A well-designed P2-Na_0.67Mn_0.85Al_0.05Zn_0.1O₂ cathode for superior sodium-ion batteries†

Xiang Ding*^a, Cuixian Hu^a, Yong Fan^a, Yue Lin^b, Jinyang Liu^c, Yibing Yang^b, Liangwei Liu^b, Haibao Ma^b, Yi Xiao^b and Lili Han*^b
aCollege of Chemistry and Materials Science, Fujian Normal University, Fuzhou 350007, China. E-mail: dingx@finu.edu.cn
^bState Key Laboratory of Structural Chemistry, Fujian Institute of Research on the Structure of Matter, Chinese Academy of Sciences, Fuzhou 350002, China. E-mail: llhan@fjirsm.ac.cn
^cCollege of Physics and Energy, Fujian Normal University, Fuzhou 350007, China

First published on 2nd August 2024

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1. Introduction

Lithium-ion batteries (LIBs) and sodium-ion batteries (SIBs) have become the favorites of intelligent driving vehicles (IDVs) and large-scale energy storage (LSES) in the recent decade. Wherein, SIBs display greater potential in next generation high-performance energy storage devices due to their high Na abundance (2.64%), lower manufacturing costs and superior performance at low-temperature.^1–3 The positive materials,⁴ consisting of layered metal oxides (e.g. [Na_0.67MnO₂], [NaNiO₂]) and polyanionic compounds (e.g. [Na₃V₂(PO₄)₃]), Prussian blue analogues (e.g. [NaFeFe(CN)₆], [Na₂MnFe(CN)₆]), and tunnel-type oxides (e.g. [Na_0.44MnO₂]) and organics have been extensively investigated in both academia and industry. Herein, the P2-Na_0.67MnO₂ (NMO) material has the merits of handy preparation and high theoretical capacity (175 mA h g⁻¹) that are considered to be of great research value.^5–7 However, NMO also suffers from the defects of the complex P2–O2 phase transition and inferior electronic conductivity.^8,9 The P2–O2 transition is owing to the Jahn–Teller effect of Mn³⁺ during the electrochemical cycling, resulting in the slippage of layers.^10,11 Besides, the fact that only one type of metal ion (Mn) is situated in the transition metal (TM) layer gives rise to the relatively low electronic density as well as transition probability.^12,13 Hence, the intrinsic electronic transfer of bulk materials is constrained.

To address the above issues,^14,15 the bulk lattice doping strategy has been extensively employed, mainly consisting of alkalis (e.g. Li⁺), alkaline-earth metals (e.g. Ca²⁺ and Mg²⁺) and multi-valent ions (e.g. Zn²⁺, Fe³⁺, Co³⁺, and Ti⁴⁺) incorporated into the Mn site, respectively. It can adjust the electronic and spatial structures by creating a multi-electronic environment and providing pillar effects, thus boosting the bulk electron conductivity, Na⁺ diffusion efficiency and structural stability. Ling et al. reported Na_0.67Cu_0.1Mn_0.9O₂ that manifested a suppressed Jahn–Teller effect as well as comparatively less phase transition. It can deliver a high capacity (222.7 mA h g⁻¹@10 mA g⁻¹; 1.5–4.2 V) with capacity retention (76%@1 A g⁻¹@300 cycles).¹⁶ Zhang's group introduced electrochemically active Fe³⁺ into the Mn-site to enhance the Na⁺ diffusion kinetics. The optimized Na_2/3Ni_1/3Mn_7/12Fe_1/12O₂ exhibited long-term cycling stability with a fading rate of 0.05% per cycle over 300 cycles.¹⁷ Most recently, a series of co-doping studies have been performed and turned out to be far more efficient for optimizing the electrochemical properties and structures than the corresponding individual-doping ones.¹⁸ Generally, co-doping has emerged as the preferable optimization technique due to the more comprehensive regulation on each lattice site, which can lead to increased delocalized electron density, ionic radius and electronegativity rooted in the introduced ions.¹⁹ Luo et al. proposed a Cu/Fe co-doped Na_0.67Mn_0.92Fe_0.04Cu_0.04O₂ cathode that can deliver superior discharge capacity (90.68%@110.5 mA h g⁻¹@100 cycles@5C).²⁰ Huang et al. adopted Zn/Mg dual-doped P2-Na_0.67MnO₂ which demonstrates a high performance of 67.2 mA h g⁻¹@10C.²¹ Hence, the synergistic effect of co-doping into the Mn-site is always more effective in performance optimization.

