English

• Large-scale energy storage systems played an increasingly critical role in renewable energy. Lithium-ion batteries, which dominated the market, can no longer meet people's environmental protection and safety demands because of serious thermal runaway, toxic and flammable organic electrolytes, and difficult recovery strategies [1–4]. Aqueous batteries, such as sodium ion batteries, ammonium ion batteries, zinc ion batteries, and aluminum ion batteries, have entered people's views with the emergence of energy and environmental problems worldwide [5–8]. Aqueous zinc ion batteries (AZIBs) were favored by researchers because of their similar ion radius to Li⁺, excellent economic and environmental benefits, and superior energy storage characteristics [9–11]. The cathode materials are a key to the electrochemical performance of AZIBs, among which MnO₂ was widely used because of its long history, diverse crystal forms, and superior theoretical capacity. However, the poor conductivity of traditional MnO₂ materials was against the rapid transfer of ions and electrons. In addition, the slow desolvation process of hydrated Zn²⁺ and the strong coulomb interaction between Zn²⁺ and MnO₂ lead to the decreased diffusion ability of Zn²⁺. More notably, the Jahn−Teller effect was considered a nightmare for manganese-based materials. The Mn³⁺ in a high-spin state led to an asymmetric distribution of electrons inside the material and seriously damaged the structural stability of MnO₂ [12–14]. Therefore, improving the electrical conductivity of MnO₂, focusing on the diffusion ability of Zn²⁺ and suppressing Jahn−Teller distortion significantly improve the electrochemical and structural properties of MnO₂. Designing a cathode material with high capacity, fast reaction kinetics, and stable structure will be a key to developing high-performance AZIBs.

  To solve and optimize a series of problems of MnO₂, many strategies have been proposed, including structural design, composite materials, and ion doping, etc. [15–17]. Among them, the doping strategy has received more attention from researchers by controlling the cost and enabling the host material to produce specific electrochemical properties. The ions doped can effectively inhibit Jahn−Teller distortion and enhance structural stability. In addition, doping will introduce vacancies and change the crystallinity, which is beneficial for the adsorption/diffusion of ions. Moreover, the adjustment of the intrinsic electronic structure will also promote the improvement of electron migration ability. These advantages of doped strategies have been demonstrated in many previous studies [12,18–21]. Among them, high valence ions were favored because of their lower substitution energy. However, it was a pity, that the research effect of high-valence ion doping was undesirability, and there were generally problems of short life and low capacity [22]. Therefore, how to enhanced the improvement effect of high-valence ion doping on the lifespan and capacity of the host material was the focus of this research.

  Vanadium oxides have already become one of the widely used cathode materials for AZIBs. This is due to the fact that the multivalent transitions of vanadium ions can provide additional capacity [23,24]. Its excellent performance in AZIBs has provided inspiration for our work [25,26]. Therefore, in this paper, vanadium oxysulfate was used as the source of vanadium ions, and the MnO₂ samples doped with different proportions of vanadium were prepared by hydrothermal method (VMO-x). The experimental results and density functional theory (DFT) showed the vanadium ions can form V−O bonds with δ-MnO₂. This optimized the charge/ion state and enhanced the electrical conductivity and electrochemical reaction. More importantly, the Mn−O bonds were lengthened, and decreased oxidation state of the Mn because vanadium ions were doped, effectively alleviating the Jahn−Teller effect caused by Mn³⁺. The VMO-5 sample exhibited a maximum specific capacity of 649 mAh/g at a capacity of 283 mAh/g, and a capacity retention rate of 79.1% after 2000 cycles at a current density of 1 A/g. Finally, the rapid reaction kinetics of VMO-5 and the energy storage mechanism of Zn²⁺/H⁺ co-intercalation was demonstrated by cyclic voltammetry (CV), ex site X-ray diffraction (XRD), and galvanostatic intermittent titration technique (GITT). This work innovatively introduced high-valent vanadium ions into MnO₂ to improve its electrochemical performance. Compared with other related works, this study has made outstanding contributions in terms of the capacity and cycle life of MnO₂ materials. In addition, this work also provided new thinking for the strong attenuation of capacity under microcurrent.

