Tailoring electronic structure to enhance the ammonium-ion storage properties of VO₂ by molybdenum doping toward highly efficient aqueous ammonium-ion batteries

Abstract

Recently, research on ammonium-ion storage has gained widespread interest, and it is still a major problem and a popular research area to produce high-performance electrode materials for aqueous ammonium ion batteries (AAIBs). Herein, the electronic structure of tunnel-like vanadium dioxide (VO₂) is tailored by molybdenum doping (denoted as VO₂-Mo) to enhance ammonium-ion storage properties toward highly efficient AAIBs. VO₂-Mo with a unique nanobelt structure is designed and synthesized by adjusting the content of Mo via a facile hydrothermal method. Density functional theory (DFT) simulations and experimental data both demonstrate that molybdenum atoms in the VO₂ structure can improve mass transfer, speed up ion transport, and accelerate kinetics, showing boosted NH₄⁺-storage properties. With 2% Mo doping, at 0.1 A g⁻¹, VO₂-Mo exhibits a specific discharge capacity of around 370 mA h g⁻¹, surpassing VO₂ (232 mA h g⁻¹) and the vanadium oxide-based materials that have been reported for NH₄⁺-storage. After approximately 6000 successive charging and discharging cycles at 2 A g⁻¹, it essentially maintains the specific capacity of 140 mA h g⁻¹. Using VO₂-Mo, polyaniline (PANI) and 1 M (NH₄)₂SO₄ as the anode, cathode, and electrolyte, respectively, a VO₂-Mo//PANI full battery was further built, and at 0.2 A g⁻¹, it reached a specific discharge capacity of up to 232 mA h g⁻¹, surpassing the performances of the most state-of-the-art AAIBs. At 89 W kg⁻¹, the VO₂-Mo//PANI battery can achieve an energy density (E) up to 133 W h kg⁻¹. This study provides new ideas for tailoring electrode materials with enhanced NH₄⁺-storage for AAIBs.

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1. Introduction

In the face of the growing energy crisis and climate change, there is an urgent need to reduce the burning of fossil fuels and develop sustainable energy sources. One of these is rechargeable batteries, which are considered to be a clean and sustainable solution.^1–5 The high energy density of conventional lithium-ion batteries (LIBs) has led to their widespread use but the safety issues of LIBs are also prominent, especially the thermal runaway problem triggered by their flammable and toxic organic electrolytes.^6–8 As a result, new batteries built on aqueous electrolytes have drawn a lot of interest recently because of their excellent security, affordability, and environmental friendliness in addition to their numerous application possibilities.^9–16 Aqueous batteries usually use metal ions as charge carriers, such as Li⁺,^17,18 K⁺,¹⁹ Na⁺,²⁰ Mg²⁺,²¹ and Zn²⁺.^22–24 These metal-ion carrier batteries have many advantages, such as abundant resource reserves and superior theoretical capacity. However, the excessive focus on metal ions has led to the neglect of non-metal ion carriers, which has resulted in insufficient research on novel non-metal ion water-based batteries and supercapacitors.²⁵ In fact, non-metal ions tend to have some unique advantages compared to metal ions, such as higher abundance, lower hydration radius and lighter molar mass.²⁶ Of these, NH₄⁺ has garnered significant interest because of its quick kinetics, lower molar mass (18 g mol⁻¹), and smaller hydration radius (3.31 Å).^27–30 What's more, the unique H-bond energy storage mechanism and tetrahedral geometry confer the topological electrochemical chemistry of NH₄⁺. Despite these advantages, the development of water-based NH₄⁺ rechargeable devices is still at a basic stage, and exploring suitable electrode materials is still a huge challenge.⁷

Previous studies on aqueous ammonium ion batteries (AAIBs) are mainly focused on cathode materials. Since Toshima's pioneering exploration of the feasibility of embedding NH₄⁺ in Prussian blue analogues (PBAs),³¹ a series of major breakthroughs have been made in the field of NH₄⁺ electrolyte energy storage. Cui et al. later conducted groundbreaking work on the use of PBAs for NH₄⁺ storage.³² In addition to PBAs, Dong et al. developed α-MnO₂ with a tunnel-like structure applied for the electrochemical storage of NH₄⁺, which had excellent electrochemical properties.³³ In addition to the above materials, other electrode materials, such as Berlin Green,³⁴ Ti₃C₂-Mxene,²⁸ MnO_x,¹² and vanadium-based materials,^27,35–37 can also be used for efficient NH₄⁺ storage. Among these, materials based on vanadium have garnered a lot of interest because of their distinctive shape, numerous changeable oxidation states, and vast resource reserves.^38–42 However, their poor electrical conductivity, slow ion diffusion kinetics and unavoidable dissolution in aqueous solutions are drawbacks that severely limit their NH₄⁺ storage performance.²⁶ For this reason, some effective strategies can be adopted to circumvent the above problems, and the common methods include cladding, interpolation, doping and defect engineering. For example, Zhang et al. developed polypyrrole intercalated VOH for efficient aqueous ammonium ion storage.⁴³ Chen et al. reported the use of interlayer engineering to promote aqueous ammonium ion storage by synergizing K⁺ and polyaniline (PANI) during vanadium oxide hydration.³⁰ Liu et al. explored a defect engineering strategy to promote ammonium vanadate electrode materials for high-performance aqueous ammonium ion supercapacitors.⁴¹ These modification strategies have made great contributions to aqueous ammonium ion energy storage. Notably, this is a particularly interesting approach in terms of defect engineering, which can change the nature of the electronic structure without significantly affecting the lattice, thus improving the electrochemical properties of the sample. In detail, oxygen defects can initiate an unbalanced distribution of electric charges, attracting more electric charges to be incorporated into the off-domain electron cloud, thus stimulating electric charge transfer and increasing the electrical conductivity.⁴⁴ Meanwhile, oxygen defects significantly facilitate the creation of additional active sites and regulated migration paths for carrier ions (NH₄⁺), which effectively accelerates the electrochemical reaction kinetics, reduces the polarization effect, and ultimately achieves reversible storage of ions.^39,41,45

