Structural, optical, thermal and conducting properties of V2−xLixO5−δ (0.15 ≤ x ≤ 0.30) systems

Lithium-doped vanadates (V2−xLixO5−δ (0.15 ≤ x ≤ 0.30)) are synthesized by melt-quench method. The physical, structural, optical, thermal and conducting properties of as-quenched samples are investigated using various experimental techniques to study their suitability for electrolyte in battery/solid oxide fuel cell application. X-ray diffraction (XRD) patterns confirm the formation of three different crystalline phases. FTIR and Raman spectra indicate that the doping of Li2O into V2O5 leads to a transition from VO5 into VO4 structural unit. The optical diffused reflectance spectra revealed that the optical band gap (Eg) decreases from 2.2 to 2.08 eV while Urbach energy (EU) increases (0.31–0.41 eV) with the addition of Li2O content in place of vanadium. The thermal stability is studied by thermogravimetric analyser (TGA). The DC conductivity of the present samples is increased from 0.08 to 0.12 Scm−1 at 450 °C with Li2O doping. These materials can be used as electrolyte for battery/solid oxide fuel cell due to their good conductivity (~0.12 Scm−1) at 450 °C.

Solid oxide fuel cells (SOFCs) are most efficient electrochemical devices which convert chemical energy into electricity with heat and water as a by-product. SOFCs are considered as eco-friendly since there is no production of sulfur or nitrogen oxides (only CO 2 ) along with fuel flexibility (H 2 , CO, hydrocarbons and biogases) [1][2][3] . Since, the high operating temperature of SOFC leads to degradation, coefficient of thermal expansion mismatch, electrode sintering and catalyst poisoning among the components of SOFCs 1-3 . Thus, a lot of research is carried out to develop new materials that can meet the requirements at lower operating temperature of SOFCs [3][4][5] . Doped V 2 O 5 materials can be a good choice for low temperature (600-800 °C) SOFCs applications due to their good conductivity at lower temperature 6 . The electrical conduction in V 2 O 5 based materials arises by exchanging of unpaired 3d 1 electrons (hopping) between V 4+ and V 5+ valence states, V 4+ -O-V 5+ -O-V 4+ . These unpaired electrons are induced polarization around vanadium ions which leads to form polaron 6,7 . Alkali metal oxide doped vanadates have been studied widely due to their numerous applications in optical and electronic devices 1,2,8 . The chemical nature and concentration of dopants play a vital role to change the oxidation state of vanadium which leads to a creation of oxygen vacancies to maintain overall electrical neutrality of the vanadate systems 5,9,10 . Punia et al. have reported the increasing trend of density and molar volume due to change in structure from VO 4 tetrahedral to VO 5 trigonal bi-pyramid with the addition of Bi 2 O 3 in zinc vanadates 11 . In addition to this, alkali ions doped materials usually show mixed electronic and ionic conductivity behavior with temperature. Doped bismuth vanadates (Bi 4 V 2 O 11−δ ) show good conductivity ~0.2 Scm −1 at 500 °C due to the higher oxygen vacancies and stabilization of high conducting γ-phase at room temperature 12 . The highest conductivity i.e. 6.6 × 10 −4 Scm −1 for Bi 4 V 2−x Al x O 11−δ (x = 0.2) system at 500 °C has been reported by Ravikant et al. 4 . Chakrabarty et al. 13 studied the doping effect of Na 2 O on conducting properties of Na 2 O-V 2 O 5 system and found an increment in conductivity (0.05 × 10 −3 -19.5 × 10 −3 Scm −1 ) with Na 2 O concentration at room temperature. A study of V 2−x Mg x O 5−δ (x = 0.05-0.30) systems has been revealed a decreasing trend in density (3.2-2.7 gcm −3 ) and thermal stability with MgO concentration. The structural transition has occurred from VO 5 to VO 4 polyhedra with MgO dopant. DC conductivity decreases from 10 −1 Sm −1 to 10 −4 Sm −1 while activation energy increases from 0.27 to 0.44 eV with an increase in MgO content in place of vanadium 6 . The optical bandgap increases whereas Urbach energy decreases with the increasing content of Bi 2 O 3 due to decreasing oxygen vacancies 14 .
