Synthesis and Electrochemical Performance of V6O13 Nanosheets Film Cathodes for LIBs

V6O13 thin films were deposited on indium-doped tin oxide (ITO) conductive glass by a concise low-temperature liquid-phase deposition method and through heat treatment. The obtained films were directly used as electrodes without adding any other media. The results indicate that the film annealed at 400 °C exhibited an excellent cycling performance, which remained at 82.7% of capacity after 100 cycles. The film annealed at 400 °C with diffusion coefficients of 6.08 × 10−12 cm2·s−1 (Li+ insertion) and 5.46 × 10−12 cm2·s−1 (Li+ extraction) in the V6O13 film electrode. The high diffusion coefficients could be ascribed to the porous morphology composed of ultrathin nanosheets. Moreover, the film endured phase transitions during electrochemical cycling, the V6O13 partially transformed to Li0.6V1.67O3.67, Li3VO4, and VO2 with the insertion of Li+ into the lattice, and Li0.6V1.67O3.67, Li3VO4, and VO2 partially reversibly transformed backwards to V6O13 with the extraction of Li+ from the lattice. The phase transition can be attributed to the unique structure and morphology with enough active sites and ions diffusion channels during cycles. Such findings reveal a bright idea to prepare high-performance cathode materials for LIBs.


Introduction
The increasing requirement for energy storage devices and novel clean energy resources has generated unparalleled research interest with the development of society [1][2][3][4].
In the recent few decades, lithium-ion batteries (LIBs) have been extensively applied in portable electronic devices and electric vehicles due to high energy density, power density, and long cyclic performance [5][6][7][8][9]. The electrode materials dominate the performance of LIBs [10]. Therefore, developing optimized electrode materials is key to improving the properties of LIBs. Layered vanadate oxides such as V 6 O 13 , V 2 O 5 , V 3 O 7 , VO 2 (B), and V 3 O 8 have been considered as a sort of applicable material because of the multiple ions oxidation states, high capacity, and layered crystal structure for ions intercalation [11][12][13][14].
Mixed-valence V 6 O 13 has been considered as a potential candidate for LIBs, which is a desired material compared with other vanadium oxides [15,16]. According to reports, it has a high theoretical capacity of 417 mA h g −1 , corresponding to 8 Li per formula unit (vanadium being reduced from an average charge of +4.33 to +3), and a theoretical energy density of 890 W h kg −1 [17][18][19][20]. The excellent performance of V 6 O 13 is attributed to its open structure, which is built up of single and double V-O layers with vanadium ions in multiple valence states (V 4+ and V 5+ ) [11,21]. Herein, the unique structure provides more ions diffusion paths and sufficient active sites, resulting in a high specific capacity and good cyclic performance.
However, there still exists plenty of possibilities for improvement in the performance of V 6 O 13 . Some achievements such as morphological regulation, structural adjustment, and so forth have been generally proposed as approaches, since the morphology and

Preparation of Films
All experimental drugs and reagents used in this experiment were analytically pure and could be used directly without further purification.
The V 6 O 13 films on ITO conductive glass were prepared via a low-temperature liquidphase deposition method and heat treatment under N 2 atmosphere. A deposition solution was prepared by dissolving 0.1630 g vanadium sulfate and 0.0840 g lactic acid in 50 mL deionized water. The deposition solution was magnetically stirred for 2 h in air at room temperature. After stirring, the pH was adjusted to 4.0 with ammonia.
A piece of ITO conductive glass (2 × 2 cm 2 , electrical resistance of 6 Ω/ ) was ultrasonically cleaned in detergent and deionized water for several times. After that, the washed ITO substrate was put face down horizontally into the deposition solution in the beaker, which was placed in a water bath at 90 • C for 3 days. When the reaction was completed, the substrate was taken out, rinsed slowly with deionized water, cooled naturally to room temperature, and dried to obtain the desired deposited precursor films. The precursor films were annealed under nitrogen atmosphere with a velocity of flow of 50 mL/min at 350 • C, 400 • C, and 450 • C at a rate of 2 • C/min for 2 h. The annealed Materials 2022, 15, 8574 3 of 13 films were obtained and named N−350, N−400, and N−450 to represent the films at the corresponding annealing temperature.

