1 Introduction

Toxic and flammable organic electrolytes, fire or explosion safety hazards, and the scarcity and high price of lithium and cobalt make lithium-ion batteries uncompetitive in sustainable and scalable energy devices, encouraging the development of aqueous multivalent metal ion (Mg2+, Zn2+, Al3+, Ca2+, etc.) batteries with the characteristics of non-toxicity, low price, high theoretical capacity, and environmental friendliness [1,2,3,4,5,6,7]. Aqueous rechargeable zinc-ion batteries (ARZIBs) have been considered one of the most promising and competitive battery candidates because of the low redox potential (Zn2+/Zn ≈ − 0.76 V vs. standard hydrogen electrode (SHE)), high specific capacity (820 mAh·g−1, 5851 mAh·cm−3), non-toxicity, considerable chemical stability, and the natural abundance of Zn metal [3, 8,9,10,11,12]. However, the development of ARZIBs is plagued by unfavorable specific capacity, poor cyclic stability, and sluggish electrochemical kinetics of the cathode hosts because of the high polarization properties of Zn2+ and strong electrostatic interaction between the cathode and Zn2+, which induces undesirable cathode dissolution into the electrolyte and insufficient insertion/extraction/diffusion of Zn2+ in cathode hosts [9, 13,14,15,16]. Therefore, exploiting suitable high-performance cathode materials is imperative for the commercial application of ARZIBs.

Two types of cathodes have been commonly reported for ARZIBs: intercalation-type (Prussian blue analogs, manganese-based oxides, vanadium-based oxides/sulfides, polyanionic compounds and MoS2, etc.) and conversion-type (Co3O4, NiO, organic compounds, etc.) [1, 2, 9, 16]. Among these, vanadium oxides with open frameworks and multiple oxidation states have become the most competitive cathode candidates owing to their high theoretical specific capacity [17,18,19]. However, narrow interlayer spacing, low conductivity, and the strong electrostatic interaction between the V–O skeleton and Zn2+ restrict the diffusion of Zn2+ and induce vanadium dissolution into the electrolyte, resulting in sluggish kinetics and severe capacity decay of vanadium oxides [16, 17, 20]. The pre-intercalation of cations and the introduction of structural water could effectively expand the interlayer spacing of vanadium oxides, and structural water could also play an electrostatic shielding role for Zn2+ and weaken the electrostatic interactions between the V–O framework and Zn2+ [10, 19, 21,22,23,24,25,26]. Nevertheless, the high atomic weight and inert electrochemical activity of the pre-intercalated metal cations hinder the advantages of the high theoretical capacity of vanadium oxides [17, 18, 27,28,29].

Recently, hydrated ammonium vanadate has attracted considerable research interest owing to the low atomic weight and large size of the NH4+ ion, and the cohesion of the layered structure can also be efficiently enhanced by hydrogen bonding between NH4+ and the V–O framework [12, 17, 20, 30,31,32]. From these perspectives, (NH4)2V10O25·8H2O (NVOH) with a 1 nm interlayer distance and NH4+/H2O pillars between the VO layers has obvious structural advantages for reversible Zn2+ (de)intercalation without volume expansion [17, 31, 33]. Furthermore, defect engineering is considered to be an effective strategy to regulate the electronic structure of electrode materials, and could be combined with other strategies to achieve multiple structural optimization, thereby improving the electrical conductivity and facilitating the ion diffusion dynamics [3, 17, 34, 35]. Therefore, exploring (NH4)2V10O25·8H2O with extended layer spacing, plenty of oxygen defects, and high electrical conductivity should be an effective strategy to improve its Zn2+ storage performance [3, 31].

