English

• In order to alleviate the energy crisis and utilize renewable energy, large scale energy storage system has received enormous attention, and it's urgent to develop a new type of low-cost energy storage device [1, 2]. Though commercial lithium-ion batteries (LIBs) have successfully dominated the market of portable products and electric vehicles, they have limited applications in the large scale energy storage system due to the uneven distribution and insufficient availability of lithium resources [3-5]. In this regard, potassium ion batteries (PIBs) hold the promise as complementary candidates to LIBs because of their abundant resource and low cost [6-9]. Furthermore, the low redox potential (−2.93 V vs. SHE) and weak Lewis acidity of K⁺ endow PIBs with high operating voltage and fast diffusion kinetics, respectively [10-12]. Nevertheless, the de/intercalation of large size of K⁺ (radius: 1.38 Å) is often accompanied by severe volume change, leading to unsatisfactory electrochemical performance, so it still remains great challenges to explore suitable electrode materials for high performance PIBs [13-15].

  To date, a variety of anode materials have been exploited for PIBs, including carbon-base materials [16-18], alloy-type materials [19, 20], organic compounds [21, 22], and transition metal oxide/chalcogenides materials [23-26], etc. Among them, transition metal oxides have been intensively studied due to their high theoretical capacity and cost effectiveness [27]. As a representative, V₂O₃ with open tunnel structure of the 3D framework can achieve fast de/intercalation of K⁺, which has gain much interest in recent years [28, 29]. Unfortunately, it is difficult to realize high capacity and good rate capability owing to its low electronic conductivity [29, 30]. Besides, the large size of K⁺ will cause structure collapse during cycling, resulting in poor cycle life. To tackle these issues, the common strategy is morphology engineering and incorporating with carbon-based materials to construct composites [31]. For example, a self-standing V₂O₃@PNCNFs electrode reported by Jiao's group shows a reversible capacity of 240 mAh/g at 50 mA/g and maintains 95.8% capacity retention after 500 cycles [28]. In addition, deficiency engineering is also an efficient method to enhance electronic conductivity of the materials [32-34]. Li et al. [24] demonstrated that constructing oxygen-rich defect amorphous shell on VO₂ surface can narrow the band gap of VO₂, boosting K⁺ storage capability. To the best of our knowledge, there is little research combining the above strategies to develop PIBs with excellent cycling stability and rate capability.

  Herein, oxygen-deficient V₂O₃ nanoparticles encapsulated in amorphous carbon shell (O_d-V₂O₃@C) are rationally designed and synthesized by a facile solvothermal method followed by calcination and reduction process. The as-prepared O_d-V₂O₃@C nanospheres are endowed with these following merits: (1) The nanostructure design guarantees the short diffusion distances and thus facilitates electron/K⁺ transfer. (2) The uniform amorphous carbon shell can enhance the electronic conductivity as well as maintain the structural stability. (3) The introduction of oxygen deficiency is capable of regulating the electronic structure of materials and greatly boosting the electrochemical performance. As a result, O_d-V₂O₃@C electrode achieved superior rate capability and excellent long-term stability (127.4 mAh/g after 1000 cycles at 2 A/g). Such a synergistic strategy by combining morphology and deficiency engineering provides an effective way to exploit transition metal oxide with high performance for energy storage application.

  O_d-V₂O₃@C was prepared by solvothermal method followed by calcination and reduction process [35]. Firstly, a homogeneous solution was obtained by dissolving 0.61 g of vanadyl acetylacetonate (VO(acac)₂) and 0.50 g of 1,4-benzenedicarboxylic acid (BDC) in 50 mL of absolute ethanol. Then it was transferred to a 100 mL Teflon-lined autoclave and heated at 180 ℃ for 12 h. After cooling down, the precipitate (V-MOF) was collected and washed with methanol for several times. The dried-down V-MOF powders were calcinated at 750 ℃ for 1 h under Ar atmosphere to get V₂O₃@C product. Finally, O_d-V₂O₃@C sample was obtained by annealing V₂O₃@C sample at 600 ℃ for 1 h under Ar/H₂ atmosphere.

