Suppressed P2-Z phase transition and Fe migration in the Na layer of an Fe/Mn-based layered oxide cathode for advanced sodium-ion batteries

Abstract

P2-type Fe/Mn-based-layered oxides have attracted extensive attention as potential candidates for sodium-ion batteries owing to their low cost, abundant reserves, and high theoretical capacity. However, when the cathode operates within the voltage range of 1.5–4.3 V, unexpected Fe-ion migration and P2-Z phase transition in this cathode cause rapid capacity degradation and remain insufficiently resolved. In this work, harmful P2-Z phase transition, as well as the migration of Fe ions, is completely mitigated. Specifically, we found that Ti-ion doping could dramatically improve the structural stability of the Fe/Mn-based cathode, hinder the formation of Fe⁴⁺O₆ in the sodium-ion layer, and promote the migration of Na⁺. The designed Na_0.67(Fe_0.5Mn_0.5)_0.495Ti_0.05O₂ (NFMT-2) cathode could exhibit a discharge capacity of 182.7 mA h g⁻¹ within the voltage range of 1.5–4.3 V (0.1C) and achieve a smaller cell volume change (1.26%) during cycling than the bare Fe/Mn-based cathode. This work presents a new approach to inhibit P2-Z unfavorable phase transition and Fe migration in the sodium layer of Fe/Mn-based layered oxides.

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1. Introduction

With the growing demand for building an environmentally friendly, low-carbon and highly efficient energy-economy structure, lithium-ion batteries are limited by lithium reserves and are distributed unevenly to fulfill this demand; thus, there is an urgent need for the development of a new battery technology to replace it.^1–3 Lithium-ion batteries, which have the same working principle and production technology as sodium-ion batteries, owing to their wide source of sodium resources, low cost, excellent low-temperature performance, and other advantages that have received widespread attention, are a perfect substitute for lithium-ion batteries.^4–6 However, sodium-ion batteries have some drawbacks, such as the larger mass and ionic radius of Na (Na vs. Li, 1.02 Å vs. 0.76 Å), leading to a lower energy density and diffusion rate.^7–9 Since the cathode is a vital component in determining the performance of batteries, it is important to explore low-cost cathode materials that can achieve rapid Na⁺ migration and superior cycling stability.

Sodium ion battery cathode materials fall into three main categories based on their different structures: layered transition metal oxides (Na_xTMO₂, TM = transition metal), polyanionic materials, and Prussian blue materials. Overall, layered transition metal oxidation has drawn much attention in recent years owing to its high energy density, simple synthesis, diverse compositions, and environmental friendliness.^10–12 Based on the coordination configuration of sodium ions and the stacking of oxygen, Delmas classified layered oxides into four structures: P2, P3, O2, and O3-type.^13,14 The letters P and O represent sodium ions in the trigonal and octahedral positions, respectively. Numbers 2 and 3 represent the number of layers stacked in the smallest repeating unit of oxygen.^15,16 In comparison with other structures, P2-type layered oxides have better Na ion transport pathways and rate properties because their Na ions are located in trigonal prismatic sites.^17–19 Moreover, cost is crucial for the application of cathodes in sodium-ion batteries. Hence, there is emerging interest in cost-effective and environmentally friendly P2-type Fe/Mn-based layered oxides. Komaba reported a binary iron–manganese-based layered oxide (P2-Na_0.67Fe_0.5Mn_0.5O₂) synthesized by applying the solid-phase method.²⁰ The material was able to integrate the advantages of the high operating voltage of Fe³⁺/Fe⁴⁺ and the large specific capacity of Mn³⁺/Mn⁴⁺, exhibiting a high specific capacity of 190.0 mA h g⁻¹ over a wide voltage range of 1.5–4.3 V. However, this Fe-containing cathode material, in which Fe ions migrate to the sodium layer, hinders sodium ion insertion.^21,22 The Jahn–Teller effect of the irreversible phase transition from P2–O2 and redox of TMs such as Fe and Mn at high voltages exacerbates the irreversible migration of TMs, causing the deterioration of cycling stability. In previous studies, low-valent elements, such as Li⁺,^23,24 Ni^2+ 25,26 and Co^3+ 27,28 were introduced into the crystal structure to mitigate the P2–O2 irreversible phase transition and the Jahn–Teller effect to improve the cycling performance and stabilize the structure of Fe/Mn-based layered oxides. Unfortunately, the layered oxide still suffers from an unfavorable P2-Z phase transition (incomplete conversion from P2 to O2) and unfavorable migration of Fe ions to the sodium layer, resulting in poor diffusivity of sodium ions, accumulation of stress during cycling and the formation of structural defects, which is the main reason for the poor cycling stability of the material.²⁹ Therefore, there is still a need for further solutions to address this challenge.

