio and sodium stoichiometry for a fixed calcination temperature of 900°C. Pure P2 structures evolve at nickel stoichiometries y < 0.5 and sodium contents x > 0.67 (Xiao et al., 2021). Access of sodium carbonate during synthesis was reported to enable the formation of the P2-structure at lower calcination temperatures, but can lead to Na₂CO₃ surface impurities (Sathiya et al., 2017).

When considering a rapid scale-up of sodium-ion batteries as a drop-in technology, the stability of sodium-ion cathode active materials in ambient atmospheres is of great interest (Kubota and Komaba, 2015). Lu and Dahn investigated the structural influence of water on P2-Na_2/3Co_zNi_1/3-zMn_2/3O₂ (z = 0, 1/6, 1/3). For cobalt substituted samples, additional reflexes in the XRD pattern corresponding to co-intercalated water were reported. In contrast to cobalt substituted samples, no water co-intercalation was reported for P2-Na_2/3Ni_1/3Mn_2/3O₂ (Lu and Dahn, 2001b). By comparing various P2-type materials, Zuo et al. proposed a reaction mechanism for the P2-type materials in contact with different humid atmospheres: In humid, CO₂-free atmospheres, P2-type layered sodium oxides undergo a sodium/proton exchange with a formation of NaOH on the surface (analogous to nickel-rich NCM materials (Shkrob et al., 2017; Pritzl et al., 2019)). With low amounts of CO₂ in the atmosphere, the composition on the particle surface changes to Na₂CO₃. At high CO₂ concentrations in the atmosphere, NaHCO₃ was proposed as the resulting surface impurity (Zuo et al., 2020). These alkaline surface impurities were reported to result in defluorination of the PVdF binder and a subsequent agglomeration of the slurry, which is complicating the electrode production (Kubota and Komaba, 2015).

In this work, we present the synthesis and electrochemical characterization of a P2-Na_xMn_3/4Ni_1/4O₂ cathode active material with spherical morphology for sodium-ion batteries. The chosen synthesis route includes mixing of a spherical hydroxide precursor with NaOH followed by calcination. From various samples, a phase diagram for Na_xMn_3/4Ni_1/4O₂ cathode active materials for sodium-ion batteries was constructed. The calcination process of phase-pure P2-Na_2/3Mn_3/4Ni_1/4O₂ was investigated in-situ with XRD and TGA-MS in order to understand the synthesis process. X-ray diffraction patterns of resulting phase pure P2 materials exhibit anisotropic broadened (10l) peaks. A stacking fault refinement to account for this anisotropic broadening was performed and is discussed in respect to the synthesis path. The prepared phase pure P2-Na_2/3Mn_3/4Ni_1/4O₂ cathode active material provides electrochemically attractive capacities in sodium half-cells. The change of the crystal structure during electrochemical (de)sodiation was investigated with operando XRD. The influence of storage in ambient air was investigated using model experiments under controlled humidity.

2 Experimental

2.1 Materials, Synthesis and Storage

A dense, spherical Mn_3/4Ni_1/4(OH)₂ precursor was prepared by co-precipitation under vigorous stirring in a continuous stirred tank reactor (CSTR, V = 1 L). Feed-in solutions were prepared from deaerated distilled water, Mn(NO₃)₂ ∙ 4 H₂O and Ni(NO₃)₂ ∙ 6 H₂O (Carl Roth), NaOH (Carl Roth) and NH₃ solution (Carl Roth). The reaction product was filtrated and washed to remove residual salt solution from the product and subsequently dried to obtain the Mn_3/4Ni_1/4(OH)₂ powder.

To construct the phase diagram in Section 3.1, varied amounts of NaOH (Carl Roth) were added as aqueous solution to the hydroxide precursor in a wet impregnation procedure. The wet impregnation process enables a homogenous distribution of the sodium source on the precursor and circumvents experimental uncertainties caused by the hydroscopic nature of NaOH. The dried, impregnated powder was calcined at 450°C in air for 5 h in a Nabertherm batch furnace and subsequently homogenized. Following, heat treatments in air at various temperatures (550–900 °C) were performed with the same Nabertherm batch furnace for 10 h. After calcination, the material was allowed to naturally cool down.

The selected material for a detailed physical and electrochemical characterization (Section 3.2 and Section 3.3) was prepared as follows: The amount of sodium hydroxide added during the wet impregnation was adjusted to result in the ratio ϕ = Na/M = 0.68 to account for a slight loss of sodium during high temperature calcination. A pre-calcination at 350°C for 5 h in air was performed in a Nabertherm bath furnace. After homogenization, the material was calcined at 900°C for 10 h in air in the same furnace. The sample was allowed to cool down naturally to 200°C, at which point the material was directly transferred to a Büchi glass oven, where it was kept overnight at 200°C and dynamic vacuum (<5 ∙ 10⁻² hPa). Following, the dried powder was directly transferred into an Ar filled glovebox (MBraun, H₂O < 0.1 ppm, O₂ < 0.1 ppm). Electrode preparation and cell assembly was done in the same glovebox without any contact to the ambient.

For the model storage experiment (Section 3.4), portions of the same material batch were taken from the glovebox and stored over a saturated Ca(NO₃)₂ ∙ 4 H₂O (Sigma Aldrich) solution in a closed container. This closed container was placed in a temperature chamber set to 30°C. Both, humidity and temperature were continuously monitored with a Testo 184H1 USB data logger. After 1, 3, 7, and 14 days of storage, samples were taken and characterized with XRD, Raman, IR, SEM and electrochemical cycling in half-cell configuration.