Electrical properties of MnO 2 doped V 2 O 5 system are also reported 15 . Tsuzuki et al. 16 have investigated V 2 O 5 -MO (M = Mg, Ca and Ba) systems and correlate the structural changes with electrical properties. In most of the cases, these doped systems are prepared by solid-state reaction/chemical method followed by slow cooling which leads to lower the conductivity due to ordering in oxygen vacancies 5 . It has been reported by some research groups that quenched samples have higher disordering/defects which leads to increase in the overall electrical conductivity. The high quenching rate encourages the reduction of vanadium which leads to higher disordering and hence may increase the overall conductivity 5,17 . Based on above discussion, it could be concluded that processing parameters and chemical nature of dopants affect the properties significantly particularly conductivity. Therefore, the motivation of the present work is to study the effect of systematic change of lower valance dopant (Li 2 O) concentration on optical, thermal and conducting properties of V 2 O 5 grown by melt-quench technique. The V 2−x Li x O 5−δ (x = 0.15-0.30) systems have been synthesized by melt-quench followed by various characterization techniques to study their structural, optical, thermal and electrical properties to check their suitability as electrolytes for SOFCs applications.

Results and Discussion
Physical properties. Various physical parameters such as density, molar volume and ionic concentration and their inter-ionic distance and polaron radius are calculated for the as-quenched samples and depicted in Table 1. Figure 1 shows the variation of density (ρ) and molar volume (V m ) with Li 2 O dopant of as-quenched samples. It is cleared (Fig. 1) that the density decreases while molar volume increases with increasing Li 2 O content. This decrement in the density is due to the replacement of the heavier V 2 O 5 (3.36 gcm −3 ) by the lighter Li 2 O (2.01 gcm −3 ) 6,18 . On the other hand, the doping of Li 1+ into V 5+ also creates oxygen vacancies. The size of oxygen vacancy is smaller as compared to the size of the oxygen anion O 2− . So, the volume of the samples with dopant also decreases which leads to increase in the density. However, a lower density of dopant Li 2 O as compared to parent V 2 O 5 influences the density of samples more efficiently than decrease the volume due to creation of vacancies in the samples 16 . Hence, the density of samples decreases with the addition of dopant (Li 2 O) in place of vanadium. On the other hand, the increase in molar volume is ascribed to the rearrangement of the lattice which leads to creating oxygen vacancies due to VO 5 → VO 4 conversion with the substitution of Li 1+ for V 5+ (as discussed in FTIR and TGA section) 18 . This increase in molar volume supports that the Li 2 O enters the network and occupy the interstitial space, which leads to an increase in free volume. Figure 2 shows the change in ionic concentration (Ni) and their inter-ionic distance (R i ) with Li 2 O concentration on the cost of vanadium for all the samples. The ionic concentration increases whereas inter-ionic distance decreases (1.8-1.4 Ǻ) with increasing Li 2 O content. Decrement in the inter-ionic distance might be increased the localization effect as also confirmed by the polarization radius values (Table 1). In general, the decrease in molar volume is responsible for the decrease in inter-ionic   Structural properties. X-ray diffraction (XRD).  Fig. 3. The shifting of peak towards higher diffraction angle side is the evidence of lattice contractions 19 . This is due to the smaller cationic radius Li 1+ (0.167 nm) in comparison to V 5+ (0.171 nm). As a result of the substitution of the larger cationic radius of V by the smaller cationic radius of Li, the lattice compression (compressive strain) arises in the Li 0.04 V 2 O 5 crystalline phase 19 . Further comparing the peak intensities of the phases, it is observed that the intensity of the major phase   Table 2. So, the volume fraction of less symmetric phases (monoclinic) increases as dopant concentration increases.