  This work used vanadium oxysulfate as a dopant to prepare the VMO-x materials by hydrothermal method (Fig. S1 in Supporting information), and a series of characterization analyses were conducted on it. The XRD results displayed that the diffraction peaks of the prepared samples were consistent with the standard card JCPDS No. 80–1098, and no excess peaks were found (Fig. 1a), which demonstrated that the samples were attributed to the layered δ-MnO₂ phase. The diffraction peaks at 12.3°, 24.8°, 36.7°, and 65.6° were attributed to the (001), (002), (111), and (310) crystal planes, respectively. With the doping of vanadium ions, the characteristic peaks belonging to δ-MnO₂ still exist. In addition, the intensity of the peaks decreased. The width increased, representing the transformation of VMO-x samples to amorphous, and the average particle size decreased, which was conducive to improving material properties.

  Figure 1

  

  Figure 1.  (a) XRD of VMO-x. (b) FTIR spectra of VMO-x. (c) Raman spectra of VMO-x. (d, e) SEM of VMO-5 at 1 µm and 200 nm scales. (f) SEM–EDX of VMO-5.

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  The Fourier transform infrared spectroscopy (FTIR) results of MnO₂ and VMO-x samples are shown in Fig. 1b. All samples demonstrated similar characteristic peaks, mainly because of the small amount of vanadium doping. The distinctive peaks at 3428 cm⁻¹ and 1596 cm⁻¹ were attributed to the vibrations of O−H groups, indicating the presence of water molecules between layers of the sample [13]. In addition, the peaks of the VMO-x were significantly enhanced because the doping of vanadium ions increased its lattice spacing and facilitated the storage of water molecules. The diffraction peaks at 1350−1385 cm⁻¹ were attributed to the C−N bonds, mainly from the C and N in the air, consistent with the results of energy dispersive the X-ray (EDX) and XPS in the following [27–29]. The asymmetric stretching vibration of the V−O−V bonds was verified to be successfully doped by vanadium ions at 670 cm⁻¹, while the infrared absorption peaks at 619 and 784 cm⁻¹ were attributed to the stretching vibration of Mn−O in the MnO6 octahedron [14,30–32].

  To investigate the influence of vanadium ions doped on the structural properties of MnO₂, Raman spectra were presented within the range of 200−1000 cm⁻¹ (Fig. 1c). The doped vanadium changed the interlayer environment of MnO₂ and reduced the Mn−O bonds coordination at 570 and 640 cm⁻¹, which had a significant inhibitory effect on Jahn−Teller. Additionally, the lengthened the Mn−O bonds in VMO-5 indicated enhanced electrochemical performance [27,33]. Furthermore, weak stretching vibrations corresponding to the V−O−V and V=O bonds were identified at 719 and 930 cm⁻¹ respectively, these findings were consistent with the FTIR results [31,34]. In conclusion, the faint signals related to vanadium-related groups found by the FTIR and Raman spectroscopy confirmed the successful doping of trace vanadium.

  The morphologies and elements of MnO₂ and VMO-5 were analyzed by scanning electron microscope (SEM) and energy dispersive X-ray (EDX). As shown in Figs. 1d and e, the VMO-5 was a typical nanoflower structure composed of 13.2 nm nanosheets, which was similar to the structure of MnO₂ (Fig. S2 in Supporting information), indicating that the doping of vanadium ions did not destroy the MnO₂ structure. The EDX of MnO₂ and VMO-5 (Fig. S3 in Supporting information and Fig. 1f), the 0.13% vanadium content in VMO-5 demonstrated the success of the doping, this small increase validates the weak changes in the FTIR and Raman results. Notably, the significant decrease in Mn content because of the substitution of part of Mn with vanadium doping. In addition, the presence of additional C in MnO₂, perhaps as influenced by the air during preparation, also explains the discovery of C−N bonds in the FTIR spectrum in MnO₂.