Materials Characterizations
The X-ray diffraction (XRD) patterns of the films were characterized by X-ray diffraction instrumentation (Rigaku Smart Lab 9kw, Tokyo, Japan) using Cu Kα radiation (λ = 1.5406 Å). Raman spectrum (Renishaw, England) was collected to obtain structural information. The further structural information of annealed films was measured by infrared spectroscopy (Nicolet 6700, Waltham, MA, USA). The microscopic morphology of the sample surface was examined with a scanning electron microscope (JEOL JSM7500F, Akishima, Japan). The chemical bonding states were characterized by X-ray photoelectron spectroscopy (Thermo ESCALAB 250Xi, Waltham, MA, USA) measurement with Al Kα source. To avoid the effects of oxide layers for the film exposed to air, etching was carried out prior to XPS testing.

Electrochemical Characterizations
The properties of films were tested in a three-electrode system with a CHI-604E electrochemical station (CH Instruments Inc., Shanghai, China) using 1M LiClO 4 /PC as the electrolyte. The prepared films were used directly as the working electrode and the effective area was 0.6 cm 2 , the platinum plate was used as the counter electrode, and the Ag/AgCl electrode in KCl saturated solution was used as the reference electrode. The three-electrode system was filled with nitrogen flow for 20 min to remove oxygen. Cyclic voltammetry (CV) was performed to investigate the electrochemical properties of the thin films. CV data were collected between −1 V and 0.4 V at different scanning rates of 1-10 mV·s −1 .

Characterization of the Obtained Films
The XRD patterns of the crystal phase structure at different annealing temperatures are shown in Figure 1. The deposited film possesses a strong peak at 9.3 • attributed to the hydrated phase of (NH 4 ) 8 11 films through heat treatment.
The FTIR spectra of the structural information are shown in Figure 2. The FTIR spectra of the films are dominated by strong absorption in the 1020-525 cm −1 region and these absorption peaks are associated with the vibrations of the bonds between V and O [3]. The bands between 1000-900 cm −1 can be attributed to V=O stretching vibrations [3,26]. The band at 1406 cm −1 is the bending pattern of N-H vibrations [27,28], which only presents in the N−350 film, representing that NH 4 + fully disappears when the annealing temperature increases, which is consistent with the above XRD results. The absorption peak is produced at around 875 cm −1 pointing to V-O stretching vibrations (shorter V-O bonds) [29]. The band at 714 cm −1 is a V-O-V stretching vibration due to bridging oxygen bonds (longer V-O bonds) [28,30]. There are no other absorption peaks, indicating almost no other impurities in the films. These results are consistent with the XRD results.
The Raman spectra of the films are displayed in Figure 3. The band at 993 cm −1 represents the V-O stretching vibration, while the one at 687 cm −1 can be ascribed to the V-O-V stretching vibration [31,32]. The peak at 523 cm −1 is attributed to the triple coordination oxygen (V 3 -O) stretching mode [30]. The peaks at 406 cm −1 and 281 cm −1 are due to the bending vibration of the V=O bond [33,34]. The additional low-frequency peaks at 139 and 189 cm −1 correspond to the stretching pattern of (V 2 O 2 ) n [35]. These peaks are The FTIR spectra of the structural information are shown in Figure 2. The FTIR spectra of the films are dominated by strong absorption in the 1020-525 cm −1 region and these absorption peaks are associated with the vibrations of the bonds between V and O [3]. The bands between 1000-900 cm −1 can be attributed to V=O stretching vibrations [3,26]. The band at 1406 cm −1 is the bending pattern of N-H vibrations [27,28], which only presents in the N−350 film, representing that NH4 + fully disappears when the annealing temperature increases, which is consistent with the above XRD results. The absorption peak is produced at around 875 cm −1 pointing to V-O stretching vibrations (shorter V-O bonds) [29]. The band at 714 cm −1 is a V-O-V stretching vibration due to bridging oxygen bonds (longer V-O bonds) [28,30]. There are no other absorption peaks, indicating almost no other impurities in the films. These results are consistent with the XRD results.   