Herein, a Ti-doped Na_0.67(Fe_0.5Mn_0.5)_0.495Ti_0.05O₂ (denoted as NFMT-2) cathode material was successfully synthesized using a high-temperature solid-phase method combined with annealing treatment. The unfavorable phase transition from P2 to Z in the high-voltage region was successfully suppressed over a wide voltage range of 1.5–4.3 V. Ti doping effectively inhibits the oxidation of Fe³⁺ to Fe⁴⁺ during charging, reduces the amount of Fe⁴⁺O₆, and hinders the migration of Fe ions to the Na layer, further improving retention during the cycling process. Additionally, Ti doping increases in the crystal plane spacing and modulates the relative thicknesses of the NaO₂ and MO₂ layers. Thus, the NFMT-2 cathode material has significantly improved cycling stability and multiplication capacity (14.3% increase in capacity retention for 100 cycles at 1C and from 39% to 58% at 1C/0.1C). Additional in situ XRD and ex situ XAS experiments demonstrated the valence alterations of Mn and Fe as well as the regulatory mechanism of the coordination environment; successful Ti doping promoted the reversible conversion of Mn³⁺/Mn⁴⁺ and Fe³⁺/Fe⁴⁺. GITT and CV tests were used to confirm the increased diffusion rate of the sodium ions. This work offers a new perspective on the investigation of transition metal migration and P2-type Fe-/Mn-based layered oxide materials.

2. Experimental section

2.1. Material synthesis

Na_0.67(Fe_0.5Mn_0.5)_1−xTi_xO₂ (x = 0, 0.02, 0.05, and 0.08, denoted as NFM, NFMT-1, NFMT-2 and NFMT-3, respectively) was prepared by applying the solid-phase method in combination with annealing treatment. The specific process was as follows: Na₂CO₃ (5% excess), Mn₂O₃, Fe₂O₃, and TiO₂ were weighed according to the molar ratio of metal elements as n(Na) : n(Mn) : n(Fe) : n(Ti) = 0.67 : (0.5 − x/2) : (0.5 − x/2) : x (x = 0; 0.03; 0.05; 0.08). After placing the weighed raw materials in a ball milling jar containing onyx balls, 50 mL of ethanol was added as a dispersant to promote homogeneous dispersion of the mixture, and mixing was carried out using a planetary ball mill (speed: 300 rpm, time: 6 h). After refinement by ball milling, the mixture was placed in a drying oven at 110 °C for 3 h to obtain a well-mixed and dried precursor. Then, the precursor was compacted and transferred to a crucible at a pressure of 20 MPa and heated up to 900 °C for 12 h at a temperature increase rate of 5 °C min⁻¹. After the calcination process was finished, the material was cooled to room temperature in a desiccator set to a constant temperature of 40 °C (cooling at a rate of approximately 80 °C min⁻¹). The final product was then immediately moved for storage into a glove box filled with argon.

2.2. Characterizations

In this paper, X-ray powder diffraction tests were carried out on the samples using the Rigaku Smartlab of Rigaku Corporation (30 mA, 40 kV), Japan, with Cu-Kα as the radiation source (λ = 1.5406 Å). The morphology and microstructure of the samples were examined via transmission electron microscopy (TEM, FEI Tecnai G2 F30, operating at 200 kV) and scanning electron microscopy (ZEISS SUPRA 55, Germany), respectively, with elemental analyses performed utilizing TEM and SEM energy spectrometers (EDS). Elemental valence states were measured by X-ray photoelectron spectroscopy (XPS, THERMO ESCALAB 250XI), and the valence states of Fe were determined by Mössbauer Spectrometry (MS-96), with ⁵⁷Co as the radioactive source. The X-ray absorption spectra (XAS) of the Mn, Fe and K-edge were measured in transmission and fluorescence mode at the XRD station of beamline 4B9A of Beijing Synchrotron Radiation Facility (BSRF). In situ X-ray diffraction Be window batteries were built, and data for the first charge/discharge cycle were obtained at a scanning rate of 0.01° s⁻¹ between 14 and 45°.