Among the vanadium oxides, vanadium dioxide (VO₂) exhibits significant promise in a range of aqueous metal ion batteries due to its tunnel-like structure (Fig. 1b),^46,47 and the research on VO₂ for AAIBs has also received attention.^39,40,48 Gong et al. developed electrode materials with power densities up to 4540 W kg⁻¹ for aqueous ammonium ion batteries by modifying the oxygen defect content in VO₂·xH₂O nanoarrays.⁴⁰ As an anode material for AAIBs, Dong et al.³⁹ produced faulty VO₂, which showed a moderate transport channel with 200 mA h g⁻¹ (0.1 A g⁻¹). The energy density of the entire battery was 96 W h kg⁻¹. Tan et al. constructed a VO₂@C composite and it showed good NH₄⁺ storage characteristics.⁴⁸ These works demonstrate that the electrochemical properties are boosted by the composite and defects, whereas, they may also be improved by other methods, such as doping, which can tailor the electronic structure of VO₂.⁴⁶ Therefore, herein, we focus on the NH₄⁺-storage of VO₂ for AAIBs enhanced by the doping strategy. Although recent works have been reported for Mo doped into VO₂ for NH₄⁺ storage, the capacity achieved by Wang et al.⁴⁹ is not high enough and the reason why this method can boost the electrochemical properties is not clearly demonstrated.⁵⁰ Besides, the latter work focused on its electrochemical properties for supercapacitors, not for AAIBs.

Fig. 1 Characterization of VO₂-Mo: (a) sample synthesis process diagram; (b) the VO₂-Mo crystal structure from several angles; (c) XRD patterns of VO₂-Mo and VO₂; (d) the small 2θ range; and (e) EPR spectra of VO₂ and VO₂-Mo.

Based on the above statements, a simple process is developed in this study to produce appropriate oxygen defects based on a mild reaction environment by introducing an appropriate amount of Mo into VO₂ (denoted as VO₂-Mo). When Mo is doped successfully, VO₂'s lattice spacing increases, which is advantageous for NH₄⁺ storage that is effective. The guest Mo atom in the VO₂ crystal structure can boost mass transfer, accelerate ion transport, and accelerate kinetics, as demonstrated by both experimental results and DFT simulations. This results in increased NH₄⁺-storage properties of the VO₂-Mo composite. It is proved that the Mo-doped sample VO₂-Mo has superior electrochemical properties and cycling stability compared to VO₂. With an excellent specific capacity of 370 mA h g⁻¹ under 0.1 A g⁻¹, the former still maintains a high capacity retention rate of 85% even after 6000 cycles. Excellent NH₄⁺ storage performance is also demonstrated by the full-cell VO₂-Mo//PANI, which is composed of polyaniline (PANI) as the cathode material and VO₂-Mo as the anode material. This strategy may open more paths in advanced NH₄⁺ energy storage materials, driving the future boom of aqueous nonmetallic ion storage systems.

2. Results and discussion

2.1. Composition and structure of VO₂-Mo

As described in Fig. 1a, molybdenum doped vanadium dioxide (denoted as VO₂-Mo) with 2% Mo atom percentage was synthesized directly by a simple hydrothermal method using V₂O₅ and (NH₄)₂MoO₄ as the raw sources, as well as anhydrous ethanol as the reductant. H₂O₂, here, plays a vital role in converting the solid V₂O₅ to the soluble coordination compound obeying the following chemical equation:⁵¹ V₂O₅ + 4H₂O₂ → 2[VO(O₂)₂]⁻ + 3H₂O + 2H⁺. This is very important for doping because the soluble 2[VO(O₂)₂]⁻ and ammonium molybdate can form a homogeneous solution and they can coprecipitate to form VO₂-Mo under hydrothermal conditions. VO₂-Mo has a typical tunneling structure, displayed in Fig. 1b, which is good for rapid ionic migration.⁵² The tunnel-like VO₂ has been demonstrated to be a promising material for Li-ion and Zn-ion storage,⁵³ as well as ammonium-ion storage.³⁹ The performance of VO₂ is enhanced because the guest atom optimized its electronic structure⁴⁶ and this property may also be suitable for ammonium-ion storage. Fig. 1c displays the X-ray diffraction (XRD) patterns, and all peaks are able to match perfectly with the standard card of VO₂ (JCPDS, No. 31-1438). Comparing VO₂ with VO₂-Mo (Fig. 1d), the (002) crystal plane (2θ about 29°) is significantly shifted to a lower angle by the introduction of Mo. The fact that the massive Mo atom is effectively doped into the crystalline structure of VO₂ in place of the V position is demonstrated by the increased crystal plane spacing. The tuned structure of VO₂ can favor the insertion and detachment of ammonium ions like other vanadium oxides.⁵⁴ It is worth noting that the introduction of Mo modulates the structure of VO₂ under hydrothermal conditions, which may make the exposure of some of its crystal facets not exactly the same. For example, the diffraction peaks at 28.8° and 30.1° correspond to the (002) and (−401) crystal facets, respectively, where the (−401) crystal facet is relatively more exposed for VO₂-Mo. The introduction of Mo does not create any impurity, but the degree of VO₂-Mo crystallinity is decreased because the XRD peaks become weak (Fig. 1c and d), which suggest that there are more oxygen defects in VO₂-Mo. In order to verify the above conjecture, we performed electron paramagnetic resonance (EPR) measurements on VO₂-Mo and VO₂, as seen in Fig. 1e, where VO₂-Mo exhibits two substantially symmetric peaks at g = 2.0, which is attributed to anionic vacancy trapping of electrons. These prove that the introduction of Mo can induce more oxygen defects, which will improve the charge storage of VO₂.⁴⁰