In this work, a series of Zn/Al co-doped P2-Na_0.67Mn_1−x−yAl_xZn_yO₂ are well-designed by regulating different dopant amounts. A range of in situ and ex situ characterization techniques and DFT calculations systematically elucidate the Zn/Al co-doping synergistic effect, occurrence of of suppressed P2–O2 phase transition and much enhanced structural stability and bulk electric/ionic conductivities during cycles. Consequently, the optimal P2-Na_0.67Mn_0.85Al_0.05Zn_0.1O₂ (NMO3) can deliver superior discharge capacity (181.85 mA h g⁻¹@0.1C) and cycling stability (83%@93.87 mA h g⁻¹@600 cycles@5C) in half-cells and sufficient energy density (334.5 W h kg⁻¹) in NMO3//HC full-cells. This efficacious optimization tactic of collaborative doping is above rubies for designing pragmatic layered P2-phase materials for SIBs.

2. Experimental section

2.1 Materials preparation

The Na_0.67MnO₂ (NMO) cathode is synthesized via a typical sol–gel method.^22,23 Na₂CO₃ (5% excess), Mn(CH₃COO)₂·4H₂O and anhydrous citric acid are dissolved in 50 mL deionized water according to the stoichiometric ratio. The as-obtained solution is continuously stirred to achieve a transparent solution. Next, it was dried at 200 °C for 12 h in an electric oven. The NMO product is obtained by programmed heating for two steps (450 °C for 6 h; 900 °C for 15 h) in a muffle finance under an air atmosphere. Similar processes are applicable to the synthesis of Na_0.67Mn_0.9Al_0.05Zn_0.05O₂ (NMO1), Na_0.67Mn_0.85Al_0.1Zn_0.05O₂ (NMO2), Na_0.67Mn_0.85Al_0.05Zn_0.1O₂ (NMO3) or Na_0.67Mn_0.8Al_0.1Zn_0.1O₂ (NMO4) with the addition of Al₂O₃ and ZnO compounds.

2.2 Materials characterization

The morphologies are obtained by transmission electron microscopy (TEM, JEM-2100) and scanning electron microscopy (SEM, Zeiss Sigma 300) with an energy dispersive spectrometer (EDS) to observe element distribution. The crystal structures of these materials are analyzed by powder X-ray diffractor (XRD, X'Pert3 Powder, 5° min⁻¹; 10–80°). In situ XRD measurements are carried out by conducting electrochemical cycling and XRD synchronously in a cell mold (Beijing Scistar Technology Co. Ltd; 5° min⁻¹; 10–50°). The valence states of the elements are determined by X-ray photoelectron spectroscopy (XPS, Thermo Fisher Scientific K-Alpha). The exact stoichiometric ratio of elements in materials is investigated by using an inductively coupled plasma optical emission spectrometer (ICP-OES, PerkinElmer, Optima8000).