  The crystal structures of the samples were identified by powder X-ray diffraction (XRD, Bruker D8 Advance, Cu Kα). The chemical valence of elements was investigated through X-ray photoelectron spectroscopy (XPS, Thermo Fisher Scientific). The morphological information was obtained from scanning electron microscopy (SEM, Hitachi SU8010) and transmission electron microscopy (TEM, Talos F200X). The Raman spectra were recorded on LabRAM HR Evolution system (514.5 nm Ar laser). Thermogravimetric analysis (TGA, STA449C) was used to test the carbon content of the samples under air flow from 30 ℃ to 700 ℃. The N₂ adsorption-desorption isotherms and pore distribution of the samples were determined with the Brunauer-Emmett-Teller method (BET, Autosorb-iQ). Electron Paramagnetic Resonance (EPR) test was utilized to characterized the oxygen vacancies.

  For the electrochemical measurements, CR2032 coin cells were used for the half-cell assembling with metallic potassium foil as counter electrode, 1 mol/L KFSI dissolved in EC/DEC as electrolyte, and glass fiber as separator. While for working electrode, the active materials were mixed with carbon black and polyvinylidene fluoride (PVDF) with mass ratio of 7:2:1 in an appropriate amount of N-methyl-2pyrrolidone solvent to obtain a uniform slurry. Then it was coat on the current collector (Cu foil) using doctor blade method. After drying up in vacuum at 80 ℃ for overnight, small discs with a diameter of 13 mm were punched as working electrode. Finally, Ar-filled glove box was used for the assemble of CR2032 coin cells. LAND CT2001A battery testing system (Wuhan, China) was used for galvanostatic charge-discharge (GCD) tests. For galvanostatic intermittent titration technique (GITT) tests, the activated coin cells were firstly charged/discharged at a pulse current density of 50 mA/g for 20 min and then held a relaxation duration of 3 h. CHI 660E electrochemical workstation was used for cyclic voltammetry tests (CV, 0.1–2.0 mV/s) and electrochemical impedance spectra (EIS, 0.01 Hz to 100 kHz).

  The V₂O₃@C composites are derived from the pyrolysis of V-MOF nanospheres that were prepared by solvothermal method (see schematic diagram of synthetic process in Fig. S1 in Supporting information). The SEM image in Fig. S2 (Supporting information) displays the spherical morphology of V-MOF with a diameter of 100–200 nm. After the calcination and reduction process at high temperature, the spherical morphology of V₂O₃@C and O_d-V₂O₃@C is well preserved (Figs. 1a and b, Fig. S3 in Supporting information). The TEM and HRTEM images further demonstrate the microstructure of the materials. It can be seen that V₂O₃ nanoparticles are embedded in amorphous carbon matrix (Figs. 1c and d). Such nano and porous structure guarantees short diffusion path and abundant reaction sites, thus promoting K⁺ reaction kinetics. The N₂ adsorption/desorption isotherms and pore distribution of the materials were conducted to verify the aforementioned microstructure. As displayed in Fig. S4 (Supporting information), a specific surface area of 270.6 m²/g with pore size range of 6–50 nm is measured for O_d-V₂O₃@C sample, showing a mesoporous character. Meanwhile, V₂O₃@C sample has a similar specific surface area (243.3 m²/g), which is consistent with the TEM result. Besides, obvious lattice fringes with d-spacings of 0.27 and 0.36 nm can be found in HRTEM images (Fig. 1d), corresponding to the d₁₀₂ and d₁₀₄ interlayer of V₂O₃, respectively. The SAED pattern of O_d-V₂O₃@C in Fig. 1e also clearly shows the existence of diffraction rings corresponding to plane (102), (104), (113) and (214), which is consistence with the XRD patterns. Furthermore, the EDX elemental mapping images presents the uniform distribution of V, O and C elements within the entire nanosphere (Fig. 1f).