  The chemical composition and valence states of the MnO₂ and VMO-5 were further analyzed using X-ray photoelectron spectroscopy (XPS). From Fig. 2a, the Mn, O, and C were found in the samples, and the vanadium was not detected in XPS because of the lower doping. Fig. 2b displayed the XPS spectrum of Mn 2p, the peaks of Mn 2p_3/2 and Mn 2p_1/2 were attributed to 642.3 and 654.2 eV, and the spin-splitting energy difference was 11.9 eV, which was concordance with the previous studies [27,35]. The enhanced energy difference of VMO-5 compared to MnO₂ (11.76 eV) indicated vanadium ion doping effectively adjusted the electronic structure of the host material and contributed to the rapid transport of ions. The two peaks of Mn 2p were deconvoluted into Mn⁴⁺ and Mn³⁺, and VMO-5 exhibited a new peak attributing to Mn²⁺ at 643.4 eV [36,37]. The Mn²⁺ demonstrated a decrease in the oxidation state of Mn, which led to enhanced structural stability and inhibited the Jahn−Teller effect of VMO-5 during charging and discharging. In addition, the vanadium ions doping led to lengthened Mn−O bonds and decreased electrostatic repulsion, which facilitated cation insertion and ion transfer kinetics, ultimately obtaining surprising capacity and multiplicity performances [38]. The O 1s spectra of MnO₂ and VMO-5 are shown in Fig. 2c, the convolution peaks of 530.0, 531.5, and 533.8 eV correspond to the Mn−O−Mn bonds, Mn−O−H bonds, and H−O−H bonds, respectively [35]. The introduction of the vanadium ions increased the lattice spacing of MnO₂ and increased the lattice water content, as demonstrated by the increase in the area of the H−O−H bonds (from 4.8% to 6.9%). The increased lattice water contributed to facilitating the diffusion of H⁺/Zn²⁺ and decreased the internal resistance of the cell, increasing the effective contact area of the electrode, and providing more electrochemical reaction interfaces [39]. The high resolution C 1s signal (Fig. 2d) showed three peaks attributed to 284.8, 286.4, and 288.8 eV, which correspond to the vibrations of C−C, C−O, and C=O bonds, respectively [40]. The C 1s spectrum of VMO-5 and MnO₂ were similar and did not undergo major alteration, which seemed to predict that the C in the samples originates from air.

  Figure 2

  

  Figure 2.  (a) XPS full spectrum of VMO-5 and MnO₂. (b) Mn 2p, (c) O 1s, and (d) C 1s XPS spectrum of VMO-5 and MnO₂.

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  The electrochemical performance of the MnO₂ and VMO-x was evaluated in a CR2032 coin cell using zinc foil as the anode and 2 mol/L ZnSO₄ + 0.2 mol/L MnSO₄ solution as the electrolyte with the working mechanism (Fig. 3a). To investigate the cycling stability and the Zn²⁺/H⁺ storage mechanism, the CV curves of the electrodes were cycled at a scan rate of 0.1 mV/s in the range of 0.8−1.8 V (Fig. 3b). The CV curves of the VMO-5 exhibited two pairs of redox peaks at 1.24 V/1.57 V and 1.37 V/1.61 V, corresponding to the two-step insertion/detachment of Zn²⁺ and H⁺ by a two-step insertion/deletion mechanisms [41]. Compared with MnO₂ (Fig. S4 in Supporting information), the doping of vanadium ions enabled VMO-5 to have smaller polarization and higher specific capacity. In addition, the increasing intensity of the reduction peaks implied an activation process, which decreased the potential difference between oxidation and reduction, making the charge/discharge process more reversible. The almost overlapping CV curves over the 5 cycles demonstrated the excellent reversibility performance of VMO-5. It was worth noting that the abnormality in the first cycle stems from the activation phenomenon of the battery. The galvanostatic charge/discharge (GCD) curves (Fig. 3c) coincided with the CV results, showing two charging/discharging plateaus. Notably, the charging/discharging plateau attributed to Zn²⁺ embedding/de-embedding disappears at a high current density, because H⁺ possesses a smaller ionic radius, a faster migration rate at high rates, and can be embedded/de-embedded easier than Zn²⁺ [42,43].