For the in situ XRD test, a specially designed Be window battery was built to monitor real-time reactions during the initial cycle. Each scan acquires data between 2θ = 10° and 45° in incremental steps of 0.02°. The structural evolution of the samples was tested in the 2θ range of 14–45° by employing Tongda TD-3500 with Cu Kα radiation (λ = 1.5418 Å).

2.3. Electrochemical test

The active material (80 wt%), acetylene black (10 wt%), and polyvinylidene fluoride (PVDF, 10 wt%) were mixed well and then applied to the aluminum foil. N-Methyl-2-pyrrolidinone (NMP) was used as a solvent, and the mixture was vacuum-dried for an entire night at 80 °C. Subsequently, the dried electrode was punched into disks with a diameter of 10 mm. The CR2032 coin battery was put together in a glove box in an argon environment. A glass fiber membrane (GF/C, Whatman) was used as the separator, and Na metal was used as the counter electrode. The electrolyte was 1 M NaClO₄ dissolved in propylene carbonate (PC) with 5% fluoroethylene carbonate (FEC). The Land system (CT2001A) was used to perform galvanostatic charge/discharge tests and galvanostatic intermittent titration technique (GITT) tests in the voltage range of 1.5–4.3 V (1C = 200 mA g⁻¹). Cyclic voltammetric (CV) curves were measured at different scan rates (0.1–1.0 mV s⁻¹) on an electrochemical workstation (CHI 660E).

3. Results and discussion

A series of NFMT (Na_0.67(Fe_0.5Mn_0.5)_1−xTi_xO₂) materials doped with different Ti contents were prepared by applying the high-temperature solid-phase method, and XRD tests were carried out to analyze the effect of Ti doping on the crystal structure of the materials. As shown in Fig. 1a, the positions and relative intensities of the diffraction peaks of the NFM sample coincide with those of the standard PDF card (No. 00-54-0894) of the P2-phase hexagonal crystal structure (space group P6₃/mmc), without the appearance of other impurity peaks, which indicates that the material with a pure P2-phase hexagonal crystal structure was successfully synthesized, and the high diffraction peak intensities indicate that the material is well crystallized. It is noteworthy that the diffraction peaks of the NFMT-1, NFMT-2, and NFMT-3 samples are essentially the same as those of the NFM samples, indicating that the P2 phase structure of the materials is not altered by a modest amount of Ti doping. Furthermore, it is observed that the diffraction peaks corresponding to the (002) crystal faces move to lower angles as the doping quantity increases. Moreover, the facet spacing d is inversely proportional to the angle θ, suggesting that both the cell parameter c and the facet spacing d increase as a result of successful Ti⁴⁺ doping. Larger crystal spacing accelerates sodium ion diffusion, offering the material superior rate properties. Four sets of samples underwent Rietveld refinement of the XRD spectra to examine in more detail the impact of Ti substitution on the crystal structure of the materials. The outcomes are displayed in Fig. 1b and Fig. S1 (ESI†). An excellent fit between the refining findings and the experimental data has been established, as evidenced by the R_wp values of all four sets of refining data being less than 10%. Table S1 (ESI†) displays the crystallographic characteristics following refinement. The lattice parameters a and c increase as the Ti doping content increases, which is attributed to the fact that the strength of the Ti–O bond (662 kJ mol⁻¹) is greater than those of the Fe–O (409 kJ mol⁻¹) and Mn–O bonds (402.3 kJ mol⁻¹), and that the radius of Ti⁴⁺ (0.605 Å) is smaller than those of Mn³⁺ (high spin, 0.645 Å) and Fe³⁺ (0.645 Å).^30–32 Upon Ti doping, the O–Na–O–TM–O–Na–O structure along the c-axis direction decreases the strength of the O–Na–O bonds. Because the TM–O bond is stronger, the corresponding d_O–Na–O becomes larger.³³ Because Na⁺ diffusion is a mechanism for sodium ion migration in a two-dimensional NaO₂ layer, which is directly related to the thickness of the NaO₂ layer, there is a possibility that iron ions adjacent to MO₂ will migrate to the NaO₂ layer to obstruct the sodium ion diffusion channel.³⁴ Therefore, Ti doping achieves the modulation of the relative thickness of the NaO₂/MO₂ layer of the NFMT-2 material and is an effective way to expand the Na layer spacing, which is conducive to the enhancement of the structural stability and the improvement of the sodium ion migration kinetics. High bonding energy Ti–O bonds, however, can enhance the material's cycling stability by stabilizing the NFMT-2 crystal structure and preventing the P2-Z phase transition from occurring during cycling. Fig. 1c schematically illustrates the crystal structure of NFMT-2. The Fe/Mn/Ti ions are randomly distributed in the octahedral sites of TMO₂ slabs. To observe the effect of different Ti doping amounts on the micro-morphology of the materials, SEM measurements were carried out on the four groups of synthesized samples, and it can be clearly observed from Fig. 1d that the micro-morphology of the materials is in the form of lamellar layers, with particle sizes of 0.5–3 μm. The particle size of the materials decreases, and the agglomeration phenomenon is suppressed with an increase in the Ti doping amount. The fine and uniform particle morphology facilitates the migration of Na ions, improves the multiplicity performance of the material and shows better structural stability during the cycling process. The crystal structures of the NFM and NFMT-2 cathodes were tested by TEM, as shown in Fig. 1e and f, and both displayed highly crystalline structures (No. 00-54-0894), with a face spacing of 0.563 nm on the (002) face of NFM and 0.568 nm on the (002) face of NFMT-2, indicating that Ti doping resulted in a larger face spacing. The thickness of the NaO₂/MO₂ layer was successfully regulated, which is consistent with the previous XRD analysis. The selected area electron diffraction (SAED) pattern along the [001] zone axis reveals the single-crystal nature of the hexagonal P2-NCLFMO structure (Fig. 1g). In addition, Fig. 1h demonstrates the EDS map of the NFMT-2 sample, and it can be clearly observed that all these elements are uniformly distributed on the surface of the cathode material. This illustrates that the sample was well controlled during the preparation process, which allowed the various raw materials to be fully mixed during the reaction process, thus ensuring the chemical homogeneity of the sample.

