Na⁺ Doping Modulates Electrochemical Behaviors of Cathode Materials in Aqueous Zinc‐Ion Batteries

Abstract

With the development of the industry, traditional lithium‐ion battery (LIB) technology is constantly improving, and new technologies are constantly breaking through, consolidating its dominant position in energy storage systems. However, the high manufacturing cost and inherent risks of organic electrolytes severely hinder the commercialization of LIBs. Therefore, developing low‐cost and high‐safety battery energy storage systems has become an urgent societal task. Research has shown that aqueous zinc‐ion batteries (AZIBs) using high‐safety and eco‐friendly aqueous electrolytes effectively reduce the risk. In this work, a facile one‐step hydrothermal approach was applied to prepare nanoflower‐structured Na‐NH₄V₄O₁₀ cathode materials. This material exhibited a specific capacity of 619.22 mAh g⁻¹ at a current density of 0.1 A g⁻¹, along with a capacity of 128.84 mAh g⁻¹ when the current density rose to 5 A g⁻¹. Meanwhile, a specific capacity of 245.63 mAh g⁻¹ was still retained even after 800 cycles at 1 A g⁻¹. This research provides valuable insights into the fabrication of cathode materials targeted at AZIBs in subsequent studies.

This figure illustrates the preparation process of Na‐NH₄V₄O₁₀. Its morphological structure presents a nanoflower shape, and the rate performance at different current densities is displayed in the lower right. The storage mechanism is the coordinated insertion of Zn²⁺/H⁺.

1. Introduction

In the field of rechargeable energy systems, aqueous zinc‐ion batteries (AZIBs) have attracted widespread attention as one of the most prospective energy storage systems in contemporary research. This is attributed to their environmentally friendly aqueous electrolytes, abundant global zinc reserves, excellent electrochemical stability, and relatively lower redox potentials of zinc metal anodes. However, they also have certain limitations, which have hindered the development and application of AZIBs. On one hand, the formation and growth of zinc dendrites may penetrate the separator membrane, leading to battery short circuits. On the other hand, the dissolution of zinc metal in aqueous electrolytes tends to trigger hydrogen evolution and corrosion reactions, which not only reduce the battery capacity but also impair the electrochemical performance of the battery [1].

The prevalent cathode materials for AZIBs encompass Prussian blue analogs (PBAs), vanadium‐based compounds, manganese‐based compounds, organic materials, and covalent organic frameworks (COFs). PBAs exhibit structural stability, yet their high crystalline water content, low electrical conductivity, and interfacial side reactions severely limit their practical applications. Manganese‐based materials such as manganese dioxide are characterized by low cost and eco‐friendliness. However, they suffer from Jahn–Teller distortion, which triggers the collapse of crystal structure of the material and results in battery capacity decay and shortened cycle life [2, 3]. Organic materials have low cost and tunable structures, but they have low operating voltage and poor conductivity. Covalent organic frameworks possess a large specific surface area and porous structure, but their complex synthesis process and poor electrical conductivity restrict their utilization in battery systems.

By contrast, vanadium‐based compounds stand out as one of the most promising cathode candidates for AZIBs, benefiting from the multivalent states inherent to vanadium. These materials feature variable valence states and can enable multielectron transfer reactions, typically delivering high specific capacities. They can be classified into vanadium oxides, metal vanadates, vanadium phosphates, and other related types. Despite these notable advantages, vanadium‐based materials still face critical bottlenecks. (1) Dissolution and phase transition issues: vanadium ions tend to leach into aqueous electrolytes, inducing structural degradation. (2) Slow reaction kinetics: the strong electrostatic interaction between Zn²⁺ and the material lattice results in sluggish ion diffusion rates. To address these challenges, researchers have adopted strategies such as interface modification, ions/crystal water doping, and carbon‐nitrogen coating to modify battery cathode materials, which effectively boost the electrochemical performance of AZIBs.