2.2 Physical and Chemical Characterization

Chemical analysis of the powder samples was done by ICP-OES measurements (Spectro Arcos SOP) in a diluted aqua regia solution. Powder X-ray diffraction (XRD) was performed on a D8Advance (Bruker) in Bragg-Brentano geometry equipped with a Cu X-ray tube and a LynxeyeXE detector. Phase quantification of the prepared samples was performed via Rietveld Refinement with Topas V6 Software. Stacking Fault Refinement was carried out using Faults Software (Casas-Cabanas et al., 2016). For Faults software, the instrumental broadening of the utilized XRD setup was derived from measurements of silicon standard NIST640c. Capillary XRD measurements were performed for ex-situ investigation of desodiated electrodes to avoid any uptake of water (Kubota and Komaba, 2015). In an Ar filled glovebox, desodiated electrodes were recovered from CR2032 coin cells and thoroughly washed in DMC (Carl Roth). The electrode composite was removed from the aluminum current collector foil and filled in glass capillaries, which were sealed inside the glovebox. Capillary XRD measurements were performed on the same lab X-ray diffractometer (Bruker). Instrumental broadening was derived from measurements of NIST660c and NIST640c standards. Evaluation of the obtained patterns was performed using Rietveld Refinement in TOPAS V6 software. All structures in this publication were drawn with VESTA Software (Momma and Izumi, 2011). Non-ambient in-situ XRD was performed on a XPert-Pro diffractometer (Malvern Panalytical) with a Cu X-ray tube and a HTK-1200N non-ambient chamber (Anton Paar). The NaOH impregnated hydroxide precursor was heated in air to mimic the calcination process. The heating was performed at 5°C min⁻¹ with 10 min hold time at the target temperature before the diffraction pattern was recorded. The applied temperature profile is presented with the diffraction data in Section 3.1. TGA-DSC-MS was used to gain further insights on the chemical reactions during the calcination process. A TGA-DSC (Netzsch STA 449C) coupled with a mass spectrometer (Netzsch QMS 403) was used for this purpose. The sample was placed in an Al₂O₃ crucible and heated at 10°C min⁻¹ in a stream of air. To account for the influence of the crucible and air stream, the DSC measurement data was corrected using a blind measurement under the same conditions. The particle architecture and morphology was characterized with a Leo 1530 VP (Zeiss) scanning electron microscope (SEM). Images were generated with an Everhart-Thornley-SE-Detector and 5 kV acceleration voltage. The volume-based particle size distribution was determined with a Mastersizer Micro (Malvern Instruments) by laser diffraction. Raman spectra were collected using a 532 nm laser at a power of 2 W and a Rayleigh-filter on a confocal Raman microscope Senterra (Bruker). Attenuated total reflection Fourier transformed infrared spectroscopy (ATR-FTIR) was performed on an ALPHA Spectrometer (Bruker) with an ALPHA-P (Bruker) ATR unit. All NaMNO spectra were normalized with respect to their highest intensity. Na₂CO₃ (Sigma Aldrich), NaHCO₃ (Fluka) and NaOH (Sigma Aldrich) were used as reference materials.

2.3 Electrochemical Characterization

Electrodes were prepared by mixing the cathode active material, SuperP-Li (Timcal) and PVDF binder (Solvay Solef P5130) in a weight ratio of 84:8:8, respectively, dispersed in an adequate amount of N-Methyl-2-pyrrolidone (NMP, Sigma Aldrich). The homogenous slurry was casted using the doctor blade technique on aluminum foil (Korff). After drying, electrodes of 12 mm diameter were punched. For materials stored in ambient conditions (3d, 14d in Section 3.4), the electrode preparation was performed at the ambient. For all other materials, these processes were performed inside an Ar filled glovebox (MBraun, H₂O < 0.1 ppm, O₂ < 0.1 ppm) to avoid any influence of the ambient. Electrodes were dried in a Büchi glass oven at 130°C and dynamic vacuum (<5 ∙ 10⁻² hPa) overnight and directly transferred in the same Ar filled glovebox. CR2032 coin cells with an aluminum oxide protected positive pole (Hohsen) were prepared with the dried electrode, two layers of glass fiber separator (Whatman GFA), 150 µl 1M NaPF₆ (Fluorochem United Kingdom) in PC (BASF) electrolyte and a sodium metal (Acros Organics) counter electrode. The typical active material loading was ∼4 mg cm⁻². A multichannel galvanostatic cell test system (CTS, Basytec) was used for the electrochemical cycling of the CR2032 coin cells. Operando XRD measurements were performed on a Stadi P (Stoe) diffractometer in transmission geometry. A curved Ge (111) crystal and a silver X-ray tube at 40 kV and 40 mA were utilized to obtain monochromatic Ag_Kα1 (λ = 0.55941 Å) radiation. CR2032 coin cells were modified with polymer windows and were loaded with a cathode electrode (active material: SuperP: PVDF weight ratio 90:5:5; active material loading ∼14 mg cm⁻²), 150 µl electrolyte (1M NaPF₆ in PC + 5% FEC), two layers of glass fiber separator (GFA, Whatman) and a sodium metal counter electrode (Acros Organics). An Interface 1010E (Gamry) potentiostat was used for electrochemical cycling of the cells.

3 Results and Discussion

3.1 Synthesis

A series of samples were prepared by mixing various amounts of NaOH as sodium source with a dense spherical Mn_3/4Ni_1/4(OH)₂ precursor and a subsequent calcination in air as described in the experimental section. The elemental composition of the obtained powders was analyzed using ICP-OES. From the ICP-OES results, a molecular formula following Na_xMn_yNi_1-yO₂ was calculated. Note, that the sodium content of the individual phases might differ, if two or more phases are present in one sample. The crystal structure of the powder samples was characterized by analysis of XRD patterns. The dependence of the obtained crystal structures on the calcination temperature and the sodium content is depicted in a color scheme in Figure 2.

FIGURE 2

FIGURE 2. Phase diagram obtained when varying the calcination temperature and the sodium content. The Mn/Ni ratio was kept constant by using the same Mn_3/4Ni_1/4(OH)₂ precursor. Refined crystal structures are visualized by red dots for P3, blue dots for P2, green dots for P3-P2, dark blue dots for O3-P2 mixtures and purple dots for mixtures of O3-P3-P2. The sodium content (x-axis) was derived from ICP-OES analysis, phase analysis was performed via Rietveld refinement of XRD patterns.

At low temperatures (<700°C), exclusively the P3 layered structure type (red dots) was formed independent of the formal sodium content in the samples (0.45 < Na/M = x < 1.06). At medium temperatures (between 700 and 800°C), a mix of P3 and P2 layered structures (green dots) was identified for low sodium content x and a mix of P2 and O3 layered structures (dark blue dots) was detected for higher sodium content x. At high temperatures (≥ 825°C) and medium to high sodium contents (0.6 < x < 1.06) the material exhibits pure P2-type structure. At high sodium content sodium carbonate impurities are detected in the diffraction patterns (confirmed by ATR-FTIR spectroscopy, not shown here). At low sodium content x, sodium-free manganese-nickel-oxides are present besides the layered oxides (Ilmenite type NiMnO₃ below 700°C and Hausmannite-type Mn₃O₄ and NiMn₂O₃ above 800°C). The obtained crystal structure and impurities only depend on the calcination temperature and the sodium content x. No dependence of the hydroxide particle size was found (different sieving fractions: < 20 µm, 20–25 μm, 25–32 μm, 32–36 μm, > 36 µm). Therefore, we conclude, that the obtained structural composition is not influenced by the length of the transport pathways for sodium into the particles.