  Figure 1

  

  Figure 1.  (a, b) SEM images, (c) TEM image, (d) HRTEM image, (e) SAED pattern, (f) EDX elemental mapping of O_d-V₂O₃@C materials.

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  The crystal structures of the samples were identified by XRD. As shown in Fig. 2a, the diffraction peaks of V₂O₃@C and O_d-V₂O₃@C can be indexed to V₂O₃ (PDF# 34–0187) without any impurity, suggesting the good crystallinity of as-prepared materials. The Raman spectra in Fig. 2b clearly shows two typical peaks (D band and G band) of the samples located at 1350 and 1600 cm⁻¹, further indicating the existence of carbon. Notably, the values of I_D/I_G ratio for V₂O₃@C and O_d-V₂O₃@C are 0.95 and 1.00, respectively, suggesting the dominant amorphous carbon component. The higher value of I_D/I_G ratio for O_d-V₂O₃@C implies the carbon coating layer contains more defects after high temperature reduction, which is beneficial for K⁺ adsorption [36]. Additionally, the carbon content of the composite can be quantitatively calculated according to the TGA curves in Fig. 2c. The result shows that carbon component of V₂O₃@C and O_d-V₂O₃@C are nearly 39.7% and 37.7%, respectively. The chemical compositions and valence of elements were studied by XPS. As shown in Fig. S5 (Supporting information), the XPS survey spectra further confirm the existence of V, O and C elements. Specifically, the V 2p spectrum of V₂O₃@C exhibits two typical peaks at 517.4 and 524.7 eV, corresponding to V 2p_3/2 and V 2p_1/2 of V³⁺, respectively. After reduction at high temperature, there are another two peaks emerging at 516.0 and 523.4 eV, which are referred to V 2p_3/2 and V 2p_1/2 of V²⁺, respectively, revealing the partial reduction of V element from 3+ to 2+ in O_d-V₂O₃@C (Fig. 2d) [37]. The O 1s spectra of V₂O₃@C and O_d-V₂O₃@C can be fitted into three peaks of V-O (530.5 eV), oxygen vacancies (531.6 eV), and C=O (533.8 eV). It is noteworthy that the intensity of peak at 531.6 eV apparently increases after the reduction, implying more oxygen defects exist in O_d-V₂O₃@C (Fig. 2e) [38]. The C 1s spectra of both samples exhibit the similar fitting result with three peaks at 284.8, 286.0 and 289.2 eV, which is ascribed to C—C, C—O and C=O bond, respectively (Fig. S6 in Supporting information). The existence of oxygen vacancies was further confirmed by EPR measurement. As shown in Fig. 2f, O_d-V₂O₃@C sample shows a stronger signal than V₂O₃@C at the g value of 2.003, suggesting more oxygen vacancies, which agrees well with XPS analysis. According to the materials characterization above, it is concluded that oxygen-deficient V₂O₃@C nanospheres were synthesized successfully.

  Figure 2

  

  Figure 2.  (a) XRD patterns, (b) Raman spectra, (c) TGA curves, (d, e) high resolution XPS spectra of (d) V 2p, (e) O 1s, and (f) EPR spectra of O_d-V₂O₃@C and V₂O₃@C materials.