In 2021, Li et al. developed the cathode material K₂V₆O₁₆·1·5H₂O, which had a specific capacity of 261 mAh g⁻¹ in a ZnSO₄ electrolyte at 0.2 A g⁻¹ [4]. In 2022, Zhou et al. fabricated the polyoxovanadate material K₄Na₂V₁₀O₂₈, whose reversible capacity retained 152.3 mAh g⁻¹ after 100 cycles at 0.1 A g⁻¹ [5]. Notably, its capacity remained almost unchanged even after 1000 cycles at 1 A g⁻¹. In 2023, Guo et al. successfully prepared a self‐supporting amorphous calcium‐doped V₂O₅ (a‐Ca‐V₂O₅) cathode for wearable AZIBs [6]. The introduction of Ca significantly reduced the formation energy of VO₂, enabling a reversible amorphous‐crystalline phase transition during battery charge–discharge processes and thereby boosting the capacity of the battery. Between 2023 and 2025, the Mu team published a series of papers [7, 8, 9, 10, 11] proposing solutions to address zinc dendrite growth and zinc dissolution, providing significant guidance for subsequent performance research on AZIBs. During the period of 2025–2026, the Wu research group conducted a series of modification studies on vanadium‐based cathode materials for aqueous zinc‐ion batteries [12, 13, 14, 15, 16]. They employed various modification methods, such as calcium ion doping, molybdenum element modification, and metal‐organic framework composite anchoring, significantly enhancing the electrochemical zinc storage performance of vanadium‐based materials. Not only did they clarify the relevant modification mechanisms, but also they laid a solid theoretical and practical foundation for the subsequent exploration of vanadium‐based electrode materials.

Research findings reveal that vanadium‐based oxides feature complex crystal structures, flexible oxygen coordination modes, and a tendency to form multicomponent doped, layered, or tunnel‐like frameworks. Their crystal structures vary with oxidation states and preparation conditions, endowing them with excellent electrochemical properties. In this work, we synthesized nanoflower‐shaped Na–NH₄V₄O₁₀ via a facile one‐step hydrothermal approach and studied how alkali metal ions intercalating into the structure affect electrode performance. It was found that partial substitution of monovalent NH₄ ⁺ in the lattice by Na⁺ achieved a theoretical charge balance. The smaller ionic radius of Na⁺ (r ≈ 102 pm) relative to NH₄ ⁺ (r ≈ 143 pm) induced slight lattice distortion, which reduced the bond energy of V─O bonds. This facilitated the escape of lattice oxygen and the subsequent formation of oxygen vacancies. These vacancies can provide additional electron transport channels and improve electrical conductivity. At the same time, pre‐embedded Na⁺ mitigated the strong electrostatic interaction during the Zn²⁺ desorption and intercalation, leading to an enhanced ionic diffusion rate for the material as well as excellent rate capability and long‐term cycling stability [17].

2. Results and Discussion

The synthesis diagram of Na–NH₄V₄O₁₀ is shown in Figure 1, and this material was prepared via a simple one‐step hydrothermal method. Figure 2a and 2b display the scanning electron microscopy (SEM) images of NH₄V₄O₁₀. Figure 2d and 2e show the SEM images of Na–NH₄V₄O₁₀. We can observe that samples both present a nanoflower‐like structure composed of secondary needle‐or plate‐like whiskers. NH₄V₄O₁₀ tends to grow more regularly, while Na–NH₄V₄O₁₀ shows curled and stacked morphology. This may be due to the doping of Na⁺, which causes lattice distortion, altering the anisotropy of crystal growth, resulting in larger thin sheet sizes of the nanoflower‐like structure, thereby increasing the specific surface area of the material, offering extra active sites, and optimizing the conductivity of the battery system. Figure 2c is the transmission electron microscopy (TEM) image of Na‐NH₄V₄O₁₀, from which it can be seen that its structure is relaxed. Figure 2f is the high‐resolution transmission electron microscopy (HRTEM) image of Na‐NH₄V₄O₁₀, and the lattice spacing of (−314) crystal plane is 0.263 nm. Figure 2g displays the energy dispersive spectroscopy (EDS) spectrum of Na‐NH₄V₄O₁₀. Figure 2h to 2k are the elemental distribution maps of Na‐NH₄V₄O₁₀; it can be seen that Na⁺ ions have been successfully doped, and the distribution of various elements is uniform.

FIGURE 1.

Illustration of the preparation process of Na‐NH₄V₄O₁₀.

FIGURE 2.

The SEM images of (a,b) NH₄V₄O₁₀ and (d,e) Na‐NH₄V₄O₁₀. Panels (c,f) The TEM image and HRTEM image of Na‐NH₄V₄O₁₀, respectively. (g–k) The elemental mapping images of Na‐NH₄V₄O₁₀, corresponding to total elements (g), N (h), Na (i), O (j), and V (k).