Fig. 1 (a) XRD patterns of NFM and NFMT materials. (b) Rietveld refinements of the XRD patterns of NFMT. (c) Schematic crystal structure of NFMT. (d) SEM image of NFMT-2. (e) HRTEM image of the NFM. (f) HRTEM image. (g) SAED pattern of NFMT-2. (h) TEM-EDS mapping images of NFMT-2.

Further, the change in transition metal valence state after Ti doping by synchrotron X-ray absorption spectroscopy (XAS) analysis on the Ti-doped sample (NFMT-2) and pristine NFM sample was investigated. The X-ray absorption near-edge structure (XANES) spectra of the Fe and Mn K-edges are shown in Fig. 2. The K-edge XANES spectra of Fe are compared with the standard reference spectra of Fe₂O₃, as shown in Fig. 2a, where the positions of the absorption edges in the NFM and NFMT-2 samples overlap, indicating that the Fe valence state remains unchanged after doping, with both being +3 valence. Fig. 2b illustrates the Fourier-transformed EXAFS spectra in the R-space. The first peak at the lower radial distance represents the TM–O coordination shell layer in the TMO₆ octahedron, while the second shell layer at the higher radial distance is associated with the TM–TM interaction. The intensity of the two scattering peaks of NFMT-2 varied after Ti doping, which was mainly attributed to the fact that Ti partially occupies the lattice sites of Fe, thus decreasing the intensity of the scattering peaks.³⁵ The wavelet-transformed (WT) EXAFS spectra of NFM and NFMT-2, (Fig. 2e and f, respectively), which visualize the change in bond lengths after doping, where the two scattering peaks near 1.5 and 2.5 Å correspond to the TM–O and TM–TM bonds. With Ti doping, the valence state and coordination environment of Fe were not significantly changed, and Ti⁴⁺ has a similar ionic radius to Fe³⁺. Thus, a small amount of doping did not cause significant changes in the crystal structure. Additionally, the K-edge XANES spectra of Mn were compared with the standard reference spectra of Mn₂O₃ and MnO₂ (Fig. 2c), and Mn in both the NFM and NFMT-2 samples was in a mixed valence state between +3 and +4 valence. The relatively low K-edge energy of Mn in the NFMT-2 sample compared to NFM suggests that Ti doping reduces the average oxidation state of Mn. The XANES spectrum of Mn was linearly fitted to show that the valence state of NFM is 3.55 and that of NFMT-2 is 3.47 (Fig. 2d). This is the result of the introduction of Ti⁴⁺ into the lattice, which undergoes the transfer of electrons to and from the surrounding atoms.³⁶ The charge transfer effect causes the average oxidation state of Mn to decrease to maintain an even distribution of charge in the transition metal layer. To further investigate the coordination structure changes in the Ti-doped samples, the Mn in R-space Fourier transformed XAFS spectra were utilized for characterization. Fig. 2g and h show the wavelet transform (WT) EXAFS spectra of NFM and NFMT-2, respectively, where the two scattering peaks around about 1.5 and 2.5 Å correspond to the Mn–O and Mn–TM bonds. The coordination environments of Mn before and after Ti doping did not change significantly, which may be attributed to the fact that the radius of Ti⁴⁺ is similar to that of Mn³⁺. It has been demonstrated that a small amount of doping allows the material to maintain a favorable crystal structure. Meanwhile, X-ray photoelectron spectroscopy (XPS) was utilized to analyze the surface chemical composition and surface valence states of the NFMT-2 samples. Fig. 3a shows the full spectrum of the NFMT-2 cathode material, with peaks detected for the following elements: Na (∼1071 eV), Mn (∼642 eV, ∼654 eV), Fe (∼711 eV, ∼724 eV), O (∼534 eV) and Ti (∼460 eV).^37,38 In addition, the plots fitted to the XPS data are shown in Fig. 3b–d. The binding energies of Ti 2p^3/2 and Ti 2p^1/2 are 457.9 eV and ∼463.6 eV, respectively, indicating that Ti ions exist only in a valence state of +4. Combined with the shift of the (002) plane in XRD, it proves that Ti was successfully doped into the lattice (Fig. 3b).³⁹Fig. 3c shows the results of the fits for the Mn element, with the Mn 2p^3/2 peaks fitted at 641.8 eV (Mn³⁺) and 643.2 eV (Mn⁴⁺), and the Mn 2p^1/2 peaks fitted at 653.3 eV (Mn³⁺) and 654.6 eV (Mn⁴⁺), meaning that the Mn element is in the intermediate valence state of the material,^40,41 which is in agreement with the XANES results. In Fig. 3d, the binding energies of Fe 2p^3/2 and Fe 2p^1/2 are 711 eV and 724 eV, respectively, indicating that the Fe element of the NFMT-2 samples is mainly +3.^42,43