The obtained results are in good agreement with the literature for the same transition metal composition Mn/Ni = 3/1. With various sodium contents x > 0.6, Xiao et al. reported exclusively P2 phases to be present at a calcination temperature of 900°C (Xiao et al., 2021). Paulsen and Dahn predicted for a sodium content of x = 2/3, a phase transition from P3 to P2 phases to occur at approximately 830°C (Paulsen and Dahn, 1999). The results presented in Figure 2 are in good agreement with both reported phase diagrams, and provide the additional, previously missing, information for a fixed transition metal ratio Mn/Ni = 3.

Due to the discussed advantages of P2-type materials, we selected phase pure P2-Na_2/3Mn_3/4Ni_1/4O₂ material for the following closer look on the synthesis process and in-depth characterization. To monitor the structural evolution during synthesis of a phase pure P2-type sodium manganese layered oxide, Mn_3/4Ni_1/4(OH)₂ precursor was mixed with NaOH solution to result in a sodium to metal ratio of Φ = Na/M = 2/3 and was calcined at different temperatures. At various stages, ex-situ powder XRD patterns were collected as presented in Figure 3. The XRD patterns of the pristine Mn_3/4Ni_1/4(OH)₂ precursor, the dried powder after impregnation with NaOH solution, the powder after calcination at 450°C in air and after 900°C in air show distinct differences from each other. Already wet impregnation with NaOH solution and drying resulted in a nearly complete transformation of the brucite-type structure of Mn_3/4Ni_1/4(OH)₂ precursor into a poorly crystalline P3-type structure. After calcination at 450°C in air, the transformation to P3 was complete. After calcination at 900°C a phase pure P2 material was obtained. For the Mn_3/4Ni_1/4(OH)₂ precursor, partial oxidation of the Mn²⁺ ions with ambient oxygen is reported in literature (Feitknecht et al., 1962; Owen, 1965; Ghosh, 2020), extracting v electrons per formula unit:

$M{n}_{3/4}N{i}_{1/4}{\left(OH\right)}_2+\frac{v}{4} O_2 \rightleftharpoons M{n}_{3/4}N{i}_{1/4}{O}_2{H}_{2-v}+\frac{v}{2}{H}_{2}O$

FIGURE 3

FIGURE 3. Ex-situ powder XRD patterns of the fresh hydroxide precursor (black), the dry impregnated powder, the powder after calcination at 450°C in air and the phase pure P2 material after 900°C in air.

After wet impregnation with NaOH solution and drying, the interlayer distance is enlarged compared to the pristine Mn_3/4Ni_1/4(OH)₂, which can be rationalized with the incorporation of sodium ions into the layered structure. Note, that the experimentally found interlayer distance differs from hydroxides, double layer hydroxides and Birnessites (Owen, 1965), but is in good agreement with P3-type sodium transition metal oxides. Based on this result, a partial exchange of the protons of the hydroxide precursor for sodium ions is proposed with insertion of u sodium ions per formula unit:

$M{n}_{3/4}N{i}_{1/4}{O}_2{H}_{2-v}+\frac{2}{3} NaOH \rightleftharpoons N{a}_{u}M{n}_{3/4}N{i}_{1/4}{O}_2{H}_{2-u-v}+\frac{u}{2}{H}_{2}O+\left(\frac{2}{3}-u\right) NaOH$

To investigate the structural changes during calcination in more detail, in-situ XRD patterns are collected during the calcination process in air and presented in Figure 4A. Upon heating, the (003)_P3 reflex continuously shifts to lower 2Θ, indicating an increase of the interlayer distance and the c parameter of the P3 structure. Beyond 200°C the crystallinity of the P3 phase increases continuously till 800°C. In this temperature range, the a lattice parameter of the P3 structure is 2.91 Å, 2.89 Å and 2.92 Å and the c lattice parameter of the P3 structure is 17.01 Å, 17.06 Å and 17.05 Å at 250, 450 and 800°C, respectively. Between 825 and 875°C the structure transforms from P3 to P2. No changes were observed for the crystal structure during the 3 hour holding time at 900°C. The P2 structure exhibits lattice parameters of a = 2.93 Å and c = 11.30 Å. During cool down, no further changes were observed in the diffraction pattern except for a thermally induced shrinkage of the unit cell. Rietveld refinements of the obtained diffraction patterns were performed and the obtained information is presented in Figure 4B. Note, that all (10l)_P2 reflexes appear broadened, which will be discussed in the following section. A waterfall diagram of the diffraction patterns is provided in the supporting information Supplementary Figure S1.

FIGURE 4

FIGURE 4. In-situ XRD calcination of Na_2/3Mn_3/4Ni_1/4O₂ in air. (A) The temperature profile is represented in the upper diagram. The obtained diffraction pattern is represented as a heat map in the lower diagrams. (B) Refined weight fractions, change of the unit cell volume, metal-metal distance (a lattice parameter) and interlayer distance (proportional to c lattice parameter).

To gain further insight into the structural transformation process, the chemical reactions during the calcination process were investigated with in-situ TGA-DSC-MS as presented in Figure 5. In the TGA-DSC-MS experiment, a first weight loss can be observed between 80—150°C, which is accompanied by an endothermal reaction and the formation of water (m/z = 18). With a reaction onset of 210°C, a second endothermal signal is observed in the DSC measurement. This reaction leads to the formation of water and carbon dioxide (m/z = 44). Whereas the evolution of water is peaking at a temperature of 300°C and is not observed beyond 420°C (m/z = 18 on baseline level), carbon dioxide is evolving till 900°C. Total mass loss after reaching 900°C accounts for 7%.

FIGURE 5

FIGURE 5. Operando TGA-DSC-MS calcination in air.