Figure 3a presents a rate error bar chart of NH₄V₄O₁₀ and Na–‍NH₄V₄O₁₀ at different current densities, which also indicates that Na‐NH₄V₄O₁₀ has better rate performance than Na–NH₄V₄O₁₀. Specifically, Na‐NH₄V₄O₁₀ exhibits reversible capacities of 619.22, 433.62, 325.92, 243.18, 188.63, and 128.84 mAh g⁻¹ at current densities of 0.1, 0.2, 0.5, 1, 2, and 5 A g⁻¹, respectively. When the current density is switched back from 5 to 0.1 A g⁻¹, the capacity of Na–NH₄V₄O₁₀ still retains a specific capacity of 334.37 mAh g⁻¹, while that of NH₄V₄O₁₀ is only 218.42 mAh g⁻¹. Figure 3d compares the long‐term cycling stabilities of Na–NH₄V₄O₁₀ and NH₄V₄O₁₀ at 1 A g⁻¹. The initial capacity decay of layered vanadium‐based materials is essentially caused by the structural reconfiguration triggered by the ion insertion/deinsertion at the beginning of the cycle. When Zn²⁺ is rapidly inserted/deinserted between layers under high current density, it will destroy the layered framework of the original NH₄V₄O₁₀, leading to irreversible lattice distortion and a decrease in layer spacing, resulting in dead zinc sites [18, 19, 20]. After 800 cycles, Na–NH₄V₄O₁₀ maintains a reversible specific capacity of 245.63 mAh g⁻¹, with a Coulombic efficiency approaching 100%. In contrast, NH₄V₄O₁₀ delivers a mere capacity of 132.58 mAh g⁻¹ after the same number of cycles, approximately half of that of Na‐NH₄V₄O₁₀. As shown in Figure S2, Na‐NH₄V₄O₁₀ delivers higher specific capacity than the pristine NH₄V₄O₁₀ over 150 cycles, which further confirms the excellent electrochemical performance of our Na‐doped material and greatly strengthens the reliability of our conclusions. These findings demonstrate that Na‐NH₄V₄O₁₀ possesses superior rate capability and long‐term cycling stability to NH₄V₄O₁₀ and even surpasses other vanadium‐based cathode materials reported in previous studies (Table S1). Figure 3b depicts the initial three cyclic voltammetry (CV) cycles of Na‐NH₄V₄O₁₀ at a scan rate of 0.1 mV s⁻¹, with the first cycle corresponding to the activation process of the material. After full activation, the subsequent two cycles overlap well, indicating relatively stable cycling performance of the Na‐NH₄V₄O₁₀ electrode. Figure 3c displays the initial three galvanostatic charge and discharge (GCD) curves of the Na‐NH₄V₄O₁₀ electrode at 1 A g⁻¹. In the initial discharge cycle, massive Na⁺ rapidly inserts into the layered Na‐NH₄V₄O₁₀ lattice, inducing transient lattice distortion and increased ion migration resistance. After structural rearrangement of the vanadium‐oxygen framework, the diffusion pathway for sodium ions is optimized, accompanied by gradual elevation of potential and continuous capacity release. The potential valley of the first cycle is considerably lower than that of the second and third cycles, implying irreversible polarization of the electrode. Moreover, the second and third cycles exhibit nearly identical profiles, indicating that the reaction kinetics achieve equilibrium after the first cycle and demonstrate the favorable reversibility of the electrode. Notably, the voltage plateaus in the GCD curves correspond well with those in the CV curves. Figure 3e presents the GCD curves of Na‐NH₄V₄O₁₀ at varying current densities. As the current density escalates from 0.1 to 5 A g⁻¹, the curve shapes and voltage plateaus remain well‐preserved, testifying to the low polarization and excellent reversibility of the electrode reaction. Figure S1 compares the third‐cycle CV curves of NH₄V₄O₁₀ and Na‐NH₄V₄O₁₀ at 0.1 mV s⁻¹. Evidently, the CV curve area of Na‐NH₄V₄O₁₀ is much larger than that of NH₄V₄O₁₀, indicating higher reversible capacity and stronger electrochemical activity of Na‐NH₄V₄O₁₀. This enhancement is mainly attributed to the Na⁺ doping, which expands the specific surface area, increases the number of active sites, accelerates the diffusion rate of Zn²⁺, and ultimately improves the overall battery performance.

FIGURE 3.

(a) The capacity comparison of NH₄V₄O₁₀ and Na‐NH₄V₄O₁₀ at different current densities. (b) The CV of the first three curves of Na‐NH₄V₄O₁₀ at a scan rate of 0.1 mV s⁻¹. (c) The first three charge–discharge profiles of Na‐NH₄V₄O₁₀ at a current density of 1 A g⁻¹. (d) The long‐term cycling performance of NH₄V₄O₁₀ and Na‐NH₄V₄O₁₀ at 1 A g⁻¹. (e) The GCD curves of Na‐NH₄V₄O₁₀ at various current densities.