Herein, oxygen-deficient (NH4)2V10O25·xH2O/GO (NVOH@GO) composites with expanded layer spacing were fabricated via a facile solution synthesis strategy. The encapsulation of (NH4)2V10O25·xH2O into graphene oxide (GO) increased the layer spacing and improved the electrical conductivity of NVOH, accelerating ion diffusion. The electron transfer and strong interaction between NVOH and GO induced more oxygen vacancies in the NVOH@GO composites, offering additional active sites for ion storage and providing additional electron transfer paths. Furthermore, the as-prepared NVOH@GO suppressed the dissolution of vanadium and effectively reduced self-discharging, indicating stable structure of the electrode and favoring superior cyclic stability. Consequently, the optimized NVOH@GO delivered a high specific capacity of 418 mAh·g−1 at 0.5 A·g−1 and a stable capacity of 238 mAh·g−1 after 10,000 cycles at 20 A·g−1. Finally, the electrochemical mechanism, including a phase transition reaction and subsequent Zn2+/H2O co-(de)intercalation processes, was elaborated using ex-situ technologies.

2 Experimental

2.1 Preparation of NVOH and NVOH@GO-x

Preparation of NVOH@GO-x: 0.4 g VCl3 and 0.4 g NH4HCO3 was dissolved in 10 and 40 mL of deionized water, respectively. After adding the VCl3 solution dropwise into the NH4HCO3 solution while stirring, the mixture was stirred at 50 °C for 4 h. A graphene oxide solution (0.0059 g·mL−1) was added to the above mixture under stirring. After ultrasonic treatment for 30 min, the uniform dispersion was heated at 80 °C for 2 h. Subsequently, 6 mL of H2O2 (30 wt%) was added dropwise to the above uniform dispersion at 80 °C, and the resulting dispersion was stirred at 80 °C for another 2 h. Finally, NVOH@GO-x was obtained by centrifugation, washed with deionized water, and dried at 100 °C for 12 h in a vacuum oven. The samples with the addition of 0.5, 1.0, and 1.5 g graphene oxide solutions were labeled as NVOH@GO-1, NVOH@GO-2, and NVOH@GO-3, respectively.

Preparation of NVOH: The NVOH preparation process was similar to that of NVOH@GO-x, except that graphene oxide was not added.

2.2 Material characterization

Scanning electron microscopy (SEM) and transmission electron microscopy (TEM) images were captured using a Zeiss Sigma 500 field-emission scanning electron microscope and JEOL JEM-2100F field-emission transmission electron microscope, respectively. X-ray diffraction (XRD) patterns were obtained on a Bruker D8-Focus X-ray diffractometer using Cu-Kα radiation (λ = 0.15406 nm). Fourier transform infrared spectroscopy (FTIR), Raman spectroscopy, and X-ray photoelectron spectroscopy (XPS) were conducted using a Bruker Smart Lab infrared spectrometer, HR-800 high-resolution confocal micro-Raman system, and PHI-1600 photoelectron spectrometer, respectively. Thermogravimetric analysis (TGA) was performed using a TG/DTA 6300 thermogravimetric analyzer in air. Inductively coupled plasma (ICP) analysis was performed on a Thermo Scientific iCAP RQ mass spectrometer. Electron paramagnetic resonance (EPR) spectra were acquired using a Bruker EMX Micro spectrometer.

2.3 Electrochemical characterization

CR2032 coin cells with metallic Zn as the counter electrode were assembled in air to investigate the electrochemical performance of the fabricated materials. To prepare the working electrode, a slurry of 70 wt% active materials, 20 wt% acetylene black and 10 wt% polyvinylidene fluoride (PVDF) with N-methyl-pyrrolidone as a solvent was coated on carbon paper and dried in a vacuum oven at 80 °C for 12 h; and the areal mass loading of the active materials was ~ 1 mg·cm−2. The electrolyte used was 2 mol·L−1 Zn(CF3SO3)2, and glass fiber was used as the separator. Cyclic voltammetry (CV) at various scan rates and electrochemical impedance spectroscopy (EIS) at frequencies of 10 mHz–100 kHz were measured on a Parstat 4000 + electrochemical workstation. Galvanostatic intermittent titration technique (GITT) and galvanostatic charge/discharge (GCD) curves were obtained using a Neware (CT-4008 T) instrument, and the GITT tests were performed using a pulse of 20 min at 50 m A·g−1 followed by a 120-min interruption.