2.3 Electrochemical measurements

The active material, carbon black (Super-P), and binder (polyvinylidene fluoride; PVDF) are dispersed in N-methyl pyrrolidone (NMP) solvent in a weight ratio of 80:10:10 to get a suspension. The typical mass loading of the active cathode materials is controlled to be about 1.5 mg cm⁻². Subsequently, the suspension is cast onto aluminum foil and dried at 70 °C for 12 h in a vacuum electric oven. The electrode is cut into circular electrode pieces with a diameter of 12 mm. The 2032-coin-cells are assembled in an argon-filled glove box for evaluating the electrochemical performance. Sodium metal is used as both the counter and reference electrodes in half-cells. A Whatman glass microfiber filter (Grade GF/F) is used as the separator and 1.0 M NaClO₄ in propylene carbonate (PC) with 5 vol% of fluoroethylene carbonate (FEC) is used as the electrolyte. The full-cells are well-assembled with NMO3 and commercial hard carbon (HC) as the cathode and anode, respectively, denoted as NMO3//HC. The HC electrode includes HC, PVDF and Super-P (80:10:10, wt%) dispersed in NMP solvent. The typical mass ratio of the cathode/anode is carefully maintained as 7:4 based on the specific capacities of cathode and anode electrodes at a low current density of 20 mA g⁻¹. In order to eliminate the existing defect behavior and improve the initial coulombic efficiency (ICE), the positive and negative electrodes were pre-treated in a half cell with three cycles. The electrolyte and separator were the same as those for the half-cell. Charge/discharge tests and intermittent constant current titration technique (GITT) tests were performed on a 2.0–4.2 V battery test system (NEWARE CT4008). Cyclic voltammetry (CV) and electrochemical impedance spectroscopy (EIS) were carried out using a CHI660E electrochemical station. Simultaneous charge/discharge and EIS tests were used to record EIS curves in the field.

3. Results and discussion

The schematic diagram of structural evolution after Al³⁺/Zn²⁺ co-doping is displayed in Fig. 1a. The P2-Na_0.67MnO₂ cathode possesses a typical layered structure attributed to the P6₃/mmc space group (PDF: #27-0751) with an ABBA oxygen stacking sequence and Mn–O-layers composed of MnO₆ octahedral units shared by the edges.²⁴ Besides, Al³⁺ and Zn²⁺ are introduced into the Mn sites to adjust the spatial and electronic structures. In the XRD patterns (Fig. 1b), the (002), (004), (100), (102) and (103) diffraction peaks located at 15.8°, 32°, 36.03°, 39.6° and 43.6°, respectively, can be indexed to the typical hexagonal P₂-phase structure without any impurities.²⁵ The highly trenchant diffraction peaks also indicate high crystallinity of these samples. Moreover, the diffraction peaks show a trend of first shifting towards smaller 2θ angles from NMO to NMO1, NMO2 and NMO4. Then, the diffraction peaks slightly shift in the reverse direction for the NMO3 sample. The Rietveld refinement patterns (Fig. 1c, d and S1†) more clearly confirm that NMO3 (a = b = 2.88002 Å, c = 11.2132 Å, and V = 80.547333 ų) has the largest cell parameters and volumes compared to the others (Table S1†).

In general, Al³⁺/Zn²⁺ co-doping can expand the interplanar spacing, thereby promoting the diffusion kinetics of Na⁺. Furthermore, the valence state of elements is determined by XPS measurements (Fig. 1e–g). The binding energies located at 1043.8 and 1020.8 eV are ascribed to the 2p_1/2 and 2p_2/3 of Zn²⁺ (Fig. 1e), respectively. The peaks situated at 643.6 and 654.5 eV and 641.7 and 653.6 eV are attributed to the 2p_1/2 and 2p_2/3 of Mn³⁺ and Mn⁴⁺ (Fig. 1f), respectively. This proves the existence of mixed valence states of Mn^3+/4+ in the NMO3 sample. Besides, the binding energies located at 74.18 and 73.18 eV are assigned to the 2p_1/2 and 2p_2/3 of Al³⁺ (Fig. 1g). Hence, the valence states of Mn, Zn and Al are +3/+4, +2 and +3 in the NMO3 sample.