Below 200°C, in-situ XRD measurements revealed a shift towards higher interlayer distances of the P3 structure, which could be caused by further sodiation of the P3 structure. In the same temperature range, the TGA-DSC-MS measurement showed an endothermic reaction and a loss of water vapor. The endothermic response is expected for evaporation of water. Therefore, analysis of XRD and TGA-DSC-MS data suggests that residual NaOH may act as a sodium source for further sodiation of a partially sodiated P3 phase with additional w sodium ions per formula unit:

$N{a}_{u}M{n}_{3/4}N{i}_{1/4}{O}_2{H}_{2-u-v}+\left(\frac{2}{3}-u\right)NaOH\rightleftharpoons N{a}_{u+w}M{n}_{3/4}N{i}_{1/4}{O}_2{H}_{2-u-v-w}+\frac{w}{2}{H}_{2}O+\left(\frac{2}{3}-u-w\right)NaOH$

From 200 to 450°C, the in-situ XRD experiment showed an increase of crystallinity of the P3 phase, a shrinkage of the a lattice parameter and increase of the interlayer distance (c lattice parameter) of the P3 structure. The operando TGA-DSC-MS experiment showed evolution of H₂O and CO₂ from 200 to 450°C. As no carbonates have been used for synthesis, CO₂ most likely originates from Na₂CO₃ formed from the reaction of NaOH and CO₂ from ambient air during impregnation and drying:

$2 NaOH+C{O}_2\rightleftharpoons N{a}_{2}C{O}_3+H_{2}O$

$NaOH+C{O}_2\rightleftharpoons NaHC{O}_3$

$2 NaHC{O}_3\rightleftharpoons N{a}_{2}C{O}_3+C{O}_2+H_{2}O$

In fact, Na₂CO₃ and NaHCO₃ is detected with ATR-FTIR on the dried hydroxide precursor after impregnation with NaOH and on the powder after calcination at 450°C in air (Supplementary Figure S2). The evolution of water vapour and carbon dioxide (200–450°C) in the TGA-DSC-MS measurement and the increasing interlayer distance in the in-situ XRD experiment, can be rationalized by further sodiation of the layered structure, with NaOH, NaHCO₃ or Na₂CO₃ as sodium source adding x sodium ions per formula unit:

$N{a}_{u+w}M{n}_{3/4}N{i}_{1/4}{O}_2{H}_{2-u-v-w}+\left(\frac{2}{3}-u-w\right)NaOH\rightleftharpoons N{a}_{u+w+x}M{n}_{3/4}N{i}_{1/4}{O}_2{H}_{2-u-v-w-x}+\frac{x}{2}{H}_{2}O+\left(\frac{2}{3}-u-w-x\right)NaOH$

The shrinkage of the a lattice parameter (metal-metal distance) in Figure 4B, suggests an ongoing oxidation of the transition metals within this temperature range.

$N{a}_{u+w+x}M{n}_{3/4}N{i}_{1/4}{O}_2{H}_{2-u-v-w}+\frac{2-u-v-w-x}{4}{O}_2\rightleftharpoons N{a}_{u+w+x}M{n}_{3/4}N{i}_{1/4}{O}_2+\frac{2-u-v-w-x}{2}{H}_{2}O$

From 450 to 800°C, in-situ XRD showed a further increase in crystallinity of the P3 structure, an increasing a lattice parameter and a constant c lattice parameter. In the TGA-DSC-MS experiment, solely CO₂ evolved above 450°C. These results suggest a further sodiation of the structure with Na₂CO₃ as sodium source. Note, that a further sodiation of the structure requires a reduction of the transition metals, which is in accordance with the increasing a lattice parameter.

$N{a}_{u+w+x}M{n}_{3/4}N{i}_{1/4}{O}_2+\frac{2/3-u-w-x}{2}N{a}_{2}C{O}_3\rightleftharpoons N{a}_{2/3}M{n}_{3/4}N{i}_{1/4}{O}_2+\frac{2/3-u-w-x}{2}C{O}_2+\frac{2/3-u-w-x}{4}{O}_2$

From 825 to 875°C, in-situ XRD showed the structural transition from P3 to P2. In the same temperature range, no significant weight loss and no DSC response occurred in the TGA-DSC-MS, indicating a pure structural transformation in absence of chemical reactions. Note that the observed weight loss of 7% is below the expected weight loss of approximately 11% for the overall reaction:

$M{n}_{3/4}N{i}_{1/4}{\left(OH\right)}_2+\frac{2}{3}NaOH+\frac{1}{3}{O}_2\rightleftharpoons N{a}_{2/3}M{n}_{3/4}N{i}_{1/4}{O}_2+\frac{4}{3}{H}_{2}O$

This experimentally observed discrepancy between the experiment and the theoretical weight loss further supports a partial sodium/proton exchange after wet impregnation and drying of the hydroxide precursor with NaOH solution.

3.2 Physical Characterization

A phase pure P2 material was prepared from the same hydroxide precursor and in similar fashion as the materials reported in the previous section. A detailed description of the material preparation can be found in the experimental section. In this section a detailed physical characterization of the material is presented. Electrochemical characterization as well as a model storage experiment based on the same material is presented in the following sections.

The chemical composition of the material as determined with ICP-OES is Na_0.64Mn_0.74Ni_0.26O₂. Powder X-ray diffraction pattern of this material is presented in Figure 6. All reflexes of the diffraction pattern can be assigned to ICDD PDF 01-070-3726 and described with space group P6₃/mmc (194) with lattice parameters a = 2.883 Å and c = 11.184 Å. Additionally, (10l) peaks are significantly broadened. Such anisotropic broadening was ascribed to stacking faults in the literature and “O3- or P3-type” stacking faults were proposed to account for anisotropic broadening of (10l) peaks (Yabuuchi et al., 2012). To the best of our knowledge, no refinement of the actual structural stacking fault model has been reported so far.

FIGURE 6

FIGURE 6. Refined powder XRD pattern of a phase pure P2 material with stacking faults (left) and schematic representation of the stacking fault sequence (right).

In the ideal P2 structure, the stacking of the transition metals follows a CC arrangement (metals stacked on top of each other) and oxygen is stacked in AB BA sequence (Delmas et al., 1980). By introducing a stacking fault with a (2/3 1/3 1/2) transition vector, the stacking sequence of the transition metal is transformed to CA resulting in broadening of the (10l) peaks (see Figure 6). In this specific layer, sodium exhibits a prismatic environment with shared edges and faces with the transition metal octahedra such as in P3-type structures (oxygen stacking AB BC CA). Note, that other stacking faults can account for a similar anisotropic broadening, but would result in octahedral environment of the sodium ions. Since octahedral environments usually result in smaller c lattice parameters and hence additional reflexes in the diffraction pattern, these stacking faults have been ruled out.