3 Results and discussion

3.1 Structural and morphological characterization

NVOH@GO-x composites were fabricated via a facile solution synthesis method using VCl3 as the vanadium source and H2O2 as the oxidant (Fig. 1a). First, VCl3 dissolved in deionized water was decomposed into VCl2, HCl, and H2V4O9 (Reaction (1)), which gradually reacted with NH4HCO3 to form ammonium vanadate. Then, GO was anchored on the ammonium vanadate, and low-valence vanadium was gradually oxidized to high-valence vanadium by H2O2, which was confirmed by a change in the color of the solution after adding H2O2 (Fig. 1b). Finally, ammonium vanadate with high-valence vanadium covered with GO was fabricated.

$${\text{8VCl}}_{{3}} + {\text{9H}}_{{2}} {\text{O}} = {\text{4VCl}}_{{2}} + {\text{H}}_{{2}} {\text{V}}_{{4}} {\text{O}}_{{9}} + {\text{16HCl}}$$
(1)
Fig. 1
figure 1

a Schematic illustration of fabrication process and b color change during H2O2 oxidation of NVOH@GO-2; structural characterization of NVOH and NVOH@GO-x: c XRD patterns; d FTIR spectra; e Raman spectra; f XPS survey spectra; high-resolution XPS spectra of g C 1s, h N 1s, i V 2p and j O 1s; k EPR spectra

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XRD patterns of NVOH and NVOH@GO-x in Fig. 1c exhibit similar diffraction peaks around 8.2°, 25.6°, 34.8°, 47.3°, 50.5°, and 60.7°, which match well with the diffraction peaks of (NH4)2V10O25·8H2O (JCPDS No. 26-0097). The main diffraction peak at 8.2° corresponding to the (001) plane of NVOH shifts to 7.9° progressively with the addition of GO, indicating enlarged layer spacing in NVOH@GO-x. No additional peaks are detected for NVOH@GO-x, indicating that the addition of GO does not affect the phase structure of NVOH. The structures of NVOH and NVOH@GO were further characterized using FTIR spectroscopy. As shown in Fig. 1d, the two bands at 523.8 and 767.9 cm−1 are assigned to the symmetric and asymmetric stretching vibration of V–O–V bonds, respectively [20, 21, 36]. The absorption peak located at approximately 1004.6 cm−1 is attributed to the V=O stretching vibration of V5+ [20, 37, 38]. The peaks at approximately 1402 and 3139.8 cm−1 are the symmetric bending and asymmetric stretching vibration of the N–H bond of NH4+, confirming the presence of NH4+ in NVOH and NVOH@GO-x [20, 36, 38]. The other two bands at 3597.8 and 1651.7 cm−1 correspond to H–O–H stretching vibration and O–H bending vibration, implying the presence of coordinated water in NVOH and NVOH@GO-x [38, 39]. The typical Raman signals at 162, 273, and 409 cm−1 are attributed to the (V2O2)n, O3–V=O, and V–O3–V bending vibrations, respectively, while the Raman peaks at 521, 711, and 1020 cm−1 are assigned to the stretching vibrations of V3–O, V2–O, and V=O, respectively (Fig. 1e) [17, 19, 20, 40, 41]. The Raman peaks at 1341.5 and 1608.3 cm−1 in NVOH@GO-x can be indexed to the D and G bands of graphitized carbon, indicating the integration of GO into NVOH [35, 42,43,44,45]. The red shifting of the O3–V=O bending vibrations suggests the elongation of V=O bonds along the c-direction, resulting in an enlarged interlayer distance in NVOH@GO-x, whereas the blue shifting of other V–O–V bonds originates from the strong interaction between V and O in NVOH@GO-x [20, 41]. XRD, FTIR, and Raman results confirm the successful fabrication of NVOH and NVOH@GO composites, and the addition of GO expands the interlayer spacing of NVOH. The enlarged layer spacing provides capacious channels for the insertion and extraction of Zn2+ from NVOH and weakens the interaction between the adjacent V2O5 layers, which favors the fast electrochemical kinetics of NVOH@GO-x [20, 27, 46].