To explore the morphology of these P2-phase materials, SEM and TEM images are obtained (Fig. 2a, b and S2†). The NMO3 material presents a regular hexagonal lamella with diameter of 3 μm as shown in the SEM image (Fig. 2a). The thickness of such a lamella is measured from the cross section to be 1 μm from the TEM image (Fig. 2b). The HRTEM image displays a lattice spacing of 0.207 nm that corresponds to the (103) crystal plane in the P2-phase NMO3 layered structure²⁶ (Fig. 2c). The SAED patterns clearly exhibit the (010), (100) and (110) lattice planes in the hexagonal P6₃/mmc space group, revealing excellent crystallinity (Fig. 2d). Moreover, the EDS-mapping images indicate the uniform distribution of Na, Mn, O, Al and Zn elements in the lamellar NMO3 material (Fig. 2e–i). The sparse bright spots of Al and Zn elements are due to the minor doping content compared to Na, Mn, and O elements.

The electrochemical properties in half-cells are tested during 2.0–4.2 V as illustrated in Fig. 3a–e. The NMO3 sample is capable of delivering an initial discharge capacity of 181.85 mA h g⁻¹ (Fig. 3a), higher than that of the others (55.9, 39.25, 18.25, and 12.3 mA h g⁻¹) (Fig. S3†). The redox peaks in the 1st, 2nd and 3rd CV curves (Fig. 3b and S4†) are consistent with the charge–discharge voltage plateaus. In the 1st charge process, the oxidation peaks at about 2.33 V and 3.094 V can be ascribed to the Mn³⁺ → Mn⁴⁺ redox, along with the extraction of Na⁺ from interlayers. Subsequently, in the 1st discharge process, the reduction peaks located at 2.0 V and 3.023 V are assigned to the Mn⁴⁺ → Mn³⁺ redox,^10,11 accompanied by the Na⁺ insertion process. These redox polarization values in CV curves for NMO3 (0.33 V) are far smaller than those of NMO (0.446 V), NMO1 (0.44 V), NMO2 (0.437 V) and NMO4 (0.429 V), respectively, revealing much better electrochemical reversibility. Besides, the modified NMO3 sample displays superior rate capabilities (e.g. 124.64 mA h g⁻¹@1C; 94.87 mA h g⁻¹@5C) compared with the others (Fig. 3c). The cycling performances at 0.1C are demonstrated in Fig. 3d. The NMO3 material shows the most competitive cycling stability (89%@181.85 mA h g⁻¹@100 cycles@0.1C) as expected. Furthermore, to investigate their long-term cycling abilities, 600 charge–discharge cycles at 5C are carried out (Fig. 3e). The NMO3 cathode can deliver a capacity of 94.87 mA h g⁻¹ initially with 83% capacity retention during 600 cycles. To the best of our knowledge, such an NMO3 material is superior to those in selected latest reported studies (Fig. 3f). On increasing the loading mass to 5 mg cm⁻² as shown in Fig. S7,† the NMO3 sample can still deliver a superior initial discharge capacity of 179.29 mA h g⁻¹, rate capabilities (e.g. 123.71 mA h g⁻¹@1C; 90.95 mA h g⁻¹@5C), and cycling stability (93%@179.29 mA h g⁻¹@30 cycles@0.1C), thereby showing good application potential. Moreover, the D_Na⁺ of these samples is evaluated by GITT tests (Fig. 3g, h and S5†). The D_Na⁺ of NMO3 ranges from 10^−9.6 to 10^−8.1 during a cycle, while they are 10^−10.5 to 10⁻⁹ for the NMO sample, respectively. As shown in Fig. 3g, h and S5,† the calculated Na⁺ diffusion coefficients (D_Na⁺) of NMO1, NMO2, NMO3, and NMO4 samples vary from 10^−10.4 to 10⁻⁹, 10^−10.3 to 10^−8.4, 10^−9.6 to 10^−8.1 and 10^−9.9 to 10^−8.1, respectively. Hence, co-doping of Zn²⁺ and Al³⁺ cations indeed boost the D_Na⁺ compared to that of bare NMO, and the highest D_Na⁺ of NMO3 indicates the fastest reaction kinetics so as to offer superior rate capability.