The FTIR spectra of the structural information are shown in Figure 2. The FTIR spectra of the films are dominated by strong absorption in the 1020-525 cm −1 region and these absorption peaks are associated with the vibrations of the bonds between V and O [3]. The bands between 1000-900 cm −1 can be attributed to V=O stretching vibrations [3,26]. The band at 1406 cm −1 is the bending pattern of N-H vibrations [27,28], which only presents in the N−350 film, representing that NH4 + fully disappears when the annealing temperature increases, which is consistent with the above XRD results. The absorption peak is produced at around 875 cm −1 pointing to V-O stretching vibrations (shorter V-O bonds) [29]. The band at 714 cm −1 is a V-O-V stretching vibration due to bridging oxygen bonds (longer V-O bonds) [28,30]. There are no other absorption peaks, indicating almost no other impurities in the films. These results are consistent with the XRD results.   V stretching vibration [31,32]. The peak at 523 cm −1 is attributed to the triple coordination oxygen (V3-O) stretching mode [30]. The peaks at 406 cm −1 and 281 cm −1 are due to the bending vibration of the V=O bond [33,34]. The additional low-frequency peaks at 139 and 189 cm −1 correspond to the stretching pattern of (V2O2)n [35]. These peaks are in agreement with the results of previous investigations. The vibrational absorption peaks are observed at similar locations of the films, indicating that the compositions are basically the same. The Raman results are constant with the IR results.

Electrochemical Performance Study
CV curves of the film electrodes at different scan rates within the range of −1 to 0.4 V are shown in Figure 5. The CV curves of the N−400 and N−450 show a pair of obvious redox peaks and two pairs of weak redox peaks, indicating a multi-step insertion/extraction of the Li + in V6O13 electrode. For the N−400, the reduction peaks are observed at about −0.38, −0.67, and −0.82 V, corresponding to the insertion of Li + into the V6O13 electrode (V 6 O 13 + xLi + + xe − = Li X V 6 O 13 ). The oxidation peaks at about −0.21 V, −0.12 V, and −0.55 V correspond to the extraction of Li + (Li x V 6 O 13 − yLi + − ye − = Li x−y V 6 O 13 ). For the N-450, the oxidation peaks are located at −0.16 V, −0.06 V, and −0.44 V, and the reduction peaks are located at −0.40, 0.63, and 0.80 V, respectively. A tiny new reduction peak at around −0.61 V becomes remarkable with the decrease of the scan rate, which may be caused by a more adequate response due to the reduced scan rate. Comparing the CV curves of the N−400 and N−450, the current of the N−400 is larger than N−450. This can be attributed to the morphology and crystallinity determined by the annealing temperature, the crystallinity annealed at 450 °C becomes complete, and the lithium ions are difficult to in-

Electrochemical Performance Study
CV curves of the film electrodes at different scan rates within the range of −1 to 0.4 V are shown in Figure 5. The CV curves of the N−400 and N−450 show a pair of obvious redox peaks and two pairs of weak redox peaks, indicating a multi-step insertion/extraction of the Li + in V 6 O 13 electrode. For the N−400, the reduction peaks are observed at about −0.38, −0.67, and −0.82 V, corresponding to the insertion of Li + into the V 6 O 13 electrode (V 6 O 13 + xLi + + xe − = Li X V 6 O 13 ). The oxidation peaks at about −0.21 V, −0.12 V, and −0.55 V correspond to the extraction of Li + (Li x V 6 O 13 − yLi + − ye − = Li x−y V 6 O 13 ). For the N-450, the oxidation peaks are located at −0.16 V, −0.06 V, and −0.44 V, and the reduction peaks are located at −0.40, 0.63, and 0.80 V, respectively. A tiny new reduction peak at around −0.61 V becomes remarkable with the decrease of the scan rate, which may be caused by a more adequate response due to the reduced scan rate. Comparing the CV curves of the N−400 and N−450, the current of the N−400 is larger than N−450. This can be attributed to the morphology and crystallinity determined by the annealing temperature, the crystallinity annealed at 450 • C becomes complete, and the lithium ions are difficult to insert/extract. In addition, a slight shift of peak position is observed with the increasing scan rate, which is mainly due to the polarization effect [36,37]. From the CV curves at different scan rates, the diffusion coefficient of Li + in the V 6 O 13 film electrode would be calculated from the peak current (I p ) and scan rate. In this equation, Ip is the peak current (A); A is the effective area of the prepared film electrode (cm 2 ); n is the number of electrons transferred in the redox reaction