FTIR spectroscopy analysis. Figure 4 shows infrared (IR) spectra at room temperature to investigate the structural changes with increasing concentration of Li 2 O dopant. The IR spectra exhibited prominent absorption bands at 1022, 1000, 957, 916, 819, 753, 602, 488 cm −1 . The absorption bands at 1022 and 1000 cm −1 are ascribed to the vibrations of isolated V=O vanadyl groups in VO 5 trigonal bipyramids [20][21][22] . This characteristic stretching band with Li 2 O doping present samples, which is associated with the layer structure of the V 2 O 5 [20][21][22][23] . The broadness of this band decreases with increasing Li 2 O content due to the decrease in the number of ionic groups and V=O bonds. Additionally, this band has a small shift to the lower wavenumber, which suggests the reduction of vanadium (V 5+ to V 4+ /V 3+ ) due to the addition of lithium content 24 . The V-O bond length increases due to the larger size of V 4+ (0.67 Ǻ) than V 5+ (0.60 Ǻ), results in requiring lower energy for the lattice vibration 25 . The high-temperature melt quenching of the present samples could be responsible for the reduction in the oxidation states of vanadium due to scarcity of oxygen in the furnace at high temperature 24 . The band at 957 cm −1 assigned to the symmetric and anti-symmetric stretching vibration of the isolated VO 2 groups in the VO 4 polyhedra 20,26 . A small kink at 916 cm −1 has appeared with the addition of Li 2 O content. It is also assigned to the asymmetric stretching vibrations of the isolated VO 2 groups in VO 4 polyhedra 20,24 . Another band at about 819 cm −1 is attributed to the vibrations of isolated [VO 4 ] tetrahedral. The broadness of this band decreases with Li 2 O content and splits into two bands (819 and 753 cm −1 ) at maximum Li 2 O content sample (VL-0.30). This new generated band is assigned to V-O-V stretching 6,20,26 . The broadening of the bands is directly related to the existence of different structural units of the same element with variable oxygens 6 20,23,27 . This band shows a small shift to the higher wavenumber with the highest doping of Li 2 O (VL-0.30 sample), which shows that some portion of V 3+ / V 4+ oxidized to V 5+ . The bond length between vanadium and oxygen decreases due to the smaller size of V 5+ as compared to V 4+ /V 3+ , which leads to strengthening of this vibration 6,25,28 . According to Dimitriv et al. 26 Table 2. Optical and conductivity parameters of the as quenched samples. www.nature.com/scientificreports www.nature.com/scientificreports/ vanadates can exist in the form of layers and chains and smaller complexes, which depends on the ratio of V 2 O 5 and Li 2 O in the samples. The Li 1+ ions can be situated between the vanadate chains, whereas some Li ions can be form independent polyhedra (LiO n ), which can be located in the chains themselves 26 . This assumption suggests two different forms of Li ions distribution in the V-O polyhedra (interstitial sites or substitutional sites) as shown in Fig. 5(a,b). In interstitial condition, the Li ions can be located between vanadate chains and layers. Li ions can interact with the isolated V=O bonds, which leads to the longer bond length (Fig. 5(a)). In substitutional condition, Li ions can occupy positions in the vanadate chain, which leads to break the some of the weaker V-O-V bonds and form new V-O-Li bridges (Fig. 5(b)) 26 . Also, there is a possibility to take place a redox reaction ( 2 ) during the melting of the batches at 900 °C. The V 5+ ions take a part as a network forming positions with VO 5 structural units which can be formed V-O-Li linkages 24 . Therefore, the influence of Li ions on the V=O bond is restricted and have an indirect manifestation which leads to preserve this band at ~1022 cm −1 26 . Thus, the V=O band is partially affected by Li 2 O doping in the present samples.
Raman spectroscopy analysis. Raman spectra of as-quenched samples (Fig. 6) 32 . This decrement in reflectivity with increasing Li 2 O dopant into V 2 O 5 lattice can be due to the reflectance onset/absorption edge towards longer wavelengths (red shift). It is good agreement with other reports 32,33 . This reflectance data is used to calculate the optical band gap by Kubelka-Munk function 34 where, F(R), K and S are the Kubelka-Munk function, absorption coefficient and scattering coefficient, respectively. The optical band gap is obtained from DRS using following equation.