Using N₂ adsorption/desorption isotherms, the porous structural information of VO₂ and VO₂-Mo is examined and is displayed in Fig. 2a. The isotherms are basically assigned to III and IV with an H3 hysteresis loop as defined by IUPAC. The specific surface area (S_BET) of VO₂-Mo, measured by the Brunauer–Emmett–Teller (BET) method, is around 43 m² g⁻¹, greater than the value of VO₂ (24 m² g⁻¹). The pore distributions of VO₂ and VO₂-Mo (Barrett–Joyner–Halenda (BJH) method) are depicted in Fig. S1 (ESI†). It can be seen that the introduction of Mo in VO₂ can induce some pores. The total pore volumes of VO₂ and VO₂-Mo are found to be 0.283 cm³ g⁻¹ and 0.394 cm³ g⁻¹, respectively. The results from the N₂ adsorption/desorption isotherms furthermore confirm that the Mo-doping can tune the structure of VO₂. Fourier transform infrared spectroscopy (FTIR) was used on the substances both before and after doping to investigate the FTIR spectra of VO₂ and VO₂-Mo, displayed in Fig. S2 (ESI†), in order to further confirm the chemical bonding composition of the products. The V=O and V–O–V bonds are responsible for the vibrational peaks at roughly 1000 cm⁻¹ and 500 cm⁻¹,⁴⁰ respectively. Compared with VO₂, the vibrational peaks in VO₂-Mo are slightly red-shifted, indicating the chemical environment change induced by Mo-doping. It is noteworthy to observe that a wavenumber of about 815 cm⁻¹ is indexed to the characteristic Mo–O vibrational peak, which strongly proves the successful doping of Mo. Fig. S3 (ESI†) depicts the Raman spectra of VO₂ and VO₂-Mo. The stretching vibration of (V₂O₂)n, the bending vibration of V=O, the bending vibration of V–O₃–V, the V₂–O stretching vibration, and the V–O stretching vibration are represented by the Raman peaks that are present in both of them at around 140 cm⁻¹, 280 cm⁻¹, 407 cm⁻¹, 695 cm⁻¹, and 992 cm⁻¹.⁵⁵ As seen in Fig. S4 (ESI†) and Fig. 2b and c, XPS spectra are used to further explore and offer information on the chemical composition and oxidation states of VO₂ and VO₂-Mo. The full XPS spectra (Fig. S4†) clearly show that the Mo 3p and Mo 3d peaks appear in VO₂-Mo, and the remaining peaks align with every one of the VO₂ peaks. The atomic percentage of Mo/V is about 1.9%, which is in line with the used raw materials. The V 2p XPS spectra of VO₂ and VO₂-Mo are compared and shown in Fig. 2b, where V 2p_1/2 and V 2p_3/2 have binding energies of about 524.2 eV and 516.8 eV for VO₂ and 523.8 eV and 516.5 eV for VO₂-Mo, respectively. What's more, the difference of V 2p_1/2 and V 2p_3/2 in these two samples is 7.3 eV. The above results prove that the V in the above samples is in the +4 oxidation state,⁴⁰ and the doping of Mo leads to a slight shift of V 2p binding energies in VO₂-Mo. The Mo 3d XPS spectra of VO₂-Mo in Fig. 2c are divided into Mo 3d_3/2 (235.8 eV) and Mo 3d_5/2 (232.6 eV). These values combined with the peak difference of 3.2 eV prove that Mo is in the +4 oxidation state in VO₂-Mo.⁵⁶ The XPS results further prove the successful doping of Mo. Scanning electron microscopy (SEM) images show that both VO₂ (Fig. S5a, ESI†) and VO₂-Mo (Fig. 2d and e) have the structures of nanobelts. The SEM elemental mapping images provide additional confirmation of the elemental composition of VO₂ and VO₂-Mo, as shown in Fig. S6 (ESI†) and Fig. 2f, respectively. They reveal the presence of elements V, O and Mo in VO₂-Mo (Fig. 2f), while only elements V and O are present in VO₂ (Fig. S6, ESI†). The elemental distributions are very homogeneous in both, which confirms the successful doping of Mo and the absence of impurities introduced during the doping process. The TEM pictures provide additional confirmation of the VO₂ and VO₂-Mo nanobelt shape (Fig. S5b, ESI;†Fig. 2g and h). A high-resolution TEM (HRTEM) picture of VO₂-Mo is displayed in Fig. 2i. It is evident that good crystallization is present, and the clear crystalline lattice spacing is around 0.357 nm, in accordance with the VO₂ (110) plane, suggestive of the successful synthesis of VO₂. In summary, the above characterization results show that the introduction of Mo atoms modulates the crystal facet index of tunneling VO₂, enlarges its spacing, and generates appropriate oxygen defects, while it also increases the specific surface area of VO₂ and enriches its microscopic pore structure, which are conducive to the enhancement of the ion transport kinetics of the carriers (NH₄⁺) during the electrochemical process and favors efficient ammonium-ion storage, and ultimately to the enhancement of the material's electrochemical performance.⁴⁰