  Figure 3

  

  Figure 3.  (a) Schematic diagram of AZIBs with zinc foil as anode and VMO-x as the cathode. (b) CV curves of VMO-5 at a scanning speed of 0.1 mV/s. (c) GCD curves of VMO-5 at a current density of 0.1−5 A/g. (d) Rate performance of MnO₂ and VMO-x. (e) EIS spectra of MnO₂ and VMO-x. (f) Cyclic performance of MnO₂ and VMO-5 at a current density of 0.1 A/g. (g) Cyclic performance of MnO₂ and VMO-5 at a current density of 1 A/g. (h) Performance Ragone diagram of VMO-5 and other types of cathode materials for AZIBs.

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  Fig. 3d exhibited the comparison of the multiplicity performance of the MnO₂ and VMO-x samples when the current densities were increased from 0.1 A/g to 5 A/g. The VMO-5 exhibited specific capacities of 303.7, 264.2, 239.0, 225.5, 206.1, 166.3, and 124.5 mAh/g at the corresponding current densities, while the specific capacities of the MnO₂ were 181.5, 153.0, 126.5, 108.2, 94.3, 70.8, and 47.1 mAh/g, respectively. In contrast, the discharge-specific capacity of VMO-5 was significantly higher than MnO₂. Furthermore, when the current density was restored to 0.1 A/g, VMO-5 exhibited an enhanced specific capacity (397.6 mAh/g), which was higher than MnO₂ (190.3 mAh/g). These results indicated that the doping strategy of vanadium ions effectively inhibited the Jahn-Teller effect, enhanced the ion diffusion capacity, and produced enhanced multiplicity performance. The electrical conductivity of the MnO₂ and VMO samples were compared by EIS, and as shown in Fig. 3e, the charge transfer resistance of MnO₂ was 263 Ω. Vanadium ion doping optimized the electrical conductivity of MnO₂ by adjusting the electronic structure, and thus VMO-5 obtained a resistance of 126 Ω, which was much lower than MnO₂.

  Vanadium ions were easier to substitute manganese ions because of their high valence state, which effectively inhibited the disproportionation reaction of Mn³⁺, and the valence change between vanadium ions provided additional capacity for the electrode materials. As shown in Fig. 3f, the MnO₂ and VMO-5 samples exhibited specific capacities of 127 mAh/g and 184 mAh/g at 0.1 A/g, respectively. Specifically, the MnO₂ exhibited a slow increase in 500 cycles with a final capacity of 278 mAh/g, while the VMO-5 capacity was greatly enhanced (649 mAh/g). The capacity enhancement attributed to the structural rearrangement of MnO₂ during the charging and discharging process which allows the vanadium ions to be deeply doped from the surface to the interior of the material. Disappointingly, the high capacity of VMO-5 was instability and displayed a surprisingly severe decrease in the subsequent cycles, probably because of the structural aberration of the samples caused by the chain reaction of the secondary doping of vanadium ions during the slow charge/discharge process of microcurrent. In addition, the structural change during secondary doping results in the inability of the by-product ZSH to dissolve effectively, which gradually accumulates on the electrode surface, resulting in a sharp decrease in capacity [44]. Encouragingly, the stability of the VMO-5 was optimized after a brief 20 cycles of activation (0.1 A/g). The cycling performances of MnO₂ and VMO-5 at 1 A/g were shown in Fig. 3g, where the initial capacity of VMO-5 was 358 mAh/g, and the VMO-5 still retained of capacity 283 mAh/g after 2000 cycles, with a retention rate of 79.1%. The MnO₂ was short-circuited at 1300 cycles, and the relevant cycling performances of the remaining samples are shown in Fig. S5 (Supporting information). One possible reason was the short activation process resting for a period, which allowed the vanadium ions introduced into the interior of MnO₂ to be stabilized efficiently during charging and discharging. In conclusion, the vanadium ion doping strategy brought a surprisingly high capacity for MnO₂, which still outperforms some previous studies (Fig. 3h and Table S1 in Supporting information) despite the challenge of severe capacity retention at small currents [45–53].

  Flexible devices have special requirements for the flexibility and safety of batteries, so VMO-5 was used as the cathode to prepare soft-pack batteries to evaluate the flexibility and safety of samples. As shown in Fig. S6 (Supporting information), the LED lights can be lit when the softpack battery is in normal condition. Unexpectedly, when the soft pack battery was bent 90° and 180°, the LED lights showed higher brightness, which fully demonstrates the extraordinary flexibility of the VMO-5 material. In addition, when the soft pack battery is damaged, the LED lights still emit normally, and the battery does not fire and heat, the safety performance of VMO-5 material has been fully verified. The satisfactory performance of the VMO-5 material in soft-pack batteries demonstrates its potential in the field of flexible devices.