As reported above, the material maintains a P3 structure during calcination up to 800°C. Between 825 and 875°C the structure transforms from P3 to P2. This structural transformation requires a breakage and reassembly of all the transition metal oxygen bonds of every second MO₂ slab. Using structural visualization with polyhedra, the P3 to P2 transformation can be described as a 60° rotation of all metal oxygen octahedra of every second MO₂ slab. The transformation from P3 to P2 most likely starts at random points within one crystallite and proceeds from these points until the growing P2 domains meet. If the oxygen stacking of the meeting P2 domains does not fit to the ideal AB BA oxygen stacking of the P2 structure, stacking faults remain. The refined stacking fault probability of 11.6% suggests an average number of nine MO₂ slabs per P2 domain.

For our material no notable change of the diffraction pattern did occur even after additional 40 h at 900°C (not shown here), which indicates that the stacking faults remain stable. From an energetic point of view, this could be rationalized, as the elimination of these stacking faults would require the breakage of all transition metal oxygen bonds of one of the neighboring P2 domains, which comes with a high energy barrier.

The particle morphology and architecture was characterized with SEM. The obtained images are presented in Figures 7A–B. The material consist of spherical, dense particles with micrometer sized crystals. A particle size distribution of this material is presented in Figure 7C. The powder offers a narrow particle size distribution (D₁₀ = 7.7 µm, D₅₀ = 13.2 µm, D₉₀ = 21.1 µm), well suited for large scale slurry and electrode processing. Due to the chosen synthesis route, the material can easily be scaled up. In fact, the production of 350 g of this material was successful.

FIGURE 7

FIGURE 7. SEM images at (A) low and (B) higher magnification. (C) particle size distribution of P2- Na_0.64Mn_3/4Ni_1/4O₂ material.

3.3 Electrochemical Characterization

After calcination, the material was processed without any further contact with the ambient. Electrochemical characterization of CR2032 coin cells in half-cell configuration is presented in Figure 8. The cells were cycled at C/10 (17.3 mA g⁻¹) for the first three cycles, followed by C/3 (57.8 mA g⁻¹) in the subsequent cycles. Every tenth cycle, a check-up cycle at C/10 (17.3 mA g⁻¹) was performed. For the initial discharge and the subsequent first cycle, the sodium content of the active material was calculated based on the active material specific charge. The voltage profile is presented in Figure 8A. During the initial sodiation (discharge) to 1.5 V, the material is sodiated to x = 0.87. During the desodiation a plateau at approximately 2.0 V arises. Further desodiation results in a smooth voltage profile till a sodium stoichiometry of x = 1/3 is reached, at which point the cell voltage sharply rises from 3.9 to 4.2 V. Further desodiation results in a voltage plateau, indicating a two phase reaction. At a sodium content of x = 0.07, the cell voltage reaches the upper cut off at 4.3 V. During the subsequent re-sodiation, the cell voltage mimics the desodiation process with small hysteresis until x ≈ 0.65. Further galvanostatic sodiation exhibits increased hysteresis between sodiation and desodiation cell voltage. At the end of cycle 1, the calculated sodium stoichiometry is x = 0.84 at 1.5 V, indicating some irreversible losses during the first cycle. At sodium stoichiometries of x = 2/3 and x = 1/2, slight shoulders in the voltage profile are evident. For the same sodium contents, sodium vacancy orderings were reported in the closely related material P2-Na_xNi_1/3Mn_2/3O₂ (Lee et al., 2013). Therefore, shoulders of the voltage profile in Figures 8A,B, might be related to sodium/vacancy ordering. A deeper investigation on sodium/vacancy ordering in P2-Na_xMn_3/4Ni_1/4O₂ is currently in preparation. In Figure 8B, selected charge and discharge curves are presented. During the first cycle, attractive specific discharge capacities of approximately 202 mAh g⁻¹ are reached. These initially high specific discharge capacities fade upon cycling. In cycle 10 and 30, the discharge capacities are 165 mAh g⁻¹ and 161 mAh g⁻¹, respectively. The capacity fade is mainly caused by a progressive capacity loss on the high voltage plateau, combined with a progressive shift to lower discharge potentials. The voltage curves between 1.5 and 4.0 V maintain their original shape. The cycle stability is presented in Figure 8C. The specific discharge capacity fades considerably during the first three cycles at C/10. During the subsequent cycles at C/3 and C/10, it stabilizes with a linear capacity fade at a much lower rate. The couloumbic efficiency of the first three cycles averages to 97.5% (Figure 8D). During the cycling at C/3, the average coulombic efficiency is 99.1%. The deviation from the average coulombic efficiency in cycle 4, 10, 11, 20, 21 and 30 are caused by the cycling protocol: the change of the specific current density between cycles causes the jumps of the coulombic efficiency. Prolonged cycling for 110 cycles with the same cycling protocol is presented in Supplementary Figure S3 of the supporting information. After 100 cycles 50% of the initial capacity remains. The origin of the capacity fade will be addressed in detail in an upcoming publication. In literature similar voltage profiles and comparable specific capacities were reported (Manikandan et al., 2016; Gutierrez et al., 2018). Several reasons for capacity fade were reported for P2-Na_xMn_1-yNi_yO₂ materials such as cracking due to anisotropic volume change (Yoshida et al., 2014; Kubota et al., 2017), exfoliation (Liu et al., 2016; Alvarado et al., 2017) and partial irreversibility of oxygen redox (Dai et al., 2020; Zhang et al., 2020).

FIGURE 8

FIGURE 8. Electrochemical characterization of phase pure P2-Na_0.64Mn_3/4Ni_1/4O₂ in sodium half-cells. (A) potential profile of initial discharge and cycle 1 (B) charge and discharge potentials of cycle 1, cycle 10 and cycle 30 (C) specific discharge capacity of the first 30 cycles (D) coulombic efficiency of the first 30 cycles. In (C,D) points represent an average value of three cycles. The standard deviation of these cells is presented as error bars. Cycles 1, 2, 3, 10, 20, 30 are performed at 17.3 mA g⁻¹ (C/10), all other cycles are performed at 57.8 mA g⁻¹ (C/3).