Figure 4a shows the TEM images of Na‐NH₄V₄O₁₀ after 100 cycles at 1 A g⁻¹. From the figure, it can be seen that the morphology of the material changes from nanoflowers to more compact spherical shapes after cycling. From Figure 4b to 4f, it can be observed that the core vanadium oxide framework and element distribution remain stable, indicating that the doping of Na further enhances the structural stability. The uniform distribution of elements ensures continuous electrochemical activity. The uniform presence of Zn elements directly proves that Zn²⁺ successfully embeds into the material lattice during charging and discharging, and the electrode reaction continues. At the same time, it also proves that the doped Na elements do not fall off during cycling and maintain the structural stability.

FIGURE 4.

After 100 cycles of charging and discharging at a current density of 1 A g⁻¹, for Na‐NH₄V₄O₁₀ (a) TEM and (b–f) the elemental mapping images, corresponding to total elements (b), Na (c), O (d), V (e), and Zn (f).

The electrochemical kinetics of Na‐NH₄V₄O₁₀ were further analyzed via CV measurements at different scan rates [17]. As illustrated in Figure 5a, CV curves of Na‐NH₄V₄O₁₀ were recorded over a voltage range of 0.1–1.7 V at various scan rates, from which two distinct redox couples were clearly observed at 0.49/0.68 and 0.85/1.15 V, respectively. As the scan rate varied, anodic peaks shifted toward higher potentials, whereas the cathodic peaks moved toward more negative potentials, yet the general shapes of the CV curves stayed almost intact, which confirms that the prepared electrode possesses outstanding electrochemical reversibility [21, 22]. The correlation between the peak current (i) and the scan rate (v) can be expressed by the following equation [23]:

FIGURE 5.

(a) CV curves of Na‐NH₄V₄O₁₀ electrode at different scan rates. (b) Pseudocapacitive contribution of Na‐NH₄V₄O₁₀ at a scan rate of 1 mV s⁻¹. (c) Pseudocapacitive contribution ratios of Na‐NH₄V₄O₁₀ at various scan rates. (d) log (i)–log (v) plots of Na‐NH₄V₄O₁₀. (e) GITT curves. (f) Single‐step GITT curve. (g) Comparison of ion diffusion rates. (h) EIS plots. (i) N₂ adsorption–desorption isotherms and pore size distribution curves of Na‐NH₄V₄O₁₀.

$$i=a\hspace{0.17em}{v}^b$$ (1)

$$\log \hspace{0.17em}\left(i\right)=b\log \hspace{0.17em}\left(v\right)+\log \hspace{0.17em}\left(a\right)$$ (2)

where i represents the peak current, v signifies the scan rate, and a and b are both constant parameters. In detail, b refers to the slope obtained from the log (i) versus log (v) plot. A b approaching 0.5 implies that the electrochemical reaction is dominated by a diffusion‐controlled process, whereas a value near 1.0 indicates that the reaction is governed by surface capacitive behavior [24]. As presented in Figure 5d, the calculated b values of Na‐NH₄V₄O₁₀ at peaks 1, 2, 3, and 4 were 0.56, 0.69, 0.6, and 0.63, respectively. These results demonstrate that the Zn²⁺ storage mechanism of the Na‐NH₄V₄O₁₀ electrode is jointly regulated by ion diffusion and capacitive effects, with the former exerting a dominant influence [25]. The respective contributions of the capacitive and diffusion‐controlled processes at a given voltage can be further quantified by the following equation [26].

$$i\left(v\right)=k_{1}v+k_2{v}^{1/2}$$ (3)

$$i\left(v\right)/v^{1/2}=k_1{v}^{1/2}+k_2$$ (4)

where k ₁ and k ₂ denote the proportions of capacitive‐controlled and diffusion‐controlled capacities in the total capacity, respectively [27, 28]. As illustrated in Figure 5b, the pseudocapacitive contribution ratio of Na‐NH₄V₄O₁₀ can reach 65% when the scan rate is set at 1 mV s⁻¹. It is worth noting that this ratio rises steadily as the scan rate goes up [29]. Figure 5c further shows that the capacitive contribution ratios stand at 34%, 43%, 50%, 56%, and 61% corresponding to the scan rates of 0.1, 0.2, 0.4, 0.6, and 0.8 mV s⁻¹, respectively.