  To investigate the charge storage and reaction kinetics of MnO₂ and VMO-5 during charging and discharging, the CV curves were analyzed at different scan rates. Fig. S7a (Supporting information) exhibited the CV curve of VMO-5, and the two pairs of redox peaks correspond to the two-step insertion mechanism of the VMO-5. As the increased in scan rate, the oxidation peak located at 1.61 V gradually disappeared, which was attributed to the suppression of Zn²⁺ embedding at large scan rates, consistent with the results of GCD curves. In addition, the capacitive behavior of the VMO-5 samples during energy storage was evaluated by determining the b-value, which was calculated as Eq. S2 (Supporting information). When b = 0.5, the electrode response was controlled by diffusive behavior, while the capacitive contribution dominated when b = 1. The log(v)−log(i) curves of VMO-5 are shown in Fig. S7c (Supporting information), and the b values of Peak 1, Peak 3, and Peak 4 were 0.71213, 0.79507, and 0.64609, respectively, after linear fitting. These results demonstrated that the response of VMO-5 was controlled by a mixture of diffusive and capacitive, and the capacitive behavior dominated. To quantitatively analyze the contribution of diffusive and capacitive behaviors, Eq. S3 (Supporting information) was used. As shown in Fig. S7d (Supporting information), the capacitive contribution has been dominated in the range of 0.1−1 mV/s scan rate and the proportion of capacitive contribution increased with the increase of the scan rate, which was attributed to the fact that H⁺ with smaller ionic radius was more suitable for insertion/detachment in large currents than Zn²⁺, and this finding was consistent with the CV and GCD results.

  To further investigate the energy storage mechanism of the VMO-5 in AZIBs, we performed ex situ XRD characterization to clarify the structural evolution of the VMO-5 during charging and discharging. Fig. 4a demonstrated the phase evolution of the VMO-5 from the initial state to the first charging and discharging process. The intense peaks appearing at 26.5° and 54.6° were attributed to the graphite foil collector, and new diffraction peaks (marked by red boxes) were found at 16.3°, 21.3°, and 58.7° during charging to 1.4 V, and gradually increased in strength during subsequent processes, indicating the generation of the byproduct Zn₄SO₄(OH)₆·nH₂O (ZSH). The formation of ZSH was attributed to the consumption of H⁺ during the charging and discharging process resulting in the localized pH increase of the electrolyte [49]. In addition, the new peaks located at 32°−36° were attributed to Zn_xMnO(OH)₂ (ZMO), which was caused by Mn²⁺ insertion [54]. When redischarged to 0.8 V, the peaks of ZSH and ZMO disappeared, demonstrating the reversible oxidation/reduction of VMO-5. The specific chemical formulas of ZSH and ZMO were as follows:

  $ 2 \mathrm{Zn}^{2+}+\mathrm{SO}_4^{2-}+6 \mathrm{OH}^{-}+\mathrm{nH}_2 \mathrm{O} \leftrightarrow \mathrm{Zn}_2 \mathrm{SO}_4(\mathrm{OH})_6 \cdot \mathrm{nH}_2 \mathrm{O} $

  (1)

  $ \begin{aligned} & \mathrm{Zn}_2 \mathrm{SO}_4(\mathrm{OH})_6 \cdot \mathrm{nH}_2 \mathrm{O}+\mathrm{Mn}^{2+} \leftrightarrow \mathrm{Zn}_x \mathrm{MnO}(\mathrm{OH})_2+4 \mathrm{H}^{+}+\mathrm{SO}_4^{2-} \\ & \quad+(2-x) e^{-} \end{aligned} $

  (2)

  Figure 4

  

  Figure 4.  (a) Ex situ XRD patterns of VMO-5 cathode in different charge and discharge states. (b) GITT profiles and corresponding diffusion coefficient of H⁺/Zn²⁺ for the VMO-5 cathode.