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  To investigate the benefits of nanostructure and oxygen vacancies, the half-cells performance of as-prepared materials were evaluated at room temperature by the techniques of cyclic voltammetry (CV), galvanostatic charge-discharge (GCD), etc. As depicted in Fig. 3a, CV curves of O_d-V₂O₃@C electrode at 0.1 mV/s exhibit a broad reduction peak at ~0.5 V in the first cycle and then vanish in the subsequent cycles, which results from the formation of solid electrolyte interface (SEI). During the deintercalation process of K⁺, there is an oxidation peak at ~1.2 V. In the subsequent cycles, the nearly overlapped CV curves suggest the highly reversibility of O_d-V₂O₃@C electrode. At the same time, the CV curves of V₂O₃@C electrode exhibits the similar trend to the former (Fig. S7 in Supporting information). Fig. 3b displays the typical GCD profiles for the first three cycles of O_d-V₂O₃@C electrode. It delivers an initial discharge/charge capacities of 721.4/262.8 mAh/g, corresponding to a coulombic efficiency (CE) of 36.4%. However, the specific discharge and charge capacities of V₂O₃@C electrode are only 664.5 and 216.2 mAh/g, respectively, showing a lower CE (32.5%) (Fig. S8 in Supporting information). More importantly, O_d-V₂O₃@C electrode is capable of exhibiting better rate performance than V₂O₃@C. As illustrated in Figs. 3c and d, the reversible capacities of 262.8, 227.8, 201.5, 179.8, 156.9 mAh/g are achieved at 100, 200, 500, 1000 and 2000 mA/g, respectively. When the current density is switched back to 100 mA/g, a reversible capacity of 261.1 mAh/g is recovered and remains stable in following cycles. In contrast, V₂O₃@C electrode exhibits inferior specific capacities at corresponding current densities. Such superior rate capability of O_d-V₂O₃@C electrode is further evidenced by the charge/discharge voltage profiles (Fig. S9 in Supporting information). In addition to better rate performance, O_d-V₂O₃@C electrode also exhibits excellent cycling stability. Specifically, O_d-V₂O₃@C electrode maintains a specific capacity of 253 mAh/g after GCD at 200 mA/g for 100 cycles, which is superior to that of V₂O₃@C (208 mAh/g) (Fig. 3d). Even cycled at 2 A/g, a reversible capacity of 127.4 mAh/g is still retained after 1000 cycles without obvious capacity decay (Fig. 3e). Moreover, the SEM images of O_d-V₂O₃@C electrode after intensely cycling are depicted in Fig. S10 (Supporting information). Impressively, O_d-V₂O₃@C electrode could maintain structural integrity during cycling. This can be ascribed to uniform amorphous carbon shell can protect V₂O₃ nanoparticles from corrosion of the electrolyte, which guarantees the satisfactory cycling stability. Carbon matrix is essential to achieve high electronic conductivity and mitigate volume change during cycling, and ultimately to achieve higher capacity and better durability. With low carbon content, the beneficial spherical carbon matrix cannot be formed. In order to study the effects of carbon matrix on K-storage performance, a carbon matrix was prepared by dissolving O_d-V₂O₃@C powder in 1 mol/L HNO₃ solution and stirring for 24 h, washing with distilled water for several times and drying overnight in a vacuum oven at 80 ℃. The SEM images are shown in Figs. S11a and b (Supporting information), where it is evident that after removal of V₂O₃, the spherical morphology is still maintained. Besides, the XRD pattern confirms the (002) peak of pure carbon without any peak referred to V₂O₃ (Fig. S11c in Supporting information). According to the carbon loading (37.7%) in the composite and the stable specific capacity (185 mAh/g at 200 mA/g) of bare carbon, the capacity contribution from carbon in O_d-V₂O₃@C composite is estimated to be 69.7 mAh/g (Fig. S11d in Supporting information).

  Figure 3

  

  Figure 3.  (a) CV curves of O_d-V₂O₃@C electrode at 0.1 mV/s. (b) GCD curves of O_d-V₂O₃@C electrode at 100 mA/g for first three cycles. (c) Rate capability of O_d-V₂O₃@C and V₂O₃@C electrodes at various current density. (d) Cycling stability of O_d-V₂O₃@C and V₂O₃@C electrodes at 200 mA/g. (e) Long-term cycling life of O_d-V₂O₃@C electrode at 2 A/g.