In order to clarify the structural evolution process behind the improved electrochemical performance of the NMO3 sample, in situ XRD patterns are obtained (Fig. 4a and b). In the charge process (2.0–4.2 V), the diffraction peaks of (002), (100), (102), (104) and (103) gradually shift towards smaller 2θ degrees by about 0.5° (Fig. 4a), respectively, accompanied by Na⁺ extraction from the Mn–O interlayers and lattice expansion. Subsequently, these diffraction peaks consistently move towards a larger 2θ angle and return to their original positions, along with Na⁺ insertion into the Mn–O interlayers and lattice contraction. Hence, derived from such Zn and Al co-doping, highly symmetrical diffraction peaks during a cycle symbolize excellent electrochemical reversibility, indicating an effectively suppressed P2 → O2 phase transition. Furthermore, for investigating the evolution of electronic conductivity in a complete electrochemical cycle, in situ EIS curves are obtained (Fig. 4b–e). In the 2nd cycle, the R_ct values progressively reduce first (charging-Na⁺ extraction-lattice expansion) and then increase later (discharging-Na⁺ insertion-lattice contraction) for the NMO3 sample (Fig. 4c), which is consistent with in situ XRD analysis (Fig. 4a). Besides, the R_ct for the NMO sample also undergoes the same trend of change in the 2nd cycle except for larger values than those of NMO3 (Fig. 4b). Moreover, the R_ct maintains the same pattern of change as before with increased values for both NMO and NMO3 samples in the 600th cycle (Fig. 4d and e). Meanwhile, it can be clearly observed that the R_ct values of the NMO3 sample are always smaller than that for the NMO sample in both the 2nd and 600th cycles (Fig. 4f and g). Such results profoundly clarify the strengthened Na-insertion/extraction kinetics which originated from a larger lattice, more stable structure, and faster ionic/electronic conductivities of the NMO3 material. The ICP-OES result of the electrolyte (1.0 M NaClO₄ in PC with 5 vol% FEC) after 600 cycles has displayed that the content of the Mn element is 0.0084 mol, which means the dissolution of Mn²⁺ (disproportionation reaction: 2Mn³⁺ ↔ Mn⁴⁺ + Mn²⁺) due to the Jahn–Teller effect.^10,11 As a result, there is about 1% Mn dissolution in the electrolyte that deteriorates the cycling performance of the battery, but to a small extent due to the co-doping of Al/Zn.

To explore the evolution of the morphology and structure after long-term cycles, ex situ XRD and SEM are carried out as shown in Fig. 5. The diffraction peaks of (002), (004), (102), (103) and (104) become slightly weaker than the pristine for the NMO3 sample, while they become very weak, and some even disappear for the NMO sample compared with the fresh sample (Fig. 5a). This suggests the much better structural stability for modified NMO3 than bare NMO. Moreover, the morphological evolution process also displays smaller cracks of bulk particles during long-term cycling (2nd → 100th → 600th) for NMO3 than for NMO (Fig. 5b). Besides, EDS-mapping images (Fig. 5c–f and S6†) and EDS line-scan curves (Fig. 5g) also confirm that the Na, Mn, Zn and Al elements are still uniformly distributed on the surface of NMO3 particles, rather than excessive migration or dissolution in electrolyte, revealing inhibited Mn³⁺ migration and phase transitions. Thus, the structural collaborative doping modification is conducive to both microscopic and macroscopic stabilities in long-term cycling.

To evaluate the practicality of the as-designed NMO3 cathode, NMO3//HC full-cells are assembled as shown in Fig. 6a–e. Fig. 6a shows the operation schematic diagram of full-cells. In detail, NMO3//HC is able to exhibit a competitive working capacity of 165.89 mA h g⁻¹ with an energy density of 334.5 W h kg⁻¹ (Fig. 6b). The energy density of the NMO3//HC full-cell is 331.48 W h kg⁻¹ taking the amount of electrolyte (60 mg) into account. Moreover, it can display reliable cycling stability (84%@140.54 mA h g⁻¹@50 cycles@0.1C) as well as superior rate capabilities (86.38 mA h g⁻¹@2C) as illustrated in Fig. 6c–e. The unexceptional cycling stability is due to the mediocre cycling performance (94.7%@244.58 mA h g⁻¹@50 cycles@0.1C) of the HC anode as shown in Fig. S8,† and we will improve this issue in the future. Hence, such an NMO3 cathode has great potential for commercial application.