TGA curves for NVOH in Fig. S1 show three weight loss steps: below 150 °C, 150–250 °C and 250–350 °C, which can be ascribed to the release of physically absorbed water, structural water, and NH3, respectively [17, 20, 47]. The weight loss above 350 °C in NVOH@GO-x is due to the decomposition of GO [48, 49]. The GO content in NVOH@GO-1, NVOH@GO-2, and NVOH@GO-3 is calculated to be 0.6 wt%, 1.7 wt%, and 3.5 wt%, respectively. For NVOH@GO-x, the non-coincident TGA curves for NH3 release show that a strong interaction between GO and NVOH exists in NVOH@GO-x, which provides additional electron transfer paths and improves electronic conductivity [50, 51].

XPS spectra in Fig. 1f reveal the co-existence of V, O, C, and N in NVOH@GO-x. The high-resolution C 1s peaks at 284.6 and 285.0 eV in Fig. 1g can be attributed to C–C/C=C and C=O bonds, respectively [42, 52]. For N 1s (Fig. 1h), the two peaks at 401.9 and 399.9 eV correspond to protonated –NH4+– and –NH– segments, respectively [19, 37, 53]. The characteristic peaks of V 2p at 517.7/525.4 eV and 516.5/524.2 eV can be assigned to V5+ and V4+, respectively (Fig. 1i) [17, 20, 47, 53,54,55]. The V4+/V5+ ratio increases from 0.14 for NVOH to 0.26 for NVOH@GO-3, implying that more V4+ and oxygen defects are present in NVOH@GO composites. The N 1s and V 2p spectra of NVOH@GO-x shift to lower binding energies in contrast to those of NVOH, which can be ascribed to electron transfer and strong interaction between NVOH and GO [51]. The O 1s spectra can be deconvoluted into three peaks at 530.3, 530.8, and 532.2 eV (Fig. 1j), which can be ascribed to the V–O bond, O–H bond, and H2O molecules, respectively [17, 56]. The obvious XPS signals for H2O molecules indicate that more adsorbed H2O is present in the NVOH@GO composites, which is conducive to the infiltration of the aqueous electrolyte. V4+ has one electron in its 3d orbital, and the EPR intensity represents the concentration of V4+ and oxygen defects [17, 18]. The stronger EPR signals with a spectral splitting factor (g) value of 1.969 in NVOH@GO-x reveal that GO addition induces more oxygen vacancies in NVOH (Fig. 1k) [18, 57, 58]. The XPS and EPR results confirm that the electron transfer and strong interaction between NVOH and GO induce more oxygen vacancies in the NVOH@GO composites, which offers additional active sites for ion storage and jump sites for charge transfer [17, 18].

The morphologies of NVOH and NVOH@GO-x are shown in Figs. 2, S2. NVOH exhibits a honeycomb structure with wrinkled nanosheets (Fig. 2a, b). In NVOH@GO-x (Figs. 2c, d, S2), the honeycomb structure can hardly be distinguished because of coverage by GO on the NVOH, and wrinkled sheets are observed in the enlarged images, indicating the GO coating on the NVOH. The TEM images in Fig. 2e, g confirm the sheet morphology of NVOH and NVOH@GO-2, respectively. High-resolution TEM (HRTEM) images exhibit the clear lattice dislocations (marked by yellow ovals) with lattice spacings of 0.99 and 1.10 nm in NVOH and NVOH@GO-2 (Fig. 2f, h), respectively, matching the (001) crystal lattice plane of (NH4)2V10O25·8H2O, which is consistent with the XRD and XPS results. Figure 2i, j displays the uniform distribution of all elements in NVOH and NVOH@GO-2, proving the successful intercalation of NH4+ in NVOH and the encapsulation of NVOH into GO. Morphological characterization verifies the formation of NVOH and NVOH@GO composites, and the introduction of GO expands the layer spacing.