Operando XRD was performed for the initial sodiation and the first full cycle to gain further insights into the sodium storage mechanism and the evolution of the crystal structure during sodium (de)intercalation. The obtained voltage profiles during the initial discharge (light blue), the first charge (red) and first complete discharge (blue) are presented with the associated diffraction patterns as a heat map in Figure 9 (left). The full diffraction patterns are presented in the supporting information Supplementary Figure S4 as a waterfall diagram. The voltage profile of this operando setup closely matches the voltage profile obtained from regular coin cells as presented in Figure 8. During the initial discharge, the (002) and (004) peaks shifted to higher diffraction angles, indicating a decrease of the c parameter. Subsequent charge to 4.0 V, lead to a shift in opposite direction, indicating an increase of the c parameter. At the 4.2 V voltage plateau, the positions of the (002) and (004) peaks remained constant and the intensity decreased. At the end of charge, an additional broad peak at 2Θ ≈ 7° (d ≈ 4.6 Å) emerged, which can be assigned to (002)_O2 peak of the O2 structure (Lu and Dahn, 2001a). In the subsequent discharge, the (002)_O2 peak disappeared and the (002) and (004) peaks of the P2 structure increased in intensity during the course of the voltage plateau. Through further discharge the (002) and (004) peaks shifted to higher angles, pointing to an increase of the c parameter.

FIGURE 9

FIGURE 9. Operando XRD: (left) measured cell voltage during the operando XRD and heat map of the obtained diffraction patterns. (right) cell voltage for the calculated sodium stoichiometry and refined lattice parameters.

The (100), (102), (103) and (110) peaks shifted to lower angles during the initial discharge, indicating an increase of the a parameter. Through the subsequent charge, these peaks shifted towards higher angles, pointing to a contraction of the a parameter. In the course of the voltage plateau at 4.2 V, the a parameter of the P2 structure remained constant. No additional peaks arising from the O2 structure can be clearly detected from the operando XRD, most likely due to severe broadening of the peaks caused by natural occurrence of stacking faults in the O2 structure (Lu and Dahn, 2001a). From our operando XRD data, no indication for a Jahn-Teller distorted P’2 structure (orthorhombic or monoclinic) (Gutierrez et al., 2018) can be derived.

Figure 9 (right) shows the correlation between cell voltage and refined lattice parameters with the calculated sodium content as x-axis. For each sodium stoichiometry, the lattice parameters are in good agreement and independent from the applied current direction, indicating structural reversibility. With decreasing sodium content, the a lattice parameter shows a nearly linear decrease between 0.84 < x < 0.35 and remains constant at sodium stoichiometries x < 0.35. The c lattice parameter of the P2 phase shows a continuous increase with decreasing sodium content between 0.84 < x < 0.35 and remains constant at sodium stoichiometries x < 0.35. Refined values of the c lattice parameter of the O2 phase emerging at the end of charge, are significantly smaller than the c lattice parameters of the P2 phase. A significant mechanical stress on the crystallites is expected due to this lattice collapse (Yoshida et al., 2014; Kubota and Komaba, 2015). Our refinements result in a linear increase of the c lattice parameters of the O2 structure during desodiation.

A solid solution of the sodium ions within the layered host structure is strongly suggested in the voltage window of 1.5–4.0 V, as phase analysis clearly shows a single phase (P2) with gradually changing lattice parameters and a smooth voltage profile in the electrochemical measurement (Robert A. Huggins, 2009). Within the solid solution regime, the course of the a lattice parameter can be rationalized with the radii of the active metal ions, which are reversibly oxidized during desodiation to compensate the charge (Mn³⁺ ∼ 0.65 Å, Mn⁴⁺ ∼ 0.53 Å, Ni²⁺ ∼ 0.69 Å, Ni³⁺_(LS) ∼ 0.56 Å, Ni⁴⁺ ∼ 0.48 Å) (Kubota et al., 2018). The characteristic change of the c lattice parameter within the solid solution regime can be rationalized with sodium ions shielding the repulsive forces of oxygen ions of neighboring MO₂ slabs. During sodiation, additional sodium ions are intercalated into the structure, increasing the positive charge stored between the oxygen ions and therefore shielding the repulsion in between neighboring MO₂ slabs. As a result, the interlayer distance (d = c/2) decreases during sodiation (discharge). During desodiation (charge), sodium ions are removed from the structure, resulting in increasing repulsive forces between the oxygen ions and an increase of the interlayer distance. In the course of the 4.2 V voltage plateau, two phases with constant lattice parameters and inverse course of intensity are simultaneously detected (P2 and O2), proving a two-phase reaction at the 4.2 V voltage plateau. For the O2 phase, the increasing c lattice parameters during desodiation are unanticipated. For a pure two-phase reaction, constant lattice parameters are to be expected (Robert A. Huggins, 2009). The change of the c lattice parameter of the O2 structure, might originate from a narrow solubility range of sodium in the O2 structure or a reductive coupling mechanism between nickel and oxygen as proposed for Na_xNi_1/3Mn_2/3O₂ at low degree of sodiation (Dai et al., 2020). Overall, our operando XRD results are in good accordance with previous reports (Manikandan et al., 2016; Gutierrez et al., 2018). Based on literature and stoichiometric considerations, the formal active redox couples are Mn³⁺/Mn⁴⁺ for x < 1/2, Ni²⁺/Ni³⁺ for 1/2 < x < 1/3 and Ni³⁺/Ni^3+δ or O^2-/Oⁿ⁻ (n = 2–δ) for x < 1/3 (Bao et al., 2017; Risthaus et al., 2019; Dai et al., 2020; Zhang et al., 2020).

Ex-situ characterization to gain further insights into the crystal structure of the O2 phase was performed. The obtained diffraction pattern for an electrode charged to a formal sodium stoichiometry of x = 1/6 is presented in Figure 10. In the obtained diffraction pattern, the sharpest peaks can be assigned to the P2 structure. Different from operando measurement, two O2 phases were detected. Both O2 phases differ in c lattice parameter. In the refinement, the following four structure models were included: a P2 structure with a sodium content of x = 1/3 (approximate onset of the two-phase reaction), an O2 structure with a c lattice parameter of 9.417 Å, a second O2 structure with a smaller c lattice parameter of 8.972 Å and aluminum (originating from the current collector foil). The c lattice parameter of the two O2 structures correspond to the endpoints of the range obtained from operando XRD. The (002) and (004) peaks of both O2 structures are clearly evident, while all other peaks are severely broadened. In contrast, only one O2 structure with an intermediate c lattice parameter was detected at the same sodium stoichiometry in the operando XRD. In the operando XRD the cell is under galvanostatic cycling conditions and therefore a certain kinetic control must be considered. For the ex-situ measurement a thermodynamic equilibrium can be assumed. The difference might arise from a certain solubility range of sodium in the O2 structure or a reductive coupling mechanism between nickel and oxygen as reported for Na_xNi_1/3Mn_2/3O₂ (Dai et al., 2020). To gain a full picture, further investigations on the processes occurring at the 4.2 V voltage plateau in Na_xMn_3/4Ni_1/4O₂ are required.