Furthermore, for clarifying the structure–activity relationship behind such synergistic doping, DFT calculations are carried out (Fig. 6f–i). Obviously, discrete islands comprising Na 1s, Mn 3d and O 2p orbitals are detected surrounding the Fermi level for all the samples that show characteristics of conductors due to the overlap of conduction bands and valence bands.³² In contrast, two additional discrete islands for Al and Zn 3p bonds are situated between the highest occupied molecular orbital (HOMO) and the lowest unoccupied molecular orbital (LUMO). Meanwhile, DOS of NMO3 (Fig. 6i) shows the largest electron cloud density located near the Fermi level than that of the NMO (Fig. 6f), NMZO (Fig. 6g) and NMAO (Fig. 6h) materials. It must be noted that the enhanced phase conductivity after Al³⁺ and Zn²⁺ doping is consistent with the evolution of R_ct values as shown in Fig. 4b–e, thereby manifesting the strongest reaction kinetics in both the bulk and interface. Thus, the introduction of Al³⁺ and Zn²⁺ elements into the structure can supply more valence bands and conduction bands with contiguous quantized energies. The free electron transition becomes much easier, thereby exhibiting better conductivity.³³

4. Conclusion

In summary, a Zn/Al co-doped P2-Na_0.67Mn_0.85Zn_0.05Al_0.1O₂ cathode is well-designed for the first time. Benefiting from the synergistic effect, the optimum NMO3 sample is capable of delivering much enhanced working capacity, rate capabilities, and cycling stability (83%@93.87 mA h g⁻¹@600 cycles@5C) in half-cells as well as energy density (334.5 W h kg⁻¹) in NMO3//HC full-cells. Moreover, the optimized structural reversibility, suppressed P2 → O2 phase transition and enhanced electrochemical reaction kinetics are profoundly clarified by a sequence of in situ and ex situ characterization methods and DFT calculations. Therefore, such innovative design can stimulate the thinking of researchers, such as adopting strategies of multi-doping into Na-, Mn- and O-sites, respectively, and (or) surface coating simultaneously to form more efficient synergistic effects, thereby further enhancing the comprehensive electrochemical performance of cathode materials for advanced SIBs.

Data availability

All relevant data are within the manuscript and its additional files.

Author contributions

Xiang Ding: conceptualization, methodology, writing – review & editing, supervision, funding acquisition. Cuixian Hu: conceptualization, methodology, writing. Yong Fan: conceptualization, methodology, writing. Yue Lin: methodology. Jinyang Liu: methodology. Yibing Yang: methodology. Liangwei Liu: methodology. Haibao Ma: methodology. Yi Xiao: methodology. Lili Han: conceptualization, methodology, supervision, funding acquisition.

Conflicts of interest

There are no conflicts to declare.

Acknowledgements

Xiang Ding, Cuixian Hu and Yong Fan contributed equally to this work. This work was financially supported by the National Key Research and Development Program of China (2022YFA1505700 and 2019YFA0210403), the National Natural Science Foundation of China (52102216), the Natural Science Foundation of Fujian Province (2022J01625 and 2022-S-002) and the Innovation Training Program for College Students (202310394020, cxxl-2023097, cxxl-2024131, and cxxl-2024136). The authors also greatly appreciate support by the Transmission Electron Microscope Platform and High-performance Computing Platform of the Fujian Science & Technology Innovation Laboratory for Optoelectronic Information of China.

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Footnote

† Electronic supplementary information (ESI) available. See DOI: https://doi.org/10.1039/d4ta04224g

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