Fig. 2 Structure and morphology of VO₂ and VO₂-Mo: (a) N₂ adsorption–desorption isotherms; (b) V 2p XPS spectra; (c) Mo 3d XPS spectra of VO₂-Mo; (d and e) SEM images of VO₂-Mo; (f) SEM elemental mapping images of VO₂-Mo; (g and h) TEM images and (i) HRTEM image of VO₂-Mo.

2.2. The VO₂-Mo electrode's NH₄⁺ storage properties and mechanism

It is widely believed that the generation of appropriate oxygen vacancies in the lattice is beneficial to increase the active sites for ion storage, accelerate ion diffusion, increase the material's electrical conductivity, and ultimately enhance the comprehensive electrochemical performances of materials for NH₄⁺ storage.^39–41 In this work, a standard three-electrode system was assembled utilizing 1 mol L⁻¹ (NH₄)₂SO₄ as the electrolyte in order to investigate the electrochemical characteristics of VO₂ and VO₂-Mo for the NH₄⁺ storage. As shown in Fig. S7 (ESI†), the three electrodes in the system were a working electrode (VO₂-Mo or VO₂), a reference electrode (Ag/AgCl), and a counter electrode (carbon rod). To investigate how the amount of Mo doping affects VO₂'s electrochemical performance, the Mo contents were adjusted to 0%, 1%, 2% and 4% (details in the ESI†). The XRD patterns in Fig. S8 (ESI†) confirm their successful preparation. The galvanostatic charge–discharge (GCD) curves of these samples at 0.5 A g⁻¹ are shown in Fig. S9 (ESI†), which proves that the most significant performance enhancement of VO₂ is achieved at a Mo content of 2%. To show the superiority of Mo doping, we focus on the comparison of VO₂ and VO₂-Mo for the NH₄⁺ storage using various tests. The cyclic voltammetry (CV) curves at various scan speeds within the potential window of −0.7–0.3 V are displayed in Fig. S10 (ESI†). The insertion/extraction of NH₄⁺, which results in vanadium's valence change (IV ↔ III), is what causes the evident redox peaks to be seen. The CV curves at 1 mV s⁻¹ are shown in Fig. 3a. Both VO₂ and VO₂-Mo have two pairs of redox peaks although a reduction peak is not obvious, which coincide with the two steps of NH₄⁺ (de)intercalation. The redox peaks of VO₂ are located at about −0.539/–0.659 V and −0.029/–0.356 V, whereas, the redox peaks of VO₂-Mo are located at about −0.570/–0.631 V and −0.172/–0.422 V. Compared with VO₂, VO₂-Mo exhibits smaller potential gaps (0.061 V and 0.250 V vs. 0.120 V and 0.327 V) and higher peak current densities. Such phenomena reveal that VO₂-Mo shows a smaller polarization because of better redox reaction kinetics and more rapid ion diffusion.⁵⁷ The redox couples of VO₂-Mo move to the center position compared to VO₂, demonstrating that Mo-doping is not only good for the redox reaction in kinetics, but also reduces the Gibbs free energy in thermodynamics, thus leading to better electrochemical performances. Fig. 3b and Fig. S11 (ESI†) show the GCD curves for VO₂ and VO₂-Mo at various current densities, respectively, which reflect their rate performances summarized in Fig. 3c. The platforms in these GCD curves correspond with the redox reactions in their CV curves. The specific discharge capacities of VO₂-Mo are around 370, 285, 225, 180, 137, and 87 mA h g⁻¹, which are higher than those of VO₂, which are about 232, 192, 143, 110, 78, and 44 mA h g⁻¹, respectively, at 0.1, 0.2, 0.5, 1, 2, and 5 A g⁻¹. After 50 cycles, the specific discharge capacities are approximately 210 mA h g⁻¹ and 310 mA h g⁻¹, respectively, when the current density reaches 0.1 A g⁻¹. The GCD curves in Fig. S11† and Fig. 3b support poor coulombic efficiency (CE), which is the cause of the decrease in particular capacities. Notably, as shown in Fig. 3d and Table S1 (ESI†), the obtained specific capacities of VO₂-Mo surpass those of many reported anode materials for ammonium ion storage, including defective VO₂ (220 mA h g⁻¹, 0.1 A g⁻¹),⁴⁰ VOPO₄·2H₂O (154.6 mA h g⁻¹, 0.1 A g⁻¹),⁵⁸ d-VO (∼200 mA h g⁻¹, 0.1 A g⁻¹),³⁹ MnAl-LDH (84 mA h g⁻¹, 0.1 A g⁻¹),⁵⁹ FVO (88.6 mA h g⁻¹, 0.1 A g⁻¹),⁶⁰ MoO₃ (115 mA h g⁻¹, 0.1 A g⁻¹),²⁹etc. In addition to the excellent specific capacity and multiple performances, the VO₂-Mo anode material also has better cycling stability compared with VO₂. As demonstrated in Fig. 3e, after approximately 6000 consecutive charging and discharging cycles at a current density of 2 A g⁻¹, the specific capacity of VO₂-Mo essentially stays constant, reaching approximately 140 mA h g⁻¹ (98% of the initial value). In contrast, the specific capacity of VO₂ stays at 37 mA h g⁻¹ (51% of the initial value). Vanadium oxide materials’ inherent features, such as their slow breakdown and structural deterioration, are what lead to the decrease in the specific capacity of VO₂.^3,8,61 Such a result illustrates that the introduced Mo can stabilize the structure of VO₂, thus leading to enhanced cycling durability. Compared with regular CE of VO₂-Mo, the coulombic efficiency (CE) of VO₂ fluctuates acutely (Fig. 3e), which also supports the conclusion of the stable structure of VO₂-Mo for highly reversible NH₄⁺ (de)intercalation. All results affirm that the overall electrochemical performances of VO₂ can be well enhanced by the introduced Mo to produce appropriate oxygen defects.