Fig. 2
figure 2

Morphological characterization of NVOH and NVOH@GO-2: SEM images of a, b NVOH and c, d NVOH@GO-2; TEM images of e, f NVOH and g, h NVOH@GO-2; TEM images and corresponding energy-dispersive spectroscopy elemental mapping for i NVOH and j NVOH@GO-2

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3.2 Electrochemical performance and reaction kinetic

CV curves at 0.1 mV·s−1 in Figs. 3a, S3a–c both display two pairs of broad peaks at 0.97/0.95 V and 0.67/0.62 V, which are related to the redox couples of V5+/V4+ and V4+/V3+ during the Zn2+ insertion/extraction process, respectively [18, 20, 27]. After the first activation cycle, the NVOH, NVOH@GO-1, NVOH@GO-2, and NVOH@GO-3 electrodes deliver average capacities of 303, 332, 349, and 342 mAh·g−1 at 0.5 A·g−1 respectively; and the GCD curves exhibit similar voltage plateaus to the CV results (Figs. 3b, S3d–f). The overlapping CV and GCD curves suggest high reversibility of the fabricated electrodes [18, 38]. Impressively, all NVOH@GO-x electrodes deliver better rate capability than NVOH through three consecutive loops at current densities of 0.5 to 20.0 A·g–1 (Figs. 3c, d, S3g–i). In particular, at 0.5, 1.0, 2.0, 5.0, 10.0, and 20.0 A·g–1, the average discharge capacities of the optimized NVOH@GO-2 electrode are 418, 396, 373, 339, 303, and 260 mAh·g–1, respectively; and the capacities at 1.0, 2.0, 5.0, 10.0, and 20.0 A·g−1 are equivalent to 95%, 89%, 81%, 72%, and 62% of that at 0.5 A·g−1, while the capacities of NVOH at the above current densities are only 336, 318, 292, 246, 205, and 159 mAh·g−1. In addition, the GCD curves of NVOH@GO-x exhibit virtually unchanged shapes at various current densities, whereas the voltage plateaus for NVOH almost vanish at 20.0 A·g−1 (Figs. 3c, S3g–i). After 2000 cycles at 5.0 A·g−1 (Figs. 3e, S4), NVOH, NVOH@GO-1, NVOH@GO-2, and NVOH@GO-3 electrodes deliver reversible capacities of 226, 299, 334, and 282 mAh·g−1, respectively, with 100% Coulombic efficiency and well-retained GCD curves. After 10,000 cycles at 20.0 A·g−1 (Fig. 3f), NVOH@GO-2 also exhibits remarkable cyclic reversibility with a high specific capacity of 238 mAh·g−1 and ~ 100% Coulombic efficiency. In addition, the NVOH@GO-2 electrode exhibits superior cyclic stability and cyclic capability with high mass loadings of 4 and 11 mg·cm−2 (Fig. S5a, b). At 50 and 100 mA·g−1, under deep charging/discharging conditions, the NVOH@GO-2 electrode delivers higher capacities of 469 and 430 mAh·g−1, respectively, with overlapping GCD curves (Fig. S5c–f). Furthermore, the 120 h standing result indicates that the NVOH@GO-2 electrode displays much lower self-discharge with 91% capacity retention after resting for 5 days (Fig. S5g, h), showing stable charge storage, which is a prerequisite for practical application of the NVOH@GO-2 electrode [59, 60]. The above results validate the high electrochemical reversibility, rapid reaction kinetics, and superior electrochemical stability of the optimized NVOH@GO-2 electrode, which can also be verified by comparing the electrochemical performance of the NVOH@GO-2 electrode in this work with other vanadium-based cathodes reported in the literatures (Table S1).