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  The reaction kinetics of MnO₂ and VMO-x were further discussed via GITT. The ion diffusion coefficient for Zn²⁺ insertion/deletion (D_Zn²⁺ ) was expressed as Eq. 4 (Supporting information). The ion diffusion coefficients of the VMO-x and MnO₂ were calculated as in Fig. 4b and Fig. S9 (Supporting information), the ion diffusion coefficients of VMO-5 were 10⁻⁶ − 10⁻⁹ cm²/s, which were significantly enhanced. These results demonstrated the doping of vanadium ions effectively increased the conductivity and enhanced the ion diffusion of MnO₂. In addition, the ion diffusion coefficients exhibited two different stages during the discharge process, further demonstrating the H⁺/Zn²⁺ co-embedding/de-embedding mechanism. Notably, the diffusion coefficient of the first stage was higher than the second stage, consistent with the faster ion diffusion ability of H⁺. In addition, by comparing the GITT of MnO₂ and the VMO-x samples after vanadium ion doping (Fig. S9), it was found that vanadium ion doping can enhance the diffusion ability of Zn²⁺ during the discharging process, which is consistent with the results of Zn²⁺ adsorption energy in the theoretical calculations below.

  DFT was utilized to elucidate the mechanisms and origins of vanadium doping in enhancing the conductivity of MnO₂ while decreasing the Jahn−Teller effect associated with Mn³⁺. Figs. 5a and b exhibited the theoretical configuration diagrams for MnO₂ and VMO-5. The incorporation of vanadium ions preserved the structural integrity, consistent with observations from SEM. Figs. 5c−e depicted the adsorption energies for Zn²⁺ on both materials, VMO-5 demonstrates an adsorption energy of −1.54 eV, exceeding that of MnO₂ was −1.13 eV. This finding theoretically corroborated that vanadium ion doping enhanced the host material's capacity for Zn²⁺ adsorption, thereby facilitating ion embedding processes. This result is consistent with the data of vanadium ion doping improving the Zn²⁺ diffusion coefficient of MnO₂ in the previous text, reflecting the consistency between the theoretical model and the experimental results. As shown in Fig. 5f, the density of states (DOS) analysis provided insights into the conductivity characteristics. MnO₂ exhibited a significant band gap of 1.482 eV indicative of its limited intrinsic conductivity, whereas VMO-5 revealed a reduced band gap of 1.091 eV, signifying enhanced electrical conductivity aligned with electrochemical behavior findings. In addition, VMO-5 showed energy levels associated with vanadium ions near the Fermi level, which strongly proved that the introduction of high valence vanadium ions changes the DOS of MnO₂, thus significantly enhancing the electrical conductivity of MnO₂ and optimizing its electron transport properties. Differential charge density analyses further clarified variations in electron interactions (Figs. 5g and h), the yellow and blue regions represent the accumulation and dispersion of charges, respectively. The doping of vanadium ions affected the charge density distribution of MnO₂, and the charge around vanadium atoms was depleted and accumulated in other parts of MnO₂, which proved the strong interaction and riched charge transfer behavior between vanadium ions and MnO₂. In conclusion, DFT proved theoretically that doped vanadium ions can effectively enhance the adsorption capacity of MnO₂ to Zn²⁺, optimize its conductivity and charge transfer behavior, and contribute to the improvement of VMO-5 cathode performance. In addition, the DFT results are consistent with the experimental results, which well proves the correctness of the theoretical model.

  Figure 5

  

  Figure 5.  DFT calculation configuration diagram: (a) MnO₂, (b) VMO-5. (c) The adsorption energies of Zn²⁺ in MnO₂ and VMO-5. E_ads of Zn²⁺ in (d) MnO₂ and (e) VMO-5. (f) The density of states (DOS) of MnO₂ and VMO-5. The differential charge density distribution of Zn²⁺ in (g) MnO₂ and (h) VMO-5.