Fig. 2 (a) XAS spectra at the Fe K-edge and (b) corresponding Fourier-transformed EXAFS spectra of NFM and NFMT-2. (c) XAS results at the Mn K-edge and (d) corresponding integration method of NFM and NFMT-2. (e) and (f) Fe and (g) and (h) Mn K-edge WT-EXAFS spectra of NFM and NFMT-2.

Fig. 3 XPS spectra of NFMT-2 (a) full spectrum, (b) Ti 2p, (c) Mn 2p, and (d) Fe 2p.

To evaluate the effect of Ti doping on the electrochemical properties of Na_0.67Fe_0.5Mn_0.5O₂ cathode materials, the electrochemical performance tests were performed in a potential window of 1.5–4.3 V. Fig. 4a demonstrates a comparison of the rate capability tests for the four groups of samples. NFMT-2 has the best rate capability, and the rate performance of NFMT-2 (1C/0.1C = 58%, 182.7 mA h g⁻¹ at 0.1C, compared with 105 mA h g⁻¹ at 1C) is better than that of NFM (1C/0.1C = 39%, 198.1 mA h g⁻¹ at 0.1C, compared with 78.2 mA h g⁻¹ at 1C). The enhancement of the rate performance is attributed to the modulation of the relative thickness of the NaO₂/MO₂ layer in the NFMT-2 material after Ti doping, and the enlarged spacing of the Na layer contributes to the enhancement of the structural stability and the improvement of the sodium ion migration kinetics. Moreover, Fig. 4b illustrates the cycling performances of NFM, NFMT-1, NFMT-2, and NFMT-3 for the first 100 cycles at 1C. The capacity retention after 100 cycles was 57.1%, 61%, 71.4%, and 64.8%, respectively, and the NFMT-2 sample with 0.05 Ti doping showed the best cycling performance, with about 20% improvement in cycling performance compared to the undoped NFM sample. Notably, the discharge specific capacity of the material decreases with the increase in the doping level, which is attributed to the fact that Ti is a nonactive element not involved in the redox reaction during the charging and discharging processes of the material. In comparison, the NFMT-2 sample with Ti doping of 0.05 was found to have the best overall performance and hence was selected for the main investigation. A comparison of the charging and discharging curves of NFM and NFMT-2 cathode materials for the first 20 laps at 0.1C is shown in Fig. 4c and d, where the first time Coulombic efficiencies of the materials are greater than 100% because both NFM and NFMT-2 are in the sodium deficient phase, and there are additional sodium ions embedded in the materials during the discharging process. The charge/discharge curves of both NFM and NFMT-2 exhibited two discharge plateaus, in which the discharge plateau of the Mn³⁺/Mn⁴⁺ redox pair was located around 2.1 V, and the voltage plateau around 3.3 V was related to the participation of Fe³⁺/Fe⁴⁺ in the redox reaction. After 20 cycles, the capacity retention of the NFMT-2 sample was 86.3%, which was significantly higher than that of the undoped NFM sample. In addition, the NFM sample has an obvious voltage decay at the voltage plateau at 3.7 V, while the NFMT-2 sample has a better reversibility in comparison. This indicates that