the process (here, for lithium batteries, the fixed value is 1); D is the diffusion coe of lithium ions (cm 2 s −1 ); C is the number of electrons in the active ion of the elec (the concentration of lithium ions in the electrolyte) (mol/cm 3 ); and v is the scan rat From Figure 6 can be seen that the Ip varies linearly with v 1/2 , the linear relationshi cates that the reaction is a diffusion-controlled process and the current in the redo cess is limited by the diffusion of ions of the electrode surface. The calculated di coefficients of Li + in the N−400 and N−450 are listed below in Table 1. The diffusio ficients of Li + are same as other literatures [3,38,39]. The Li + diffusion coefficient of is larger than N−450 for both the anode and cathode, suggesting that the diffusion of the N−400 is easier than the N−450, which is due to the surface effect of the t nanosheets. Moreover, the diffusion coefficient of the cathode is significantly high anode, which means it is easier to embed lithium ions than to remove them and this may lead to lithium ion retention in the interlayer. Hence, the lithium ions may pa retained in the interlayer to cause capacity fading. The retention of lithium ions m one of the reasons for the polarization of the CV curves.   In this equation, Ip is the peak current (A); A is the effective area of the prepare film electrode (cm 2 ); n is the number of electrons transferred in the redox reaction the process (here, for lithium batteries, the fixed value is 1); D is the diffusion coe of lithium ions (cm 2 s −1 ); C is the number of electrons in the active ion of the elec (the concentration of lithium ions in the electrolyte) (mol/cm 3 ); and v is the scan rat From Figure 6 can be seen that the Ip varies linearly with v 1/2 , the linear relationshi cates that the reaction is a diffusion-controlled process and the current in the red cess is limited by the diffusion of ions of the electrode surface. The calculated di coefficients of Li + in the N−400 and N−450 are listed below in Table 1. The diffusio ficients of Li + are same as other literatures [3,38,39]. The Li + diffusion coefficient of is larger than N−450 for both the anode and cathode, suggesting that the diffusio of the N−400 is easier than the N−450, which is due to the surface effect of the nanosheets. Moreover, the diffusion coefficient of the cathode is significantly high anode, which means it is easier to embed lithium ions than to remove them and thi may lead to lithium ion retention in the interlayer. Hence, the lithium ions may pa retained in the interlayer to cause capacity fading. The retention of lithium ions m one of the reasons for the polarization of the CV curves.  In this equation, I p is the peak current (A); A is the effective area of the prepared V 6 O 13 film electrode (cm 2 ); n is the number of electrons transferred in the redox reaction during the process (here, for lithium batteries, the fixed value is 1); D is the diffusion coefficient of lithium ions (cm 2 s −1 ); C is the number of electrons in the active ion of the electrolyte (the concentration of lithium ions in the electrolyte) (mol/cm 3 ); and v is the scan rate (V/s). From Figure 6 can be seen that the I p varies linearly with v 1/2 , the linear relationship indicates that the reaction is a diffusion-controlled process and the current in the redox process is limited by the diffusion of ions of the electrode surface. The calculated diffusion coefficients of Li + in the N−400 and N−450 are listed below in Table 1. The diffusion coefficients of Li + are same as other literatures [3,38,39]. The Li + diffusion coefficient of N−400 is larger than N−450 for both the anode and cathode, suggesting that the diffusion of Li + of the N−400 is easier than the N−450, which is due to the surface effect of the thinner nanosheets. Moreover, the diffusion coefficient of the cathode is significantly higher than anode, which means it is easier to embed lithium ions than to remove them and this result may lead to lithium ion retention in the interlayer. Hence, the lithium ions may partly be retained in the interlayer to cause capacity fading. The retention of lithium ions may be one of the reasons for the polarization of the CV curves. The cycling performance of films is examined at a scan rate of 10 mV/s from −1 to 0.4 V in Figure 7. The CV curves of the N−400 similarly display one obvious pair and two in-distinctive pairs of redox peaks, which is consistent