The galvanostatic intermittent titration technique (GITT) was further employed to investigate the reaction kinetics of the Na‐NH₄V₄O₁₀ electrode [30]. The initial cycles of Na‐NH₄V₄O₁₀ and NH₄V₄O₁₀ were measured with a current pulse of 0.1 A g⁻¹ for 10 min, and the obtained results are presented at the bottom of Figure 5e; the settings of the remaining parameters are all the same. Herein, the voltage profiles exhibit a jagged shape due to the relaxation periods after charging and discharging processes [31]. Figure 5f displays the local magnification of the GITT curve of Na‐NH₄V₄O₁₀. The diffusion coefficient of Zn²⁺ can be determined via the following equation [32]:

$${D_{\mathrm{Zn}}}^{2\text{+}}\text{ = }4L^2/\pi \tau {(\text{Δ}{E}_s/\text{Δ}{E}_{\tau })}^2$$ (5)

where L denotes the Zn²⁺ diffusion length, namely, the electrode thickness; τ is the duration of the galvanostatic current pulse (s); ΔE _s represents the steady‐state potential difference of the plateau step; and ΔE_τ denotes the total voltage variation during the galvanostatic current pulse. As shown in Figure 5g, Na‐NH₄V₄O₁₀ exhibits a faster ion diffusion rate, thus delivering superior Zn²⁺ storage performance to NH₄V₄O₁₀. Figure 5h presents the electrochemical impedance spectroscopy (EIS) plots of NH₄V₄O₁₀ and Na‐NH₄V₄O₁₀, which were collected and matched with the equivalent circuit model. The linear section at low frequencies is linked to Zn²⁺ migration, while the semicircle at high frequencies corresponds to the charge transfer resistance (R _ct) [33]. The R _ct value of Na‐NH₄V₄O₁₀ (130.8 Ω) is much lower than that of NH₄V₄O₁₀ (225.9 Ω), reflecting excellent charge transfer capability [34, 35]. The N₂ adsorption–desorption isotherms and pore size distribution curves of Na‐NH₄V₄O₁₀ are displayed in Figure 5i, while those of NH₄V₄O₁₀ are provided in Figure S3, and the specific surface area and average pore size data for the two samples are summarized in Table S2. The pore size distribution curves verify the existence of a mesoporous structure in Na‐NH₄V₄O₁₀. Comparative analysis reveals that the specific surface area of Na‐NH₄V₄O₁₀ (30.1464 m² g⁻¹) is much larger than that of NH₄V₄O₁₀ (16.2603 m² g⁻¹). This phenomenon can be ascribed to the morphological alterations of Na‐NH₄V₄O₁₀ after modification, which serve to expand the specific surface area, furnish more active sites, augment the contact interface between electrode material and electrolyte, and consequently confer improved electrochemical properties on the material [36].

Figure 6a compares the X‐ray power diffraction (XRD) patterns of samples before and after Na⁺ doping, and it can be seen that both Na‐NH₄V₄O₁₀ and NH₄V₄O₁₀ show good matching with the standard reference card (PDF#31−0075). The diffraction peaks located at 25.5°, 33.9°, 47.0°, 49.8°, and 61.3° are assigned to the (110), (−311), (−205), (020), and (−712) crystal planes, respectively, confirming that Na⁺ doping into the layered structure does not damage the original host lattice framework. However, the intensity of XRD peaks of the Na⁺‐doped sample decreases to a certain extent compared with that of NH₄V₄O₁₀, which is due to the introduction of lattice defects during the partial substitution of NH₄ ⁺ by Na⁺, thus forming Na‐based vanadium oxides with lower crystallinity. This is favorable for the generation of more active sites for Zn²⁺ storage [37].

FIGURE 6.

(a) XRD patterns. (b–e) XPS spectra of Na‐NH₄V₄O₁₀: (b) survey spectra, (c) N 1s, (d) V 2p. (e) Comparison of the O 1s spectra of NH₄V₄O₁₀ and Na‐NH₄V₄O₁₀. (f) EPR spectra of NH₄V₄O₁₀ and Na‐NH₄V₄O₁₀.