the doping of Ti can inhibit the participation of Fe³⁺/Fe⁴⁺ in the redox reaction, effectively alleviate the irreversible migration of Fe³⁺/Fe⁴⁺ and improve the cycling performance of the material. In Fig. 4e, the dQ/dV peak (∼3.4 V) of NFM is displaced towards a lower voltage, indicating that the irreversibility of Fe³⁺/Fe⁴⁺ redox increases with the increase in the number of cycles, leading to an increase in the proportion of Fe⁴⁺ and a rapid capacity decay. In contrast, NFMT-2 exhibits a highly reversible dQ/dV curve in Fig. 4f, with the corresponding Fe³⁺/Fe⁴⁺ reduction peak shifting to higher voltages throughout the cycle, indicating a decrease in the Fe⁴⁺ proportion. Moreover, the dQ/dV peaks (1.75–2.75 V) for Mn³⁺/Mn⁴⁺ of the NFM materials were shifted to lower voltages, and Mn³⁺ with the Jahn taller effect caused the structural transition, whereas NFMT-2 displayed greater reversibility.

Fig. 4 (a) Rate capabilities at variable rate currents of NFM and NFMT, and (b) cycling performance at 1C of NFM and NFMT cathodes. (c) and (d) Charge-discharge curves of NFM and NFMT-2 electrodes for the first 20 cycles, respectively, and (e) and (f) the corresponding dQ/dV curves.

Improvement of sodium ion migration kinetics and electrochemical polarization of NFMT-2 materials by slowing down the irreversible migration of Fe³⁺/Fe⁴⁺ was further explored by combining GITT and CV tests. Fig. 5a and c show the GITT curves of NFM versus NFMT-2 during charging and discharging, and Fig. 5b and d show the corresponding ion diffusion coefficient plots, respectively. The polarization of the NFMT-2 is significantly reduced compared to the potential profile of the NFM anode. In addition, it should be noted that the sodium ion diffusion coefficient (D_Na⁺) of the NFMT-2 anode was 5.29 × 10⁻¹⁰ cm² s⁻¹ during charging and 2.99 × 10⁻¹⁰ cm² s⁻¹ during discharging, which were both above the NFM (charging: 4.12 × 10⁻¹⁰ cm² s⁻¹; discharging: 2.99 × 10⁻¹⁰ cm² s⁻¹). GITT calculation details are presented in the ESI.† The enhanced sodium migration kinetics of the NFMT-2 cathode material can be attributed to the effective expansion of the Na⁺ layer spacing after Ti doping, which is conducive to enhance the structural stability and acceleration of Na⁺ migration. D_Na can also be calculated from the CV curves of different scanning rates (Fig. 5e and f). The D_Na values of the O1, O2, R1, and R2 peaks are all at 10⁻¹⁰ cm² s⁻¹ (calculation details are described in the ESI†), matching well with the GITT results. In summary, the GITT and CV results indicate that the NFMT-2 electrode exhibits faster Na⁺ diffusion kinetics.

Fig. 5 Electrode process kinetics of the NFM and NFMT-2 electrodes: (a) and (c) charge–discharge GITT profiles in the first cycle, and (b) and (d) the corresponding Na⁺ diffusion coefficients (D_Na). (e) CV curves at varying scan rates from 0.1 to 1 mV s⁻¹ of NFMT-2 and (f) the corresponding fitting curves between peak currents (i_p) and the square root of scan rates (v^1/2).