with the results in Figure 5. The integral area of CV curves can represent capacity to some extent (C = ( idV)/2vV). For the N−400, the CV curves for the first 50 cycles are largely overlapping, with little capacity fading, the peak area starts to fade slowly for the next 50 cycles. The integral area of the CV curves maintains 83% after 100 cycles. The main oxidation peak located at −0.22 V stays at the same position for the 100 cycles, while the main reduction peak located near −0.65 V continuously shifts toward the right and locates at −0.59 V at the end of the 100th cycle. For the N−450, the integral area of CV curves gradually decrease, which is maintained at 61% after 100 cycles. The main oxidation peak at −0.13 V shifts gradually toward −0.22 V at the first 10 cycles, and the peak position keeps steady at −0.22 V for the next 90 cycles. The main reduction peak located at −0.63 V continuously shifts toward −0.59 V, which is similar to the N−400. The displacement of redox peaks can be attributed to some irreversible phase transitions during cycling. Overall, the electrochemical capacity and cycling performance of the N−400 is better than the N−450, which is ascribed to the morphology, crystalline, and phase transitions.  The cycling performance of films is examined at a scan rate of 10 mV/s from − V in Figure 7. The CV curves of the N−400 similarly display one obvious pair and distinctive pairs of redox peaks, which is consistent with the results in Figure 5. Th gral area of CV curves can represent capacity to some extent ( = (∫ )/2 ). N−400, the CV curves for the first 50 cycles are largely overlapping, with little c fading, the peak area starts to fade slowly for the next 50 cycles. The integral area CV curves maintains 83% after 100 cycles. The main oxidation peak located at − stays at the same position for the 100 cycles, while the main reduction peak locate −0.65 V continuously shifts toward the right and locates at −0.59 V at the end of th cycle. For the N−450, the integral area of CV curves gradually decrease, which is tained at 61% after 100 cycles. The main oxidation peak at −0.13 V shifts gradually −0.22 V at the first 10 cycles, and the peak position keeps steady at −0.22 V for the cycles. The main reduction peak located at −0.63 V continuously shifts toward − which is similar to the N−400. The displacement of redox peaks can be attributed t irreversible phase transitions during cycling. Overall, the electrochemical capaci cycling performance of the N−400 is better than the N−450, which is ascribed to th phology, crystalline, and phase transitions. The fitted and normalized electrochemical impedance test is used to further ev the electrochemical of the as-prepared V6O13 film electrode. The Figure 8 shows an alent circuit model according to the simulation. The model commonly employed studies consists of serial resistance (Rs), charge transfer (Rct), a Warburg diffusion e (Wo), and capacitive element (CPE). Generally, the impedance spectra can be divid a semicircle in the high-frequency range that is assigned to the charge-transfer res at the electrode-electrolyte interface and a straight line in the low-frequency ran implies the Li + diffusion-controlled process in the solid electrode [40,41]. The The fitted and normalized electrochemical impedance test is used to further evaluate the electrochemical of the as-prepared V 6 O 13 film electrode. The Figure 8 shows an equivalent circuit model according to the simulation. The model commonly employed in LIB studies consists of serial resistance (R s ), charge transfer (R ct ), a Warburg diffusion element (W o ), and capacitive element (CPE). Generally, the impedance spectra can be divided into a semicircle in the high-frequency range that is assigned to the charge-transfer resistance at the electrode-electrolyte interface and a straight line in the low-frequency range that implies the Li + diffusion-controlled process in the solid electrode [40,41]. The results demonstrate that the R ct value of the N−400 is smaller than the N−450, indicating that the electrochemical reaction occurs more easily in the N−400 film. The reason can also be attributed to the morphology and crystallinity. Meanwhile, it is obvious that the R ct value of both films decreases after cycling, which may be relevant to the gradual activation of the film electrodes that can expose more active sites after cycling, and phase transitions that occur during cycling.