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  To further interpret the better rate capability of O_d-V₂O₃@C electrode, the electrochemical kinetics were studied through CV, GITT and EIS measurements. As shown in Fig. 4a, the CV curves of O_d-V₂O₃@C electrode show distinct redox peaks at different scan rates with increasing peak intensity from 0.1 mV/s to 1.0 mV/s. Generally, the peak current (i) vs. scan rate (v) comply with the following equation (Eq. 1) [39]:

  (1)

  Figure 4

  

  Figure 4.  (a) CV curves at various scan rates of O_d-V₂O₃@C electrode. (b) log(i) vs. log(v) plots and corresponding linear relationship of O_d-V₂O₃@C electrode. (c) Normalized contribution ratios of capacitive contribution at various scan rates of O_d-V₂O₃@C electrode. (d) GITT curves at a pulse current density of 50 mA/g of O_d-V₂O₃@C and V₂O₃@C electrodes. (e) The calculated K⁺ diffusion coefficients upon discharge and charge process of O_d-V₂O₃@C and V₂O₃@C electrodes.

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  where the value of b can evaluate the charge storage behavior of the electrode. As depicted in Fig. 4b, the b values of O_d-V₂O₃@C electrode are close to 1, indicating the capacity is dominated by pseudocapacitance. Besides, the percentages of capacitance and diffusion contributions can be quantitatively calculated according to Eq. 2 [39]:

  (2)

  where k₁v and k₂v^1/2 represent the capacitive contribution fraction and the diffusion one, respectively. As a result, the capacitive contributions occupied ~86.8% of the total capacity at the scan rate of 1.0 mV/s (Fig. S12 in Supporting information). Besides, the ratio of capacitive contribution has a gradual upward tendency with the increase of scan rates, which benefits the high rate performance (Fig. 4c). Although V₂O₃@C electrode also exhibits dominant surface-controlled capacitive behavior (Fig. S13 in Supporting information), the proportion of capacitive contribution is lower than that of O_d-V₂O₃@C electrode. Therefore, the introduction of oxygen vacancies promotes the capacitance-type charge storage, which accounts for the superior rate capability of O_d-V₂O₃@C electrode. Fig. 4d presents the voltage response of the electrodes under a pulse current density of 50 mA/g. Since the voltage is linearly proportional to τ (Fig. S14b in Supporting information), the K⁺ diffusion coefficient (D_K⁺) can be obtained by Eq. 3 [40]:

  (3)

  where τ is duration of pulse current, n_m is the number of moles, ΔE_S and ΔE_τ are the voltage variation shown in Fig. S14a (Supporting information). The calculated D_K⁺ for O_d-V₂O₃@C and V₂O₃@C electrodes upon discharge/charge process are shown in Fig. 4e. Apparently, O_d-V₂O₃@C electrode exhibits a much higher diffusion coefficient during the whole process, demonstrating the fast electrochemical kinetics. According to the equivalent circuit model (Fig. S15b in Supporting information), impedance spectra of V₂O₃@C and O_d-V₂O₃@C electrodes before cycling can be reasonably fitted and depicted in Fig. S15a (Supporting information). As summarized in Table S1 (Supporting information), the values of Ohm impedance (R_s) and charge transfer resistance (R_ct) of O_d-V₂O₃@C electrode are 3.2 and 1501 Ω, respectively, which are apparently lower than that of V₂O₃@C electrode (6.5 and 2361 Ω). These results suggest the fast reaction kinetics of O_d-V₂O₃@C electrode, leading to the better rate capability.