Then, the structural evolution of the NFMT-2 material during the initial desodiation/sodiation process was monitored utilizing in situ XRD techniques at a voltage interval of 1.5–4.3 V, with the test diffraction angle set in the range of 14–45°. In Fig. 6a, the sodium ions are gradually removed from the crystal structure during charging up to 4.3 V. The (002) and (004) peaks shift to a lower angle while the transition metal is oxidized, increasing the c-axis spacing and decreasing the a-value of the transition metal layer spacing, which is induced by an increase in the electrostatic repulsion between neighboring O–O layers as oxidation progresses. The diffraction peaks show a displacement in the opposite direction of charging during the discharging process. Notably, the cell volume changes only 1.26% throughout the charging and discharging processes (Fig. 6b), exhibiting excellent low-strain characteristics. Moreover, throughout the test, no other peaks appeared or the original peaks disappeared, indicating no phase transition in the NFMT-2 material. This is completely different from the ex situ XRD of the NFM samples, which shows a Z-phase (incomplete transformation of the P2–O2 phase) during deep desodiation (Fig. 6d). The change in structure in the range of 1.5–4.3 V is displayed in Fig. 6e. The results indicated that the NFM-2 material successfully inhibited the deleterious phase transformation of the material during the cycling process and improved the structural stability of the material during the cycling process.

Fig. 6 Na-storage mechanisms: (a) In situ XRD patterns of the NFMT-2 electrode during the first charge/discharge process at 0.1 C. (b) The evolution of lattice parameters and unit-cell volume upon charge/discharge during the first cycle for NFMT-2. (c) Magnified contour plot of the (002) diffraction peak. (d) Ex situ XRD patterns of NFM for each point of the electrochemical profile. (e) Structure evolution diagrams of NFMT-2.

To investigate the migration of Fe ions due to the successive desodiation/sodiation of sodium ions, the NFM and NFMT-2 electrodes were first treated with 5 charge/discharge cycles and then charged to 4.3 V to the material, after which Mössbauer Spectrometry tests were performed (Fig. 7a and b). The Na⁺ detachment of NFM and NFMT-2 causes the diphasic asymmetry with an additional spectral weight to shift to negative velocities, which is consistent with the results in the literature.^44–46 The spectra of NFM and NFMT-2 were fitted into two components: the Fe³⁺O₆ (octahedral environment) and Fe⁴⁺O₆ components. The main Fe³⁺O₆ component has a homogeneous isomerization shift of about 0.32–0.37 mm s⁻¹ and a quadrupole split of 0.68–0.86 mm s⁻¹. The second component, Fe⁴⁺O₆, exhibited a homogeneous isomerization shift of 0.34–0.42 mm s⁻¹ and a quadrupole splitting of 0.62–0.73 mm s⁻¹. All components are consistent with an iron-based sodium ion-layered oxide. Based on the literature, one of the reasons for the capacity decay of iron-based cathode materials during cycling is that the migration of iron ions hinders the embedding of sodium ions.^21,22,47 When the NFMT-2 sample was charged to 4.3 V, the Fe⁴⁺O₆ component was 27.42%, while the NFM sample was 31.15% (Table S2, ESI†). According to the DFT analysis of Fe-based sodium-ion battery cathode materials by Li, the high-spin state Fe⁴⁺O₆ with nickel ferrous-layered structure leads to the migration of Fe to the sodium layer.⁴⁸ Therefore, Ti doping can effectively inhibit Fe³⁺ from being oxidized by Fe⁴⁺ during charging, reduce the amount of Fe⁴⁺O₆, and hinder the migration of Fe ions to the sodium layer. To clarify the charge compensation mechanism, ex situ XAS tests were performed on the NFMT-2 samples. The valence changes in transition metals during charging and discharging can be analyzed by comparing the shifts of the pre-edge peaks of the X-ray absorption near edge structure (XANES) spectra at different voltages. Fig. 7d, e and g, h show the ex situ XANES spectra of the Mn K and Fe K-edges, respectively. During charging, the NFMT-2 Mn K-edge XANES spectra show a significant shift corresponding to the Mn³⁺ to Mn⁴⁺ oxidation process, suggesting that Mn is involved in the charge compensation process from the pristine state to 4.3 V. Then, at the end of the discharge, the K-edge spectrum of Mn returns to the lower energy region, showing that Mn³⁺/Mn⁴⁺ is a reversible process in redox. The XANES spectra of Fe³⁺ and Fe⁴⁺ are similar, with relatively weak shifts in the pre-edge peaks. However, a shift in the K-edge spectra of Fe to the higher energy region can be observed through the magnified image in Fig. 7g, which corresponds to the oxidation of Fe³⁺ during this charging process. At the end of the discharging, the K-edge spectra of Fe return to the energy position of the original sample. The Fourier-transformed extended X-ray absorption fine structure (EXAFS) spectra at the Mn and Fe K-edges are displayed in Fig. 7f and i, respectively. The peaks at lower radial distances correspond to the coordination of transition metal (TM) and oxygen (O) in the first shell layer, while the second shell layer at higher radial distances is related to TM–TM interactions. Through the relative changes in bond lengths, the coordination environment changes in the material crystals during the charging and discharging processes can be effectively reflected.⁴⁹ Benefiting from the more stable crystal structure after Ti doping, the local environments of Mn and Fe did not change significantly during the charging and discharging processes and showed high structural reversibility, which is consistent with the results of the in situ XRD.