Materials 2022, 15, x FOR PEER REVIEW 9 of attributed to the morphology and crystallinity. Meanwhile, it is obvious that the Rct val of both films decreases after cycling, which may be relevant to the gradual activation the film electrodes that can expose more active sites after cycling, and phase transitio that occur during cycling. To verify the morphology of the prepared films after 100 cycles, the SEM images a shown in Figure 9. The edges of the nanosheets of the N−400 have changed from flat petal nanosheets to irregular nanosheets with cut edges, and the surface is rougher an some nanoparticles have grown on the surface. Figure 9c,d shows that the edges of t nanosheets have changed from smooth petal-shaped to rectangular stacked nanoshee with a slightly rougher surface than before. In summary, the overall morphology of t films was basically unchanged, but there was also a little self-aggregation and crushin which means that the films were cyclically stable. To verify the crystal phase structure and conjectures of the N-400 film during cycl the ex-situ XRD patterns at different charge/discharge states are shown in Figure 10. Com paring them with the N-400 film before cycling, there is an obvious difference. After t To verify the morphology of the prepared films after 100 cycles, the SEM images are shown in Figure 9. The edges of the nanosheets of the N−400 have changed from flatter petal nanosheets to irregular nanosheets with cut edges, and the surface is rougher and some nanoparticles have grown on the surface. Figure 9c,d shows that the edges of the nanosheets have changed from smooth petal-shaped to rectangular stacked nanosheets with a slightly rougher surface than before. In summary, the overall morphology of the films was basically unchanged, but there was also a little self-aggregation and crushing, which means that the films were cyclically stable. attributed to the morphology and crystallinity. Meanwhile, it is obvious that the Rct value of both films decreases after cycling, which may be relevant to the gradual activation of the film electrodes that can expose more active sites after cycling, and phase transitions that occur during cycling. To verify the morphology of the prepared films after 100 cycles, the SEM images are shown in Figure 9. The edges of the nanosheets of the N−400 have changed from flatter petal nanosheets to irregular nanosheets with cut edges, and the surface is rougher and some nanoparticles have grown on the surface. Figure 9c,d shows that the edges of the nanosheets have changed from smooth petal-shaped to rectangular stacked nanosheets with a slightly rougher surface than before. In summary, the overall morphology of the films was basically unchanged, but there was also a little self-aggregation and crushing, which means that the films were cyclically stable. To verify the crystal phase structure and conjectures of the N-400 film during cycles, the ex-situ XRD patterns at different charge/discharge states are shown in Figure 10. Comparing them with the N-400 film before cycling, there is an obvious difference. After the  To verify the crystal phase structure and conjectures of the N-400 film during cycles, the ex-situ XRD patterns at different charge/discharge states are shown in Figure 10. Comparing them with the N-400 film before cycling, there is an obvious difference. After the first discharge process (Li + insertion), the peak of V 6  Furthermore, the peak intensity of Li0.6V1.67O3.67, Li3VO4 becomes stronger and the peak at 25.3° of V6O13 basically disappears again, which is due to the insertion of Li + . Then, after the hundredth charge process (Li + extraction), the peak intensity of Li0.6V1.67O3.67, Li3VO4, and VO2 decreases and the peak at 25.3° of V6O13 appears again. The ex-situ XRD results indicate there are phase transitions during Li + insertion and extraction. The process evolution mechanism is speculated as follows: Firstly, the phase starts to change when Li + are inserted after the first discharge, the new phases of Li0.6V1.67O3.67 and Li3VO4 appear. That process is accompanied by a phase transition from V6O13 to VO2, the degree of this phase transition is very small, so the transformation of the crystal phase structure has not yet been achieved. After the first charge process, the The ex-situ XRD results indicate there are phase transitions during Li + insertion and extraction. The process evolution mechanism is speculated as follows: Firstly, the phase starts to change when Li + are inserted after the first discharge, the new phases of Li 0.6 V 1.67 O 3.67 and Li 3 VO 4 appear. That process is accompanied by a phase transition from V 6 O 13 to VO 2 , the degree of this phase transition is very small, so the transformation of the crystal phase structure has not yet been achieved. After the first charge process, the transformation from V 6 O 13 to VO 2 still has not yet been completed. The Li 0.6 V 1.67 O 3.67 and Li 3 VO 4 still exist but the peak intensity becomes weak because the phase transition is not fully reversible with the extraction of Li + . Then after the hundredth discharge process, the transition from V 6 O 13 to VO 2 has been completed, so the peak of