  Moreover, four probe method and density functional theory (DFT) calculation is conducted to further understand the impacts on the electronic property of V₂O₃ after introducing oxygen defects. In order to explore the effects of various oxygen vacancy concentrations on electrical conductivity, O_d-V₂O₃@C materials with different H₂ treatment times have been synthesized, and their electrical conductivity were measured by four probe method. As shown in Fig. S16 (Supporting information), the electrical conductivity gradually increases with the extension of treatment times, but then decreases when treatment time is too long. It can conclude that appropriate amount of oxygen vacancy has great influence on the electronic structure of V₂O₃@C. The open tunnel structure of O_d-V₂O₃@C materials is illustrated in Fig. 5a. As shown in Fig. 5b, the density of states (DOS) of both samples clearly illustrate that the conduction band come across the Fermi levels (0 eV), unveiling the metallic properties of them. Interestingly, the intensity of DOS for O_d-V₂O₃ is substantially higher than that of V₂O₃, which proves the enhanced electronic conductivity of O_d-V₂O₃. In addition, the band structure also reveals the same conclusion. As shown in Figs. 5c and d, O_d-V₂O₃ possesses a narrower band gap (2.35 eV) than that of V₂O₃ (2.59 eV), thereby promoting the charge transfer [41, 42].

  Figure 5

  

  Figure 5.  (a) Structural illustration of O_d-V₂O₃. (b) DOS of V₂O₃ and O_d-V₂O₃. The band structure of (c) V₂O₃ and (d) O_d-V₂O₃.

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  In conclusion, a rational and feasible approach has been proposed to synthesize O_d-V₂O₃@C anode materials with high performance. The MOF derived carbon coating uniformly wrapping V₂O₃ nanoparticles can shorten the diffusion distances to enhance the electronic conductivity of materials, as well as alleviate the large volume variation and ensure structural integrity. Moreover, the as-prepared O_d-V₂O₃@C materials with rich oxygen defects is capable of tuning the electronic structure of V₂O₃ and greatly boosting K⁺ storage capability. When applied as anode for KIBs, O_d-V₂O₃@C electrode delivers a high capacity of 127.4 mAh/g after cycling for 1000 cycles at high current density of 2 A/g, exhibiting outstanding electrochemical performance. This study provides an effective way for developing advanced anode materials for KIBs.

  Declaration of competing interest

  There are no conflicts to declare.

  Acknowledgments

  We gratefully acknowledge the financial support from the National Natural Science Foundation of China (Nos. 51922042 and 51872098) and Fundamental Research Funds for Central Universities, China (No. 2020ZYGXZR074) and the Scientific and Technological Plan of Qingyuan City, China (2019DZX008).

  Supplementary materials

  Supplementary material associated with this article can be found, in the online version, at doi:10.1016/j.cclet.2022.03.095.

  
  
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  Figure 1  (a, b) SEM images, (c) TEM image, (d) HRTEM image, (e) SAED pattern, (f) EDX elemental mapping of O_d-V₂O₃@C materials.

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  Figure 2  (a) XRD patterns, (b) Raman spectra, (c) TGA curves, (d, e) high resolution XPS spectra of (d) V 2p, (e) O 1s, and (f) EPR spectra of O_d-V₂O₃@C and V₂O₃@C materials.

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  Figure 3  (a) CV curves of O_d-V₂O₃@C electrode at 0.1 mV/s. (b) GCD curves of O_d-V₂O₃@C electrode at 100 mA/g for first three cycles. (c) Rate capability of O_d-V₂O₃@C and V₂O₃@C electrodes at various current density. (d) Cycling stability of O_d-V₂O₃@C and V₂O₃@C electrodes at 200 mA/g. (e) Long-term cycling life of O_d-V₂O₃@C electrode at 2 A/g.

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  Figure 4  (a) CV curves at various scan rates of O_d-V₂O₃@C electrode. (b) log(i) vs. log(v) plots and corresponding linear relationship of O_d-V₂O₃@C electrode. (c) Normalized contribution ratios of capacitive contribution at various scan rates of O_d-V₂O₃@C electrode. (d) GITT curves at a pulse current density of 50 mA/g of O_d-V₂O₃@C and V₂O₃@C electrodes. (e) The calculated K⁺ diffusion coefficients upon discharge and charge process of O_d-V₂O₃@C and V₂O₃@C electrodes.

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  Figure 5  (a) Structural illustration of O_d-V₂O₃. (b) DOS of V₂O₃ and O_d-V₂O₃. The band structure of (c) V₂O₃ and (d) O_d-V₂O₃.

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