Fig. 7 Ex situ ⁵⁷Fe-Mössbauer spectroscopy for (a) NFM and (b) NFMT-2 charging to 4.3 V. (c) GCD curves and corresponding points of NFMT-2 electrode for taking Ex situ XAS measurement. Ex situ XANES spectra of the NFMT-2 K-edge: (d) Mn K-edge for the charging process, (e) Mn K-edge for the discharging process, (g) Fe K-edge for the charging process, and (h) Fe K-edge for the discharging process. (f) Mn and (i) Fe K-edge Fourier-transformed EXAFS spectra at different voltages of NFMT-2.

4. Conclusion

In summary, a series of Na_0.67(Fe_0.5Mn_0.5)_1−xTi_xO₂ (x = 0, 0.02, 0.05, 0.08) materials doped with different Ti contents are prepared by applying the high-temperature solid-phase method. On the one hand, Ti doping effectively inhibits Fe³⁺ from being oxidized to Fe⁴⁺, thereby alleviating the migration of iron ions to the sodium layer and achieving the complete suppression of the unfavorable phase transition in a wide voltage interval of 1.5–4.3 V. XAS and Mössbauer Spectrometry tests further demonstrated that the ratio of Fe⁴⁺O₆ was reduced. On the other hand, compared with Fe–O and Mn–O bonds, stronger Ti–O bonds widen the crystal plane spacing of the Na layer, thereby modulating the apparent relative thickness of the NaO₂/MO₂ layer and improving the migration rate of Na⁺, as confirmed by XRD refinement and HR-TEM. Moreover, GITT and CV tests reveal that the NFMT-2 material possesses fast reaction kinetics, achieving excellent rate capability (182.7mA h g⁻¹ at 0.1C, 105 mA h g⁻¹ at 1C) and good cycling performance (71.4% capacity retention after 100 cycles at 1C, a 14.3% improvement over NFM). In situ XRD showed that the P2 pure phase was maintained throughout the desodiation/sodiation process with only a 1.26% volume change. Ex situ XAS tests demonstrated the reversible conversions of Mn³⁺/Mn⁴⁺ and Fe³⁺/Fe⁴⁺. This work provides a new direction for the development of P2-phase Fe-/Mn-based-layered oxides and contributes to the advancement of next-generation low-cost, high-performance sodium-ion batteries.

Data avaialability

Data are available from the authors upon request.

Conflicts of interest

The authors declare no conflict of interest.

Acknowledgements

This work was supported by the National Natural Science Foundation of China (52274295, 11775226), Natural Science Foundation of Hebei Province (E2021501029,A2021501007), the Fundamental Research Funds for the Central Universities (N2423051, N2423053, N2423005, N2302016, N2423019), the Science and Technology Project of Hebei Education Department (QN2024238), the Basic Research Program Project of Shijiazhuang City for Universities Stationed in Hebei Province (241790937A), the Science and Technology Project of Qinhuangdao City (202302B006), the Open Research Subject of Key Laboratory of Advanced Electrode Materials for Novel Solar Cells for Petroleum and Chemical Industry of China (2004A052). Hebei Provincial Doctoral Candidate Innovation Ability Training Funding Project (CXZZBS2024176). Hebei Province Science and Technology Research and Development Platform Special Innovation Capability Enhancement Plan Project (24464402D). Thanks eceshi (www.eceshi.com) for the TEM test. The authors also would like to thank the shiyanjia lab (www.shiyanjia.com) for the XPS test. The authors wish to thank facility support of the 4B9A beamline of Beijing Synchrotron Radiation Facility (BSRF).

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Footnote

† Electronic supplementary information (ESI) available. See DOI: https://doi.org/10.1039/d4tc04961f

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