Fig. 3
figure 3

Electrochemical performance: a CV curves at 0.1 mV·s−1, b GCD curves at 0.5 A·g−1, and c GCD curves at different current densities of NVOH@GO-2; d rate performance and e cyclic performance at 5.0 A·g−1 of NVOH and NVOH@GO-x; f cyclic performance at 20 A·g−1 of NVOH@GO-2

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CV curves were measured at various scan rates to investigate the Zn2+ intercalation kinetics of the as-prepared electrodes. The redox peaks in the CV curves at 0.1 and 0.8 mV·s−1 of the NVOH@GO-2 electrode display narrower voltage gaps and higher peak current densities than those of the NVOH electrode (Figs. 4a, S6a), suggesting lower polarization and faster redox reaction kinetics in the NVOH@GO-2 electrode [20, 27, 51]. As the scan rates increase, the peak centers of the redox couples change significantly and the V4+/V3+ redox pairs are suppressed in NVOH (Fig. S7a), while the CV curves of the NVOH@GO-x electrodes remain similar, with slightly shifted peak positions (Figs. 4b, S7b, c), further validating the low polarization and fast electrochemical kinetics of the NVOH@GO-x electrodes [34, 61]. It is generally accepted that the relationship between current (i) and scan rates (ν) can be expressed by Eq. (2), where a and b are adjustable parameters. Values for b of 0.5 and 1 illustrate diffusion-controlled and surface-controlled pseudocapacitive processes for charge storage, respectively [34, 62]. The calculated b values of peaks 1–4 for NVOH and NVOH@GO-x range from 0.57 to 0.92 (Figs. 4c and S7d–f), which suggests that the redox reactions are controlled by a combination of capacitive-controlled and diffusion-controlled processes [35, 62]. The proportion of capacitive-controlled and diffusion-controlled capacities can be quantitatively determined by Eq. (3), where k1 and k2 are the percentage coefficients of capacitive and diffusion contribution in the total capacity, respectively. The fitted CV curves (shadow area) at 1 mV·s−1 in Figs. 4d, S7g–i demonstrate that the percentages of capacitive contributions are 73%, 69%, 56%, and 40% in NVOH, NVOH@GO-1, NVOH@GO-2, and NVOH@GO-3, respectively. Figure 4e shows the contribution ratios of diffusion-controlled and capacitive-controlled capacities in the total capacity at different scan rates. Although the capacitive contribution increases with increasing scan rate, the NVOH@GO-x electrodes display lower capacitive contribution ratios, which indicates that the NVOH-GO interface modulates the electrostatic interaction and accelerates ion diffusion into the NVOH framework [16, 35, 61].

$$i = a\nu^{b }$$
(2)
$$i = k_{{1}} \nu + k_{{2}} \nu^{{{1}/{2}}}$$
(3)
Fig. 4
figure 4

Electrochemical kinetics: a the fifth CV curves of NVOH and NVOH@GO-2 at 0.1 mV·s−1; b CV curves at various scan rates, c lgi versus lgv plots and d capacitive contribution (shadow region) to total current at 1.0 mV·s−1 of NVOH@GO-2; e individual contribution ratios of capacitive and diffusion-controlled behavior at various scan rates; f GITT profiles and calculated Zn2+ diffusion coefficient (D) during g discharging process and h charging process of NVOH and NVOH@GO-2