FIGURE 10

FIGURE 10. Refinement of ex-situ XRD pattern of desodiated electrode at x = 1/6 obtained in capillary transmission. Measured data as red crosses, refined pattern as black line, difference as grey line.

3.4 Effect of Storage

To evaluate the stability of our spherical P2-Na_0.64Mn_3/4Ni_1/4O₂ material in ambient atmospheres, a model storage experiment was performed. A schematic illustration of the experimental setup is presented in Figure 11A. The material was stored at controlled temperature in a closed container over a saturated Ca(NO₃)₂ solution. The deliquescence relative humidity (DRH) of Ca(NO₃)₂ ∙ 4 H₂O at 30°C was reported to be 46.7% (Adams and Merz, 1929; Guo et al., 2019). Using a humidity and temperature tracker, we monitored the storage conditions. The average temperature was 27.6°C and the relative humidity was 48.3% with high accuracy over the course of the 14 days storage experiment. Samples are taken after 1, 3, 7 and 14 days of storage.

FIGURE 11

FIGURE 11. (A) Schematic representation of model storage experiment, (B) XRD patterns of pristine and stored materials, (C) representative Raman spectra of pristine and stored materials, (D) ATR-FTIR spectra of pristine, stored and reference materials and (E) zoom-in into ATR-FTIR spectra.

To check the bulk stability over the course of the model storage experiment, powder X-ray diffraction was performed. The obtained diffraction patterns are presented in Figure 11B. No additional peaks are evident from the diffraction patterns, indicating a structural stability under applied conditions. Refined lattice parameters and stacking faults are reported in Table 1. No significant change of the a lattice parameter and the stacking fault probability is observed, but the c lattice parameter decreases during the storage experiment. The change in c lattice parameter can be rationalized with a sodium/proton exchange and the formation of amorphous surface NaOH, NaHCO₃ or Na₂CO₃ due to the reaction with absorbed water and the ambient atmosphere (Zuo et al., 2020). To obtain further information on the local structure, confocal Raman spectroscopy was performed on the stored materials and representative spectra are presented in Figure 11C. The bands with Raman shifts between 200 cm⁻¹ and 700 cm⁻¹ can be assigned to A_1g, E_1g and E_2g modes of the P2-type material (Iliev et al., 2004). To the best of our knowledge, so far the band assignments for P2-type sodium transition metal layered oxides in literature are not consistent (Iliev et al., 2004; Lemmens et al., 2004; Qu et al., 2006; Yang et al., 2006; Singh et al., 2015; Wang et al., 2016). From the obtained Raman spectra, no change of the local crystal structure is evident over the course of the model storage experiment. For all stored materials, sharp Raman bands around 1080 cm⁻¹ corresponding to the υ₁ symmetric stretching vibration of the carbonate group in Na₂CO₃ (Buzgar and Ionut Apopei, 2009) were evident in certain local areas, while in other areas no sodium carbonate bands could be detected. This suggests locally formed sodium carbonate impurities on the particles. ATR-FTIR was performed to allow a semi-quantitative comparison. With ATR-FTIR many particles are simultaneously probed, providing statistically averaged spectra. The obtained spectra of stored materials with reference materials are presented in Figures 11D,E. Sodium carbonate is the dominant surface species for all stored materials. The intensity of the Na₂CO₃ bands clearly increases with storage time.

TABLE 1

TABLE 1. Lattice parameters and stacking fault probabilities derived from refinement of X-ray diffraction patterns of pristine and stored materials.

SEM images of the pristine, 3 days stored and 14 days stored material are shown in Figure 12. As presented above, the pristine material exhibited a clean surface with micrometre sized crystals. After 3 days of storage, the surface of the material was covered with a surface film and small dots appeared on the surface. After 14 days of storage, the surface of the material resembled the pristine one, but needle like structures were attached to the spherical particles. Such needle like structures on sodium layered oxides were previously reported for materials stored in humid air (Duffort et al., 2015; Zuo et al., 2020). Based on EDX and XRD, these needles are reported to consist of NaHCO₃ or Na₂CO₃ type compounds (Duffort et al., 2015; Zuo et al., 2020).

FIGURE 12

FIGURE 12. SEM images of (A) pristine, (B) 3 days stored material and (C) 14 days stored material.

Characterization results on a similar material (P2-Na_0.69Mn_3/4Ni_1/4O₂) stored at higher relative humidity in a similar setup (60–68% RH, approx. 30°C) are presented in the supporting information Supplementary Figure S5. The bulk structure remains intact and no water co-intercalation is observed in XRD patterns (Supplementary Figure S5A) independent of the high relative humidity. No change of the local structure is apparent from Raman spectra (Supplementary Figure S5B). The dominant surface species is Na₂CO₃ as determined with ATR-FTIR (Supplementary Figures S5C,D). The amount of adsorbed water is determined with Karl-Fischer titration (Supplementary Figure S5E). Up to 0.45 %wt. of water is absorbed during the first days of storage. With prolonged storage, the amount of absorbed water decreases slightly and the amount of surface Na₂CO₃ increases. Overall, results obtained during storage at 60–68% relative humidity (Supplementary Figure S5) agree well with the above-described results at 48% RH (Figure 11).

Our results strongly suggest the uptake of water with subsequent sodium/proton exchange and the formation of Na₂CO₃ on the particle surface as a result of humid ambient storage. These results are in good accordance with the proposed reaction mechanism in literature (Zuo et al., 2020). The formed surface films were not uniformly distributed on the particle surface. Upon prolonged storage, needle like sodium carbonate compounds formed on the particle surface.