X‐ray photoelectron spectroscopy (XPS) measurements were performed to further characterize the chemical elemental composition and electronic valence states of Na‐NH₄V₄O₁₀ [38]. Figure 6b displays the survey XPS spectra of Na‐NH₄V₄O₁₀, where the characteristic signals of N, V, O, Na, and C elements are clearly detected, which aligns well with the findings from EDS mapping characterization [39]. For the XPS survey spectra of NH₄V₄O₁₀, refer to Figure S4. As illustrated in Figure 6c, the N 1s spectrum presents two distinct characteristic peaks located at 399.89 and 401.59 eV, which are respectively attributed to neutral nitrogen moieties (–NH–) and positively charged nitrogen species (–NH⁺–) [21]. As shown in Figure 6d, the V 2p spectrum can be split into four individual peaks: the peaks at 515.89 and 517.3 eV are assigned to the V 2p_3/2 orbital, whereas the other two peaks at 523.6 and 525.1 eV correspond to the V 2p_1/2 orbital [40]. Figure 6e shows the O 1s spectra of NH₄V₄O₁₀ and Na‐NH₄V₄O₁₀, which can be fitted into three signal peaks at 529.99, 531.64, and 533.08 eV, corresponding to V─O bonds, oxygen vacancies (O_v), and O─H bonds. The emergence of the peak at 531.64 eV points to the existence of oxygen vacancies [41]. Through calculation, it can be found that the oxygen vacancy ratio of NH₄V₄O₁₀ is 22.46% and that of Na‐NH₄V₄O₁₀ is 25.34%, which is higher than that of NH₄V₄O₁₀, indicating that the doping of Na⁺ indeed introduces oxygen vacancies. Figure 6f shows the comparison of electron paramagnetic resonance (EPR) spectra of Na‐NH₄V₄O₁₀ and NH₄V₄O₁₀. It can be seen from the figure that a paramagnetic signal appears at g = 1.96 in Na‐NH₄V₄O₁₀. The introduction of metal will cause the g of oxygen vacancies to shift [42]. The signal intensity of Na‐NH₄V₄O₁₀ is higher, indicating that Na doping increases the concentration of oxygen vacancies in the sample, effectively regulating the crystal structure and thereby enhancing the performance of the material.

The Zn²⁺ storage mechanism of Na‐NH₄V₄O₁₀//Zn batteries was further clarified via ex situ XRD and ex situ XPS characterization [43, 44]. Figure 7a presents the first‐cycle charge–discharge curve of the Na‐NH₄V₄O₁₀ electrode at a current density of 5 A g⁻¹, along with the XRD patterns of this electrode measured at different voltages. The characteristic peaks of NH₄V₄O₁₀ can be detected at all tested voltages, indicating a robust crystal structure with no obvious structural collapse during cycling. Notably, the strong diffraction peaks around 45° and 75° are derived from the stainless‐steel foil substrate [45]. Herein, “C” and “D” are abbreviations for charging and discharging processes, respectively. The discharging process is associated with Zn²⁺ intercalation into the crystal lattice, accompanied by the reduction of vanadium valence (namely, V⁵⁺ is reduced to V⁴⁺). A decrease in intensity of characteristic diffraction peaks is observed as the discharge voltage declines from 0.85 to 0.1 V, arising from lattice distortion induced by Zn²⁺ intercalation. When discharged to 0.38 V, three distinct Zn‐containing diffraction peaks emerge at around 20°, which can be well indexed to Zn₄SO₄(OH)₆·4H₂O (ZSH, JCPDS No. 44–0673) and Zn₃(OH)₂V₂O₇·2H₂O (ZVO, JCPDS No. 50–0570). These newly formed byproducts are generated by the synergistic coinsertion of H⁺ and Zn²⁺, attributed to water dissociation in the aqueous electrolyte, producing equimolar H⁺ and OH^‒. Since free H⁺ cannot stably exist in the electrolyte, spontaneous intercalation of H⁺ into the cathode lattice is essential to maintain charge neutrality, directly verifying the Zn²⁺/H⁺ cointercalation reaction mechanism. During the subsequent charging process, the aforementioned three Zn‐related diffraction peaks completely vanish at 1.7 V, which is due to the reversible extraction of Zn²⁺ and H⁺ as well as the dissolution of these by‐products [46]. The extracted Zn²⁺ then diffuses through the electrolyte and is electrochemically reduced and deposited as metallic Zn on the Zn anode. Meanwhile, the intensities of the characteristic diffraction peaks recover gradually from 0.87 to 1.23 V during charging, implying lattice restoration after Zn²⁺ deintercalation. For the (110) crystal plane, the diffraction peak shifts toward higher angles during discharge, caused by reduced lattice spacing induced by the electrostatic interaction between intercalated Zn²⁺ and V─O bonds. In contrast, the peak position restores to its initial location during charging, demonstrating gradual lattice spacing recovery with Zn²⁺ deintercalation. This phenomenon also confirms the highly reversible Zn²⁺ intercalation/deintercalation behavior of the Na‐NH₄V₄O₁₀ electrode [47, 48]. The well‐preserved lattice structure with negligible structural variation before and after cycling serves as a critical structural foundation for the superior electrochemical performance of Na‐NH₄V₄O₁₀//Zn batteries [49, 50].

FIGURE 7.

(a) The first‐cycle charge–discharge curve of the Na‐NH₄V₄O₁₀ electrode at 5 A g⁻¹ and the XRD patterns under different charge–discharge voltages. Ex situ XPS maps of (b) Zn 2p and (c) V 2p under different charging and discharging states.