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EIS and GITT were conducted to further examine the reaction kinetics and solid-state diffusion dynamics, respectively. A semicircle in the high-frequency region and a slanted line in the low-frequency region are observed in all the Nyquist plots (Fig. S6b), which are related to the ohmic resistance (Re), charge-transfer resistance (Rct), and Warburg impedance (Zw). The Re and Rct values of the NVOH@GO-x electrodes are much lower than those of NVOH (Table S2), demonstrating the improved electrical conductivity and accelerated reaction kinetics of NVOH@GO-x electrodes [21, 41, 57]. The GITT profiles in Fig. 4f indicate that the NVOH@GO-2 electrode delivers less overpotential during charging/discharging, further indicating its fast kinetics. The ion diffusion coefficients can be calculated according to Eq. (4), where τ, mm, Vm, S, ΔEs, and ΔEτ are the current pulse time, moles, molar volume, electrode area, steady-state voltage change, and voltage change during the pulse, respectively [19, 35, 57]. The calculated diffusion coefficients of NVOH@GO-2 during charging and discharging are higher than those of NVOH (Fig. 4g, h), indicating the efficient diffusion dynamics of the NVOH@GO-2 electrode. The CV, EIS, and GITT results reveal that the electron transfer and the strong interaction between NVOH and GO, together with oxygen vacancies, improve the electronic conductivity and enhance the ion diffusion kinetics of the NVOH@GO-x electrodes, contributing to their superior electrochemical performance.

$$D = \frac{4}{\pi \tau }\left( {\frac{{m_{{\text{m}}} V_{\text{m}} }}{S}} \right)^{2} \left( {\frac{{\Delta E_{{\text{s}}} }}{{\Delta E_{\tau } }}} \right)^{2}$$
(4)

3.3 Electrochemical reaction mechanism

The electrochemical reaction mechanism of the NVOH@GO electrode was characterized using a series of ex-situ experiments. Ex-situ XRD patterns and the corresponding GCD curves are shown in Fig. 5a–c. At all investigated charging/discharging states, the diffraction peaks corresponding to (NH4)2V10O25 vanish, and new diffraction peaks emerging at 12.3°, 20.9°, 30.1°, 32.0°, 34.2°, and 36.5° can be indexed to Zn3V2O7(OH)2·2H2O (JCPDS No. 50-0570), which is generated by the coordination of Zn2+-intercalated V–O with H2O molecules, suggesting that Zn2+ and H2O are co-intercalated into NVOH@GO [20, 26, 63]. In addition, as can be observed in Fig. 5b, the (001) plane of Zn3V2O7(OH)2·2H2O shifts to lower 2θ values during the discharging process and moves back during the charging process, indicating the reversible Zn2+ (de)intercalation process of Zn3V2O7(OH)2·2H2O, which is consistent with the literatures [14, 64].

Fig. 5
figure 5

Electrochemical reaction mechanism of NVOH@GO-2 electrode: a, b ex-situ XRD patterns at selected states and c corresponding GCD curves of initial two cycles at 0.05 A·g−1; d Raman spectra at different states; XPS spectra of e Zn 2p, f V 2p, g O 1s and h N 1s at different states; i schematic diagram of reaction mechanism for NVOH@GO electrode; j V concentration in electrolyte after electrodes were immersed in 2 mol·L−1 Zn(CF3SO3)2 for 50 days