V 6 O 13 is barely detected. When Li + are extracted again after the hundredth charge process, the VO 2 partly changes backward to V 6 O 13 because of incomplete reversibility. Additionally, other phases of Li 0.6 V 1.67 O 3.67 and Li 3 VO 4 exhibit regular changes with Li + insertion/extraction, and the peak intensity increases with ion insertion and decreases with ion extraction. Figure 11 shows the Raman spectra after 100 cycles. The Raman peaks of the N−400 and N−450 are basically same, but there is an obvious difference with the film before cycling. A red shift of these Raman peaks both in N−400 and N−450 is observed after 100 cycles. The shift is attributed to the Li + between the V-O layers, which results in a lattice expansion, and the split peak at 139 cm −1 may also related to it [30]. transformation from V6O13 to VO2 still has not yet been completed. The Li0.6V1.67O3.67 and Li3VO4 still exist but the peak intensity becomes weak because the phase transition is not fully reversible with the extraction of Li + . Then after the hundredth discharge process, the transition from V6O13 to VO2 has been completed, so the peak of V6O13 is barely detected. When Li + are extracted again after the hundredth charge process, the VO2 partly changes backward to V6O13 because of incomplete reversibility. Additionally, other phases of Li0.6V1.67O3.67 and Li3VO4 exhibit regular changes with Li + insertion/extraction, and the peak intensity increases with ion insertion and decreases with ion extraction. Figure 11 shows the Raman spectra after 100 cycles. The Raman peaks of the N−400 and N−450 are basically same, but there is an obvious difference with the film before cycling. A red shift of these Raman peaks both in N−400 and N−450 is observed after 100 cycles. The shift is attributed to the Li + between the V-O layers, which results in a lattice expansion, and the split peak at 139 cm −1 may also related to it [30]. In order to obtain valence information for V, O, and Li of the films, XPS was carried out on the N−400 film before cycling and after 100 CV cycles. Figure 12a reveals that the N−400 film at different states both possess peaks of V and O, proving their coexistence directly. The obvious peak appearing at about 450 eV is the In of ITO substrate, which is due to the etching. Figure 12b shows the O 1s spectra of the N−400 film in different states. The O 1s spectra shift to the right after cycling, which can be attributed to the following two reasons: (1) The Li + between the V-O layers bonding with O; (2) changes in the oxygen environment caused by etching onto the ITO glass substrate. Figure 12c shows that Li + exist in both charge/discharge states, and the Li 1s is located at about 56 eV, which is the same as in other literatures [42,43]. The valence changes of V are shown in Figure 12d-f. It is obviously that the ratio of V 5+ /V 4+ changes in different states. The N−400 film before cycling delivers a V 4+ /V 5+ ratio of 1.9, which is basically consistent with V 4+ 4V 5+ 2O13, and is corresponding to the average valence of V (+4.33) in V6O13. After the hundredth discharge process, the V 4+ /V 5+ ratio turns down to 1.3, because of the formation of the new Li3VO8 phase during cycling, which leads to the increase of the valence state of V. After the hundredth charge process, the V 4+ /V 5+ ratio turns back to 1.4, which is caused by a part of the new Li3VO8 phase reversing back to V6O13. The results of the ex-situ XPS further prove the findings of the ex-situ XRD results and phase transitions during cycling. In order to obtain valence information for V, O, and Li of the films, XPS was carried out on the N−400 film before cycling and after 100 CV cycles. Figure 12a reveals that the N−400 film at different states both possess peaks of V and O, proving their coexistence directly. The obvious peak appearing at about 450 eV is the In of ITO substrate, which is due to the etching. Figure 12b shows the O 1s spectra of the N−400 film in different states. The O 1s spectra shift to the right after cycling, which can be attributed to the following two reasons: (1) The Li + between the V-O layers bonding with O; (2) changes in the oxygen environment caused by etching onto the ITO glass substrate. Figure 12c shows that Li + exist in both charge/discharge states, and the Li 1s is located at about 56 eV, which is the same as in other literatures [42,43]. The valence changes of V are shown in Figure 12d-f. It is obviously that the ratio of V 5+ /V 4+ changes in different states. The N−400 film before cycling delivers a V 4+ /V 5+ ratio of 1.9, which is basically consistent with V 4+ 4 V 5+ 2 O 13 , and is corresponding to the average valence of V (+4.33) in V 6 O 13 . After the hundredth discharge process, the V 4+ /V 5+ ratio turns down to 1.3, because of the formation of the new Li 3 VO 8 phase during cycling, which leads to the increase of the valence state of V. After the hundredth charge process, the V 4+ /V 5+ ratio turns back to 1.4, which is caused by a part of the new Li 3 VO 8 phase reversing back to V 6 O 13 . The results of the ex-situ XPS further prove the findings of the ex-situ XRD results and phase transitions during cycling.