To evaluate the influence of storage on the electrochemical performance of the material, coin cells were prepared from the material stored for 3 and 14 days. The electrochemical results of these cells are compared with the pristine material in Figure 13. Voltage curves for the initial sodiation and the first cycle are presented in Figure 13A. The OCV of the pristine, 3 days stored and 14 days stored material was 2.48, 2.68 and 2.69 V, respectively. Higher charge uptake during the initial galvanostatic sodiation was observed for the stored materials compared to the pristine material. For the first cycle, the voltage profiles of all three materials are very similar. The cycling stability is presented in Figure 13B. During the first cycle all three materials offered similar specific discharge capacities. In the subsequent cycles, both stored materials exhibited slightly less capacity fading, resulting in higher discharge capacities for the stored material. The coulombic efficiency, as presented in Figure 13C, is very similar for all three materials.

FIGURE 13

FIGURE 13. (A) charge and discharge potential profiles of cells with pristine, 3 and 14 days stored material, (B) corresponding specific discharge capacity for the first 30 cycles and (C) coulombic efficiency of the first 30 cycles. For (B,C) every point represents an average of three cells, the standard deviation is presented as error bars.

Overall, the electrochemical cycling of the materials in coin cells seems not affected by the prolonged storage in humid ambient conditions. But the increased OCV and the higher capacity during the initial discharge, indicates a loss of active sodium due to storage. In a full cell with an initially sodium-free anode, a loss of active sodium directly translates into reduced capacity and reduced energy density of the full cells. Additionally, a negative effect of the formed alkaline surface films on the processability in PVdF based slurries is anticipated (Kubota and Komaba, 2015) and a more pronounced gassing in full cells can be expected as reported for lithium-ion batteries (Zhang, 2014; Jung et al., 2018; Sicklinger et al., 2019; Teichert et al., 2020). From a process point of view dry room conditions could be used to avoid the reaction with humid atmosphere, however this is associated with higher cost for the processing. From a materials point of view, a stabilisation via compositional variation (Zuo et al., 2020) or application of protective surface coatings allow direct processing of sodium ion cathode materials in standard large scale production.

4 Conclusion

In this work, we report a study on the synthesis, electrochemistry and ambient storage of Na_xMn_3/4Ni_1/4O₂ layered oxides as cathode materials in sodium-ion batteries. Analogue to large scale production of layered oxides for LIBs, spherical dense particles were manufactured via co-precipitation of a hydroxide precursor which subsequently was calcined with various amounts of sodium hydroxide at a number of temperatures.

By structural analysis of a series of samples, a phase diagram was constructed. The obtained phase diagram is in good agreement with previous reported results (Paulsen and Dahn, 1999; Xiao et al., 2021) and provides the previously missing information for Mn/Ni = 3. At intermediate sodium content and high calcination temperature, phase pure P2 materials form.

Further investigations on the synthesis process of P2-Na_0.64Mn_3/4Ni_1/4O₂ were performed by ex-situ XRD, non-ambient in-situ XRD and TGA-DSC-MS:

(i) After wet impregnation and drying in ambient conditions, the material exhibits a mixture of the hydroxide precursor and a poorly crystalline P3 structure.

(ii) From room temperature to 450°C, a sodiation and crystallization of the P3 structure together with a complete oxidative removal of protons from the material is proposed.

(iii) From 450 to 800°C, further sodiation of the P3 structure together with a reduction of the transition metals is proposed.

(iv) Between 800 and 875°C, the structure is transformed from P3 to P2.

The obtained P2 material exhibits anisotropic broadening of the (10l) reflexes. A stacking fault model incorporating P3-type sodium sites was proposed. These stacking faults most likely remain between neighbouring P2 domains, during the structural transformation from P3 to P2.

The material reached an attractive initial specific discharge capacity of 202 mAh g⁻¹ in electrochemical half-cells at specific currents of 17.3 mA g⁻¹. A capacity fade associated with the 4.2 V voltage plateau occurred, which will be discussed in an upcoming publication. In accordance with literature, we found a solid solution regime (P2 structure) between 1.5 and 4.0 V (0.86 > x > 1/3) and a two-phase reaction (P2—O2) at the 4.2 V voltage plateau.

The stability of the material during prolonged ambient storage was investigated. Even after 14 days of storage, no hydrated phases were detected. In accordance with literature (Zuo et al., 2020), we found Na₂CO₃ on the particle surfaces of stored materials. No effect of these sodium carbonate surface impurities was found on the electrochemical cycling in half-cells.

With this publication, we provide fundamental insights into the synthesis of Na_xMn_3/4Ni_1/4O₂ layered oxides. The developed synthesis route is analogous to the highly scalable production process of NCM materials and leads to sodium cathode materials with technically relevant particle specifications. The obtained phase pure P2 materials, are solely based on low-cost and abundant raw materials, offer promising electrochemical performance and a high storage stability in ambient atmospheres.

Data Availability Statement

The raw data supporting the conclusions of this article will be made available by the authors, without undue reservation.

Author Contributions

LP synthesized the sodium containing materials, conducted most of the physical characterization and the electrochemical characterization, performed data analysis and drafted and edited the manuscript. NJ performed the TGA-DSC-MS measurements and assisted with evaluation of the obtained data. CG and MM performed the operando XRD measurement. ML provided the possibility for non-ambient XRD and contributed with his knowledge and many discussion to the quality of this manuscript. MM, SP, and MW-M provided guidance for materials and electrochemistry and contributed with their knowledge to the quality of this manuscript. PA supervised the work and co-edited this manuscript.

Funding

This work was funded by the German Federal Ministry of Education and Research (BMBF) in the projects TRANSITION (03XP0186C), ExcellBattMat (03XP0257A and 03XP0257C) and the German Research Foundation (DFG) in the project POLIS Cluster of Excellence (ProjectID 390874152).

Conflict of Interest

The authors declare that the research was conducted in the absence of any commercial or financial relationships that could be construed as a potential conflict of interest.

Publisher’s Note

All claims expressed in this article are solely those of the authors and do not necessarily represent those of their affiliated organizations, or those of the publisher, the editors and the reviewers. Any product that may be evaluated in this article, or claim that may be made by its manufacturer, is not guaranteed or endorsed by the publisher.

Acknowledgments

The authors thank Samuel Blessing for the non-ambient in-situ XRD measurement, Wolfgang Weirather for his work on the hydroxide precipitation, Gisela Arnold for ICP-OES measurements, Adam Reupert for the introduction to the operando XRD setup, Premsyl Beran for help with the refinement of XRD instrument parameters, Gemma Maskell for language editing and all colleagues for many productive discussions.

Supplementary Material

The Supplementary Material for this article can be found online at: https://www.frontiersin.org/articles/10.3389/fenrg.2022.910842/full#supplementary-material

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