Ex situ XPS characterization was further adopted to investigate the variations in elemental valence states during the reaction process. Figure 7b illustrates the peak position changes of Zn under different voltage states. In the fully discharged states, the Zn 2p peak intensity increases significantly, which is attributed to the formation of byproducts (ZSH and ZVO), proving the participation of H⁺ in the reaction. When charged to 1.7 V, the peak intensity declines owing to the dissolution of byproducts and the reversible extraction of Zn²⁺ [46], which is consistent with the results of ex situ XRD. Figure 7c depicts the changes in V 2p peak positions under diverse voltage states, with V⁵⁺ and V⁴⁺ being the primary species during charging. Upon discharging to 0.1 V, the peak intensity decreases, and V⁵⁺/V⁴⁺ are converted to V³⁺, with V⁵⁺ eventually disappearing. This is due to the reduction reaction that occurred during the insertion of Zn²⁺. When charged to 1.7 V, V³⁺ completely disappears, while V⁵⁺ appears and takes the dominant position. This proves the reversibility of Zn²⁺/H⁺ intercalation and deintercalation, which is a key factor contributing to the excellent performance of the battery.

The electronic properties of NH₄V₄O₁₀ and Na‐NH₄V₄O₁₀ were studied using density functional theory. Figure 8a and 8b show the total and partial density of states of NH₄V₄O₁₀ and Na‐NH₄V₄O₁₀, respectively. From the figures, it can be seen that both graphs do not have an energy gap. However, the electronic state density near the Fermi level of Na‐NH₄V₄O₁₀ is significantly higher (To obtain a more accurate density of states, a more precise K‐point of 6 × 9 × 9 was adopted). This indicates that the doping of Na effectively optimizes the electronic conductivity of the material and accelerates the electrochemical reaction kinetics. Figure 8c and 8d are schematic diagrams of the crystal structures of NH₄V₄O₁₀ and Na‐NH₄V₄O₁₀. The Na atoms were successfully embedded in the vanadium oxide layered framework, proving that the doping of Na does not damage the main crystal structure. Figure 8e and 8f are the difference change density diagrams of NH₄V₄O₁₀ and Na‐NH₄V₄O₁₀. In the pure phase system, the charge is mainly concentrated between the V─O chemical bonds. This demonstrates the strong covalent effect of the V─O bond. However, in Figure 8f, there is a significant charge dissipation around Na, and a large amount of charge is transferred to the surrounding vanadium oxide layer. This intuitively proves that Na acts as an electron donor and undergoes a strong charge transfer with the matrix, regulating the valence state of the vanadium atom and the overall electron distribution. Mechanistically, this explains the intrinsic reason for the improvement in electronic conductivity after doping with Na⁺. The bond lengths of V─O in Table S3 can also be used to prove that the effect of the V─O bond weakens after Na doping, which is consistent with the aforementioned results.

FIGURE 8.

Density functional theory (DFT) calculation results. Total density of states (DOS) and partial DOS of (a) NH₄V₄O₁₀ and (b) Na‐NH₄V₄O₁₀. The structure of (c) NH₄V₄O₁₀ and (d) Na‐NH₄V₄O₁₀. Charge density difference (CDD) of (e) NH₄V₄O₁₀ and (f) Na‐NH₄V₄O₁₀.

3. Conclusion

In summary, we synthesized a nanoflower‐shaped Na‐NH₄V₄O₁₀ electrode via a facile hydrothermal method, which exhibits outstanding electrochemical performance. Pre‐embedded Na⁺ weakens electrostatic interactions during the Zn²⁺ intercalation/deintercalation process, stabilizes the layered structure, facilitates the insertion and extraction of Zn²⁺, and thus improves the ion diffusion rate. Na⁺ doping causes slight lattice distortion and oxygen defects, providing more electron transfer pathways and enhancing electrical conductivity. While its highly reversible Zn²⁺/H⁺ costorage behavior is further confirmed via ex situ XRD and ex situ XPS characterization, providing valuable insights for subsequent investigations into AZIBs.