Fig. 3 Electrochemical properties of VO₂ and VO₂-Mo: (a) comparison of CV curves at 1 mV s⁻¹; (b) GCD curves of VO₂-Mo at different current densities; (c) rate performances; (d) comparison of the capacity with literature studies; and (e) cycling performances.

Fig. 4a displays the CV curves at five distinct scanning rates of 0.2–1 mV s⁻¹ in order to better evaluate the kinetics of VO₂-Mo for NH₄⁺ intercalation/deintercalation. The oxidation and reduction peaks invariably change slightly at different scanning speeds because of the polarization effect, but the shapes of the CV curves maintain a high degree of similarity, indicating that the material has good stability and remarkable electrochemical kinetics.⁶² The previous studies proved that these CV curves’ peak current (i) and scanning rate (v) should match with the following two equations:⁶³

i = av^b (1)

Fig. 4 Electrochemical kinetic studies of VO₂-Mo: (a) CV curves obtained at various scan rates; (b) the fitting log(i) vs. log(v) curves; (c) the capacitive contribution graph; (d) contribution ratios; (e) Nyquist charts and (f) the relationship of Z‘ − ω^−1/2.

When a and b are both constants, then eqn (1) can be changed to:

log i = log a + b log(v) (2)

The relationship between i and v is evident in eqn (2), where b denotes the slope relationship between log(i) and log(v). In particular, it is demonstrated that surface capacitance predominates when b is close to 1, and that ion diffusion's contribution to capacity prevails when b is close to 0.5. As shown in Fig. 4b, Peak 1 and Peak 2's computed b-values of 0.98 and 0.94, respectively, demonstrate that the surface capacitance contribution and a small amount of diffusion characteristics are the primary drivers of the ammonium ion storage process in VO₂-Mo. The diffusion and capacitance contribution ratios of VO₂-Mo in ammonium ion storage at different scanning rates can be calculated using eqn (3):³⁰

i = k₁v + k₂v^1/2 (3)

where k₁v represents capacitive behavior, k₂v^1/2 represents ion diffusion control, k₁ represents percentage contribution in response to capacitive behavior, and k₂ represents diffusion control percentage. The contribution ratios of diffusion and capacitance are shown in Fig. 4d, while Fig. 4c displays the capacitive contribution at 1 mV s⁻¹. When the CV scanning rate increases from 0.2 mV s⁻¹ to 1 mV s⁻¹, the percentage of capacitive contribution progressively increases from 72% to 93%. This is also similar to the trend when the b value is close to 1. The result proves that VO₂-Mo has excellent charge reaction kinetics, which is an important reason for its excellent rate performance and highly reversible capacity. In order to reveal the intrinsic reason for this phenomenon, the conductivity of the VO₂-Mo and VO₂ electrode materials was examined using electrochemical impedance spectroscopy (EIS), as shown in Fig. 4e. The electrode's charge transfer resistance (R_ct) is represented by the circular section of the Nyquist curve at high frequency, and the Warburg diffusion impedance associated with NH₄⁺ diffusion is represented by the linear portion at low frequencies.⁶⁴ It can be seen that the radius of the VO₂-Mo electrode is much smaller and its slope is steeper than that of VO₂, proving the low intrinsic resistance, VO₂-Mo's higher mass-transfer resistance and quick ion migration for NH₄⁺ storage.⁶⁵ The ion diffusion rate is indicated by the slope b generated by the diffusion of ions in the low-frequency domain, as demonstrated by Z′ − ω^−1/2 fitting curves (Fig. 4f) with b values of 12.7 and 9.3. The aforementioned findings clearly show that in the VO₂-Mo material, ammonium ions migrate more quickly and the rate of redox reaction is likewise faster. This indicates that the oxygen vacancies generated in VO₂ by employing Mo are favorable for electron transfer and accelerate ion diffusion, which explains well why the integrated electrochemical performances of VO₂-Mo are much better than that of VO₂.