Conclusions
In summary, nanosheet-like (NH4)8(V19O41(OH)9)(H2O)11 precursor films were synthesized via a simple low temperature liquid phase deposition method, and the V6O13 nanosheets film was successfully obtained by the following heat treatment under N2 atmosphere. The V6O13 films were directly used as cathodes for LIBs without adding binders and conductive agents. The results showed that the film annealed at 400 °C had the best electrochemical performance, with diffusion coefficients of 6.084 × 10 −12 cm 2 s −1 (Li + insertion) and 5.464 × 10 −12 cm 2 s −1 (Li + extraction), and with excellent cycling performance during CV cycles, which remained at 82.7% of capacity after 100 cycles. The ex-situ XRD results revealed the mechanism. There existed phase transitions with the insertion/extraction of Li + . The V6O13 partly transformed to Li0.6V1.67O3.67, Li3VO4, and VO2 with the insertion of Li + into the lattice, and Li0.6V1.67O3.67, Li3VO4, and VO2 partly reversibly transformed backwards to V6O13 with the extraction of Li + . The phase transition was ascribed to the structure and nanosheet morphology of V6O13, which provided more ion diffusion paths and sufficient active sites during cycling. Moreover, there were enough space and diffusion channels in the structure for the phase transition. This work may provide an Figure 12. XPS spectra of the N−400 film: XPS spectrum before and after 100 cycles (a); O 1s before and after 100 cycles (b); Li 1s before and after 100 cycles (c); V 2p before cycling (d); after 100th discharge state (e); after 100th charge state (f).

Conclusions
In summary, nanosheet-like (NH 4 ) 8 (V 19 O 41 (OH) 9 )(H 2 O) 11 precursor films were synthesized via a simple low temperature liquid phase deposition method, and the V 6 O 13 nanosheets film was successfully obtained by the following heat treatment under N 2 atmosphere. The V 6 O 13 films were directly used as cathodes for LIBs without adding binders and conductive agents. The results showed that the film annealed at 400 • C had the best electrochemical performance, with diffusion coefficients of 6.084 × 10 −12 cm 2 s −1 (Li + insertion) and 5.464 × 10 −12 cm 2 s −1 (Li + extraction), and with excellent cycling performance during CV cycles, which remained at 82.7% of capacity after 100 cycles. The ex-situ XRD results revealed the mechanism. There existed phase transitions with the insertion/extraction of Li + . The V 6 O 13 partly transformed to Li 0.6 V 1.67 O 3.67 , Li 3 VO 4 , and VO 2 with the insertion of Li + into the lattice, and Li 0.6 V 1.67 O 3.67 , Li 3 VO 4 , and VO 2 partly reversibly transformed backwards to V 6 O 13 with the extraction of Li + . The phase transition was ascribed to the structure and nanosheet morphology of V 6 O 13 , which provided more ion diffusion paths and sufficient active sites during cycling. Moreover, there were enough space and diffusion channels in the structure for the phase transition. This work may provide an inspiration for enhancing the performance and studying the process mechanisms of vanadium oxides as cathode materials for LIBs.

Conflicts of Interest:
The authors declare no conflict of interest.