4. Experiment Section

4.1. Synthesis of NH₄V₄O₁₀ and Na‐NH₄V₄O₁₀

All chemical reagents used in this study were of analytical grade and utilized directly without further purification or modification. Detailed information about the chemicals is provided in the experimental section of the supporting information. The target sample was synthesized via a one‐step hydrothermal approach. Specifically, 0.32 g of NH₄V₄O₁₀ was dispersed into 40 mL of deionized (DI) water, after which the mixture was heated to 60°C and stirred thoroughly to achieve homogeneity. Subsequently, 0.58 g of C₂H₂O₄·2H₂O was incorporated into the aforementioned dispersion, with continuous stirring until a clear greenish solution was formed. Next, 0.0178 g of Na₂CO₃ (at a molar ratio of Na to V being 1:8) was introduced into the aforementioned clear solution and stirred until fully dissolved. The resulting homogeneous solution was transferred into a 50 mL polytetrafluoroethylene‐lined hydrothermal autoclave, and the reactor was maintained at 180°C for 3 h. After the reaction finished, the obtained precipitate was alternately rinsed three times with DI water and anhydrous ethanol and then dried in a vacuum oven at 60°C overnight. For the preparation of NH₄V₄O₁₀, the procedure was exactly the same as described earlier, except that Na₂CO₃ was not added in the final step.

4.2. Characterization

We characterized the surface morphology and crystalline structure of the as‐synthesized materials by means of SEM (ZEISS Sigma 300, Germany), TEM (JEOL JEM‐F 200, Japan), and HRTEM (JEOL JEM‐F 200, Japan). Moreover, XRD (Rigaku SmartLab SE, Japan) tests and XPS (Shimadzu/Krayos AXIS Ultra DLD, Japan) analyses were employed to investigate crystal structures formed and elemental valence states of the materials.

4.3. Electrochemical Performance Tests

A homogeneous mixture was prepared by thoroughly blending the synthesized precursor, conductive carbon black, and polyvinylidene fluoride at a mass ratio of 7:2:1. Subsequently, 0.375 g of N‐methylpyrrolidone was introduced into the mixture, and the resulting system was magnetically stirred for 3 h to form a uniform slurry. The slurry was cast onto 1 × 1 cm stainless steel foils and dried at 60°C in a convection oven, yielding cathode sheets with an average active material loading of 1.2 mg. For battery assembly, 2 M ZnSO₄ served as the electrolyte, and Zn foils were adopted as the anode. The battery was assembled following a specific sequence, after which its electrochemical performance was evaluated. CV and EIS tests were carried out on a CHI660E electrochemical workstation. Besides, GITT and GCD, along with assessments of rate capability and long‐term cycling stability, were tested via the NEWARE testing system.

4.4. Calculation Method

All spin‐polarized DFT calculations [51, 52] were implemented via the Vienna Ab initio Simulation Package (VASP). In this computational framework, plane‐wave basis sets combined with the projector augmented‐wave (PAW) method [53, 54] were utilized to treat core electrons. The Perdew–Burke–Ernzerhof (PBE) functional [55] within the generalized gradient approximation (GGA) was adopted to characterize exchange‐correlation interactions. Grimme’s DFT‐D3 dispersion correction method was included to accurately account for van der Waals interactions [56]. A uniform plane‐wave cutoff energy of 450 eV was set for all computational cases. For Brillouin‐zone sampling, a Γ‐centered Monkhorst–Pack k‐point grid with a size of 2 × 3 × 3 was applied. Considering the strong correlation characteristics of d‐orbit electrons in vanadium atoms, the DFT + U correction method was adopted, and the effective Hubbard U value was determined as 3.25 eV in accordance with previous studies [57]. All atomic positions were fully relaxed until the maximum force on each atom was less than 0.02 eV/Å, and the total energy was converged to within 10⁻⁵ eV. The electron density difference, $\Delta \rho$, is defined as $\Delta \rho ={\rho }_{\ast \mathrm{M}}-{\rho }_{\ast }-{\rho }_{\mathrm{M}}$, where ${\rho }_{\ast \mathrm{M}},{\rho }_{\ast }$, and ${\rho }_{\mathrm{M}}$ correspond to the electron densities of the surface with adsorbed M molecules, the bare isolated surface, and the isolated M molecule, respectively. Moreover, the doping formation energy, E_f, was calculated using the formula: $E_f=(E_2+E_{\mathrm{NH}4})-(E_1+E_{\mathrm{Na}})$, where $E_2$ represents the total energy of the Na‐doped structure, $E_1$ is the total energy of the pristine(undoped) structure, $E_{\mathrm{NH}4}$ is the energy of the NH₄ species under vacuum conditions, and $E_{\mathrm{Na}}$ is the energy of bulk Na.

Funding

This work was supported by the Horizontal Research Project Foundation of Shanxi Province (2510400002HX).

Conflicts of Interest

The authors declare no conflicts of interest.

Supporting information

Supplementary Material

Data Availability Statement

The data that support the findings of this study are available from the corresponding author upon reasonable request.