Several ex situ characterization techniques, including XRD, FTIR, and XPS tests, were used to reveal the mechanism of NH₄⁺-storage behavior in VO₂-Mo electrodes during charging and discharging. Fig. 5f shows the five chosen states of A/D and pristine labeling on the GCD curves of the VO₂-Mo electrode. The ex situ XRD patterns in different states (Fig. 5a) reveal that the VO₂-Mo electrode shows the characteristic diffraction peaks of VO₂ (JCPDS, No. 31-1438) before and after applying the potential except some peaks from the substrate Ti foil. This proves the highly stable structure of VO₂-Mo for the ammoniation/deammoniation processes. The enlarged XRD patterns show the shift of the (110) peak during charging and discharging (Fig. 5b). When fully discharged to −0.7 V (from pristine to A), the (110) peak of the electrode material shifts to a higher angle, indicating that the ammonium ions are clearly inserted into the lattice structure of the material, and conversely, at a full charge state of 0.3 V (A to B), the (001) peak returns to a lower angle, indicating that the ammonium ions are deintercalated from the lattice of VO₂-Mo. More impressively, when the applied potential is again reduced to −0.7 V (B to C) and improved to 0.3 V (C to D), the corresponding peak also returns to a higher (lower) degree, which corresponds to the state of A (B), strongly demonstrating that the NH₄⁺ insertion/deinsertion reaction is highly reversible. It is also noteworthy that no impurities are detected in all XRD patterns, which implies that the whole charging/discharging process is very efficient and does not produce any by-products. Fig. 5e shows the ex situ FTIR spectra in different selected states, and the peaks associated with the V–O stretching vibrations appear in the low-wave number region (below 1100 cm⁻¹). When NH₄⁺ enters the tunneling structure of VO₂-Mo during the discharge process, H atoms combine with the O atoms to form hydrogen bonds. The nitrogen–hydrogen bond's bending vibration is attributed to the peak at 1394 cm⁻¹, and its stretching vibration is attributed to the peak at 3195 cm⁻¹. The peaks at 2800 and 1078 cm⁻¹ represent the change of the crystal structure during charging and discharging and correspond to hydrogen atoms in NH₄⁺ bound to oxygen atoms in the VO₂-Mo main skeleton.⁶⁶ When ammonium ions deintercalate (charge) or intercalate (discharge), their vibrational peaks’ intensity correspondingly reduces or increases. To further investigate the bonding mechanism of the VO₂-Mo electrode for storing NH₄⁺, we performed ex situ XPS tests for both full charge and full discharge processes. The ex situ entire XPS spectra of full charging and full discharging procedures are displayed in Fig. S12a (ESI†) and the results show that there are slight differences in the strengths and positions of the peaks corresponding to N 1s, V 2p, and Mo 3d under the full charging and full discharging conditions, which suggests that they may be related to the mechanism of ammonium ion storage. The ex situ XPS spectra of N 1s, V 2p, and Mo 3d during full charging and full discharging are shown in Fig. 5c, d and Fig. S12 (ESI†), respectively. The intensity of N 1s is obviously much weaker in the fully charged state than that in the fully discharged state, and the spectra of N 1s are mainly attributed to the protonated state (–NH⁺–, 400.4 eV) and the –NH– segment (402.0 eV) (Fig. 5c).⁴¹ The strong evidence that the charging and discharging processes are accompanied by the embedding (discharge) and detachment (charge) of NH₄⁺ by hydrogen bond is provided by the intensity variation of NH⁺–, which is consistent with the results reported by the previous report.⁶² In addition, the oxidation state changes reflected by the ex situ XPS spectra of V 2p (Fig. 5d) and Mo 3d (Fig. S12†) can again confirm the above mechanism: accompanied by NH₄⁺ embedding in the discharge process, the contents of +3 (V/Mo) increase. Nevertheless, the contents of +4 (V/Mo) decrease accompanied by NH₄⁺ detachment in the charging process.¹² Based on the above characterization, it can be fully demonstrated that the energy storage mechanism of the VO₂-Mo electrode is as follows: during charging and discharging, NH₄⁺ is delocalized/embedded from the tunneling structure and accompanied by the breaking and generation of hydrogen bonds, during which the metal V and Mo also produce corresponding valence changes to balance the charge. The VO₂-Mo electrode's NH₄⁺ storage mechanism can be visually represented as illustrated in Fig. 5f, which is based on a thorough conclusion presented above. The reaction during charging and discharging is expressed by the following equation: VO₂-Mo + xNH₄⁺ + xe⁻ ↔ Mo-(NH₄)_xVO₂-Mo. The NH₄⁺ storage mechanism of the VO₂-Mo electrode is demonstrated by the creation and disruption of hydrogen bonds, which is in keeping with the previously reported V-based materials for ammonium ion storage.^{27,38,41,54,67}

Fig. 5 Electrochemical mechanism of VO₂-Mo for ammonium ion storage: (a and b) the five points’ equivalent ex situ XRD patterns; (c and d) the ex situ XPS spectra; (e) the ex situ FTIR spectra corresponding to the five points; and (f) mechanism diagram of ammonium ion storage during charging/discharging processes, inserting the GCD curves.