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To further confirm the reversible Zn2+ (de)intercalation of Zn3V2O7(OH)2·2H2O, ex-situ Raman spectroscopy and XPS were performed. Three Raman peaks at 248, 369, and 498 cm−1 are assigned to Zn2+–O2− symmetrical vibrations, two Raman peaks at 141 and 315 cm−1 are ascribed to V–O bending vibrations, and two Raman peaks at 806 and 866 cm−1 are assigned to V–O stretching vibration (Fig. 5d), suggesting the presence of Zn3V2O7(OH)2·2H2O [14, 21, 64, 65]. The Raman peak at 498 cm−1 for the fully charged electrode and the Raman peak at 315 cm−1 for the fully discharged electrode disappear because the extraction of Zn2+ weakens the Zn–O vibration and enhances the V–O vibration [14]. In both the fully discharged and charged states, obvious Zn peaks located at 1021.9 and 1045.0 eV ascribed to intercalated and surface-adsorbed Zn2+ are observed in Fig. 5e, indicating the presence of Zn-containing compounds and the reversible Zn2+ (de)intercalation of Zn3V2O7(OH)2·2H2O [14, 40, 64, 66, 67]. V5+ at 517.7 and 525.4 eV accompanied by V4+ at 517.2 and 524.4 eV are de-convoluted in the V 2p XPS spectra of the charged and discharged electrodes (Fig. 5f), and the charged electrode displays a higher V5+/V4+ ratio due to Zn2+ de-intercalation [17, 28, 29, 38]. The O 1s spectra consist of three parts (Fig. 5g), corresponding to V–O bonds (530.6 eV), H2O molecules (532.1 eV), and Zn–O bonds (533.0 eV), respectively [38, 64, 65]. In the discharged state, the O 1s XPS peak for H2O molecules increases owing to the co-intercalation of Zn2+ and H2O, while the weakened peak for V–O bonds can be ascribed to the insertion of Zn2+, in agreement with the ex-situ Raman results [23, 38]. The N 1s spectra exhibits two peaks at 402.1 and 400.5 eV (Fig. 5h), ascribed to protonated –NH4+– and –NH– segments, respectively, which implies that the intercalated NH4+ may be present as NH3 [19, 53]. Moreover, the intensity of the –NH4+– peak increases, and the N 1s XPS spectrum shifts to higher binding energies in the discharged state, indicating that NH4+ is involved in the redox reaction [16, 53]. The two charged electrodes display similar ex-situ Raman peaks and XPS spectra (Fig. 5d–h), indicating the high reversibility of the electrode in subsequent cycles. The ex-situ XRD, Raman, and XPS results reveal the charge storage mechanism of the NVOH@GO electrode, including a phase transition from (NH4)2V10O25 to Zn3V2O7(OH)2·2H2O during the first discharge process and a reversible Zn2+/H2O co-(de)intercalation of Zn3V2O7(OH)2·2H2O during the subsequent cycle, as described in Fig. 5i.

The structural stability of the electrode materials was evaluated by a static soaking experiment in 2 mol·L−1 Zn(CF3SO3)2 for 50 days and 2 mol·L−1 ZnSO4 for 10 days. Although the electrolyte soaked with NVOH and the NVOH@GO-2 electrode remains transparent and colorless during the static soaking experiment (Figs. 5j, S8), the V concentration in the electrolyte immersed with NVOH is much higher than that immersed with the NVOH@GO-2 electrode, suggesting that increased interaction between V and O in NVOH@GO-x induced by GO addition suppresses the dissolution of vanadium into the electrolyte, indicating the stability of the electrode and favoring its superior cyclic stability [38, 65].

4 Conclusion

In summary, a solution synthesis strategy was successfully employed to engineer oxygen-deficient (NH4)2V10O25·xH2O/GO composites, which exhibited satisfactory electrochemical performance as cathodes for ARZIBs. In particular, the optimized NVOH@GO delivered an extraordinary rate capability of 303 mAh·g−1 at 10 A·g−1 and an ultra-stable capacity of 238 mAh·g−1 after 10,000 cycles at 20 A·g−1. The superior electrochemical performance of the NVOH@GO composites can be attributed to the following favorable features: the encapsulation of (NH4)2V10O25·xH2O into GO enlarged the layer spacing and improved the electrical conductivity of NVOH, which provided spacious channels for Zn2+ (de)intercalation and accelerated ion diffusion, favoring fast electrochemical dynamics; the electron transfer and strong interaction between NVOH and GO provided additional electron transfer paths; abundant oxygen vacancies offered additional active sites for ion storage and jump sites for charge transfer, leading to a high-specific capacity; and inhibition of vanadium dissolution and limited self-discharge contributed to the superior cyclic stability. In addition, the electrochemical mechanism, including the phase transition reaction and subsequent Zn2+/H2O co-(de)intercalation process, was elucidated using ex-situ XRD, XPS, and Raman techniques. This work not only provides a strategy for the construction of V-based cathodes with dissolution inhibition, but also proposes a charge storage mechanism.