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  The Jahn−Teller distortion caused by Mn³⁺ had always been an obstacle to the high performance of MnO₂-based materials. Mn³⁺ in a high spin state had a very large magnetic moment, which led to the asymmetric distribution of electrons between the layers of MnO₂. Structurally, it was manifested as the elongation of the Mn−O bonds in the longitudinal direction and the shortening of the Mn−O bonds in the transverse direction, so that the linear MnO₂ arrangement elongated along the axial direction and produced Jahn−Teller distortion. Here, high valence vanadium ions were doped into MnO₂. The low substitution energy of vanadium was used to reduce the oxidation state of Mn, thereby effectively suppressing the Jahn−Teller effect and enhancing the structural stability of MnO₂. In addition, the multivalent transition of vanadium ions also brought additional capacity to MnO₂, making it exhibited unexpected results in terms of electrochemical performance. Therefore, the VMO-5 cathode demonstrated a long cycle life of 2000 cycles, an excellent specific capacity of 283 mAh/g, and a capacity retention rate of 79%. The DFT theoretical calculation proved that doped vanadium ions effectively enhanced the adsorption capacity of MnO₂ to Zn²⁺, and optimized its electrical conductivity, and charged transfer behavior. In addition, VMO-5 demonstrated excellent flexibility and safety, exhibiting great potential in flexible devices. This work will provide a feasible and effective path for the modification of layered MnO₂.

  Declaration of competing interest

  The authors declare that they have no known competing financial interests or personal relationships that could have appeared to influence the work reported in this paper.

  CRediT authorship contribution statement

  Le Li: Writing – original draft. Shaofeng Jia: Writing – review & editing. Shi Yue: Writing – review & editing. Yuanyuan Yang: Writing – review & editing. Chao Tan: Writing – review & editing. Conghui Wang: Writing – review & editing. Hengwei Qiu: Writing – review & editing. Yongqiang Ji: Writing – review & editing. Minghui Cao: Writing – review & editing. Zige Tai: Writing – review & editing. Dan Zhang: Writing – review & editing, Supervision.

  Acknowledgments

  This work was supported the National Key Research and Development Program of China (No. 2024YFA1409900), the National Natural Science Foundation of China (Nos. 62101296 and 52303335), the China Postdoctoral Science Foundation (Nos. 2021M702656 and 2023M730099), the Natural Science Foundation of Shaanxi Province (Nos. 2021JQ-756 and 2021M702656), the Graduate Innovation Fund of the School of Mechanical Engineering, Shaanxi University of Technology (No. SLGJX202404).

  Supplementary materials

  Supplementary material associated with this article can be found, in the online version, at doi:10.1016/j.cclet.2025.111009.

  
  
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  Figure 1  (a) XRD of VMO-x. (b) FTIR spectra of VMO-x. (c) Raman spectra of VMO-x. (d, e) SEM of VMO-5 at 1 µm and 200 nm scales. (f) SEM–EDX of VMO-5.

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  Figure 2  (a) XPS full spectrum of VMO-5 and MnO₂. (b) Mn 2p, (c) O 1s, and (d) C 1s XPS spectrum of VMO-5 and MnO₂.

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  Figure 3  (a) Schematic diagram of AZIBs with zinc foil as anode and VMO-x as the cathode. (b) CV curves of VMO-5 at a scanning speed of 0.1 mV/s. (c) GCD curves of VMO-5 at a current density of 0.1−5 A/g. (d) Rate performance of MnO₂ and VMO-x. (e) EIS spectra of MnO₂ and VMO-x. (f) Cyclic performance of MnO₂ and VMO-5 at a current density of 0.1 A/g. (g) Cyclic performance of MnO₂ and VMO-5 at a current density of 1 A/g. (h) Performance Ragone diagram of VMO-5 and other types of cathode materials for AZIBs.

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  Figure 4  (a) Ex situ XRD patterns of VMO-5 cathode in different charge and discharge states. (b) GITT profiles and corresponding diffusion coefficient of H⁺/Zn²⁺ for the VMO-5 cathode.

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  Figure 5  DFT calculation configuration diagram: (a) MnO₂, (b) VMO-5. (c) The adsorption energies of Zn²⁺ in MnO₂ and VMO-5. E_ads of Zn²⁺ in (d) MnO₂ and (e) VMO-5. (f) The density of states (DOS) of MnO₂ and VMO-5. The differential charge density distribution of Zn²⁺ in (g) MnO₂ and (h) VMO-5.

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