2.3. VO₂-Mo//PANI aqueous NH₄⁺ batteries

To further evaluate the practicality of VO₂-Mo electrodes, as shown in Fig. 6a, we attempted to fabricate an aqueous NH₄⁺ cell. The structural characterization and electrochemical characterization of PANI are described in Fig. S13 (ESI†). PANI's CV curves display a few redox reactions, which indicates that it can store NH₄⁺. Its operating potential range is −0.3–0.7 V and PANI's specific discharge capacities at current densities of 0.2, 0.5, 1, 2, and 5 A g⁻¹ are approximately 110, 98, 89, 80, and 67 mA h g⁻¹, respectively. The above results predict that PANI as a cathode can be better paired with an anode composed of VO₂-Mo.⁶⁸Fig. 6b and Fig. S14a (ESI†) show the CV curves of VO₂-Mo//PANI and VO₂//PANI aqueous NH₄⁺ batteries at different scanning speeds, respectively, and both of them can operate well in the selected voltage window up to 1.4 V. The redox peaks correspond with the (de)intercalation of NH₄⁺ discussed above. The integrated area of the VO₂-Mo//PANI battery is significantly bigger than that of the VO₂//PANI battery, as shown by the comparison of the CV curves of the two batteries at 1 mV s⁻¹ (Fig. S14b†), proving the superior performances of VO₂-Mo//PANI battery. In order to investigate the intrinsic reasons for the advanced VO₂-Mo//PANI battery, the conductivities of VO₂-Mo//PANI and VO₂//PANI are respectively investigated by EIS (Fig. S15, ESI†). As expected, the VO₂-Mo//PANI battery has lower electrical conductivity and better charge/ion transfer behavior, which can achieve efficient and rapid ammonium ion storage. The rate performances and corresponding GCD curves of the VO₂-Mo/PANI battery at various current densities are shown in Fig. 6c and d. The redox peaks in CV curves match with the platforms in these GCD curves (Fig. 6b). The complete battery shows reversible specific discharge capacities of approximately 232, 171, 135, 102, and 62 mA h g⁻¹ for current densities of 0.2, 0.5, 1, 2, and 5 A g⁻¹, respectively, which is encouraging. The specific discharge capacities of the full battery rapidly recover to about 220 mA h g⁻¹ and 235 mA h g⁻¹ when the current densities are restored to 0.2 and 0.1 A g⁻¹. The VO₂-Mo//PANI battery achieves the specific discharge capacity as high as ca. 232 mA h g⁻¹ at 0.2 A g⁻¹, which surpasses that of the state-of-the-art AAIBs, such as, MnAlLDH//PTCDI (60 mA h g⁻¹, 0.1 A g⁻¹),⁶⁹ MnO₂//d-VO (181 mA h g⁻¹, 0.1 A g⁻¹),³⁹ γ-MnO₂//WO₃ (106 mA h g⁻¹, 0.5 A g⁻¹),⁷⁰ Mn₃(PO₄)₂·7H₂O//MoO_x (98 mA h g⁻¹, 0.1 A g⁻¹),⁷¹ NiHCF//1,5-NAPD (143 mA h g⁻¹, 0.1 A g⁻¹)⁷² and so on. The cycling stability with CE at 5 A g⁻¹ is shown in Fig. 6g. After 6000 cycles, the complete battery reaches a specific discharge capacity of roughly 55 mA h·g⁻¹, with an 85% cycle retention, as well as the CE around 100% during the repeated charge/discharge processes. Such results prove the superior cycling stability of the VO₂-Mo//PANI battery, in line with the results of the three-electrode system (Fig. 3e). To further illustrate the merits of the VO₂-Mo//PANI battery, Ragone plots compared with the previously reported batteries for ammonium ion storage are shown in Fig. 6e. Based on the weight of VO₂-Mo, the discharge energy density (E) of VO₂-Mo/PANI battery can reach up to 133 W h kg⁻¹ at a power density (P) of 89 W kg⁻¹. As shown in Table S2 (ESI†), the achieved E and P values are superior to those of previous NH₄⁺ storage devices, including ACC@VPP//PTCDI,³⁵ MoO₃@C//MoO₃@C,⁷³ NVO//AC,⁵⁴ MnAlLDH//PTCDI,⁶⁹ KVO/PANI//AC,³⁰ MoO₃//CuFe PBA,²⁹ PVO//AC,⁷⁴ MnO₂//d-VO,³⁹etc. The above characteristics prove the good energy output of the designed VO₂-Mo//PANI battery. Furthermore, a LED is powered by a series battery system to demonstrate the viability of this ammonium ion battery's practical application. As shown in Fig. 6f, the red LED has been successfully lit continuously for more than 30 minutes, showing the stable operating characteristics of the designed VO₂-Mo//PANI battery. In summary, the entire battery exhibits strong electrochemical properties and stable cycling, making it suitable for large-scale energy storage system production in the future.

Fig. 6 The complete cell performance of VO₂-Mo/PANI: (a) schematic diagram of a VO₂-Mo/PANI battery; (b) CV curves of a VO₂-Mo//PANI battery at different scan rates; (c) rate performances of a VO₂-Mo//PANI battery; (d) GCD curves of a VO₂-Mo//PANI battery at different voltage windows; (e) Ragone plots; (f) demonstration of the VO₂-Mo//PANI battery to power an LED light; and (g) cycling stability and CE of a VO₂-Mo//PANI battery.

2.4. DFT calculations

Both the three- and two-electrode systems attest that the Mo atom in VO₂ can boost the NH₄⁺-storage properties of VO₂. We conducted the density functional theory (DFT) calculations to uncover the underlying reason, whose information is shown in the ESI† in detail. The structures of VO₂ and VO₂-Mo models at different orientated views are depicted in Fig. 7a–d, which clearly represent the tunnel-like structure of VO₂. They are built, optimized and used to calcaulate the band structures, density of state (DOS), partial DOS (PDOS) and electron density difference during the DFT calculations. The band structures of VO₂ and VO₂-Mo are shown in Fig. 7e and f. The bandgap of VO₂-Mo (Fig. 7f) decreases to 1.81 eV compared with VO₂ (1.99 eV, Fig. 7e), which demostrates that the introduction of Mo increases the electronic conductivity and lowers the electron transport barrier, in line with the EIS spectra disucssed in Fig. 4e, f and Fig. S15.† Besides, in the band structures of VO₂-Mo, the energy bands become flat and the bandwidth decreases, which imply the current carrier concentration increases. It is good for the flow of electrons, enhancing the electrochemcial properteis of VO₂-Mo.⁷⁵ Fig. S16a (ESI†) shows the DOS curves of VO₂ and VO₂-Mo and Fig. S16b and c (ESI†) depicts their corresponding PDOS curves of O 2p and V 3d. After doping with the element Mo, we can find that both vanadium 3d orbitals and oxygen 2p orbitals move closer to the Fermi energy level, and the increase in the energy level of the electrons in the two orbitals means that their increased activity is one of the causative factors for the instability of the binding of positively charged protons with the ammonium ion. This feature can decrease the strength of the H-bond and the weak H-bond can greatly boost the NH₄⁺-storage.⁷⁶ The electron density differences of VO₂ and VO₂-Mo are compared in Fig. 7g and h, which reflect the change of electron cloud. After the introduction of Mo, both positive and negative charges are increased, which implies the electrons’ localization of VO₂-Mo. This feature can modulate the strength of the H-bond. Therefore, the electronic conductivity is enhanced and the electron transport barrier is lowered by Mo atoms in VO₂, resulting in the boosted NH₄⁺ kinetics during the charge and discharge processes. The structural models of VO₂ and VO₂-Mo for calculating the adsorption energy with ammonium ion are shown in Fig. S17 (ESI†). VO₂-Mo exhibits a higher adsorption energy than that of VO₂, which supports the point that the H-bond becomes weak after the introduction of Mo atoms, thus enhancing the NH₄⁺ kinetics. All the experimental and DFT results unequivocally prove that the Mo-doping strategy can serve as a useful tool to boost ammonium-ion storage properties of materials for AAIBs.

Fig. 7 The results of DFT calculations: (a and b) the structures of VO₂; (c and d) the structures of VO₂-Mo; (e) the band structure of VO₂; (f) the band structure of VO₂-Mo; (g) the electron density difference of VO₂; and (h) the electron density difference of VO₂-Mo.

3. Conclusion

In conclusion, we demonstrate that VO₂'s electronic structure is tailored by Mo doping (VO₂-Mo) to enhance ammonium-ion storage properties toward highly efficient AAIBs. VO₂-Mo is successfully prepared by adjusting the content of Mo via a facile hydrothermal method. The introduction of Mo can increase the oxygen defects in VO₂-Mo, which improves the electronic structure properties and provides more active sites and smooth migration paths for ammonium ions. Both experimental results and DFT calculations reveal that the guest Mo atom in the VO₂ crystal structure can speed up ion transport, increase mass transfer, and improve kinetics, leading to improved NH₄⁺-storage properties. VO₂-Mo (2% Mo) exhibits a specific discharge capacity of about 370 mA h g⁻¹ at 0.1 A g⁻¹, higher than that of VO₂ (232 mA h g⁻¹) and the reported vanadium oxides for NH₄⁺-storage. The capacity of 140 mA h g⁻¹ basically remains unchanged after about 6000 consecutive charging and discharging processes at 2 A g⁻¹. The creation of a hydrogen bond between NH₄⁺ and VO₂-Mo, which is corroborated by ex situ XRD, FTIR, and XPS, is the mechanism for storing NH₄⁺. The VO₂-Mo//PANI complete cell outperforms the most advanced electrode materials for AAIBs, achieving a specific discharge capacity of up to 232 mA h g⁻¹ at 0.2 A g⁻¹. At 89 W kg⁻¹, the E value of the VO₂-Mo//PANI battery can reach up to 133 W h kg⁻¹. In addition to meeting the anticipated future need for the large-scale production of aqueous NH₄⁺ energy storage systems, our work offers fresh insights into the creation of tailored electrode materials with increased NH₄⁺-storage for AAIBs.

Data availability

The data supporting this article have been included as part of the ESI.†

Conflicts of interest

There are no conflicts to declare.

Acknowledgements

This work is supported by the Doctoral Research Start-up Fund of Hubei University of Science and Technology (BK202504) and the Natural Science Foundation of Liaoning Province (2023-MS-115).

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Footnote

† Electronic supplementary information (ESI) available. See DOI: https://doi.org/10.1039/d4qi01910e

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