Where, k, E g and ν h are the proportionality constant, optical band gap and incident photon energy, respectively. The exponent 'n' depends on the transition, where n = 1/2, 2, 3/2 and 3 for allowed direct, allowed indirect, forbidden direct and forbidden indirect transition, respectively 34 . The direct optical band gap of all the samples is obtained using the Tauc plot ( Fig. 7(b)) by extrapolating the straight-line portion of the plots of Thermal properties. Figure 8(a) represents the DTA curves of lithium-doped V 2 O 5 as-quenched samples in the temperature range of 30-800 °C. In DTA curves, all the samples exhibit three distinct endothermic peaks at a higher temperature (≥571 °C), which result from the melting points (Tm 1 , Tm 2 and Tm 3 ) of three different phases as observed in XRD patterns. Figure 8(b) shows the TG curves of the Li 2 O doped V 2 O 5 as-quenched samples, which exhibit typical weight loss behavior. TGA thermograph shows a total weight loss of ~5% in the temperature range of 30-800 °C. The weight loss can be considered in three steps as marked I (30-120 °C), II (120-600 °C) and III (600-700 °C) in the TGA graph. The I step is assigned to the departure of adsorbed species and loss of water molecules 34 . The II step is associated to the reduction of V 5+ to V 4+ /V 3+ ions which can be shown by the following equation 24,34 . www.nature.com/scientificreports www.nature.com/scientificreports/ In the III step, a little weight gain is observed in all the samples which can be related to the oxidation of V 3+ to V 5+ as written by the following equation 37,38 .
This weight gain during the oxidation at the higher temperature can be associated with the combination of O 2 and its vacancies in the present samples. The reduction of these systems is expected due to vanadium is believed to have a mixed valence state (V +3 , V +4 and V 5+ ) 24,37,38 . It is also observed that weight loss increases with the addition of Li 2 O on the cost of V 2 O 5 which leads to a decrease in the thermal stability of the present samples (I and II steps). It is manifestation of waterphilic nature of the samples and it increases with an increase in Li 2 O concentration. The reduction of vanadium from V 5+ state to lower valance state is also increases with the addition of Li 2 O content. The chemical stability also decreases with Li 2 O concentration which is also manifested by Urbach energy and FTIR analysis of the samples. All the samples exhibit good thermal and chemical stability in the temperature range of ~300-550 °C without any phase transition ( Fig. 8(a,b)) except VL-0.30 sample. However, sample with the lowest content of Li 2 O (7.5 mol%, VL0-0.15) shows the highest thermal and chemical stability in the temperature range of ~300-550 °C in comparison to other samples. Therefore, these samples can be used as electrolytes for solid oxide fuel cells or batteries applications.
Conducting studies. It is a well-known fact that hopping mechanism and band conduction are responsible for the electrical conduction in these types of materials 6,39 . The frequency dependent conductivity, σ ac of the samples is determined using the following relation 6,39 .
where, ε 0 ε′, ε″, f and tanδ are the permittivity of free space (8.854 × 10 −12 Fm −1 ), dielectric constant, dielectric loss, frequency (in Hz) of the applied electric field and tangent of loss, respectively. Figure 9(a-d) illustrates the AC conductivity (σ ac ) variation with frequency and temperatures (50-450 °C, step size of 50 °C). The conductivity is almost frequency independent at low frequency region (plateau region) and corresponds to the DC conductivity (σ dc ) of the samples 40 . At the same time, the conductivity is dependent on frequency (σ ac ) at higher frequency region (dispersion region) and increases with increasing frequency 6,15,39,40 . The plateau region signifies the randomness of ionic diffusion which is due to the long-range order transport and appreciable diffusion of mobile ions in response to the applied ac field while the dispersion region signifies the conductivity relaxation phenomena in present samples 6,7,15,40,41 . It is also observed that σ ac increases with increasing temperature (50-450 °C) for all the samples which supports the semiconductor nature of the samples 6,42 . Generally, the conductivity in such type of materials depends on the concentration and mobility of mobile ions/charge carriers i.e. electrons and holes 6,42 . The mobility of mobile ions/charge carriers increases with increasing temperature which leads to increasing overall conductivity as observed in the present samples. It can be understood by the fact that thermal energy (temperature) excites more electrons from valence band to conduction band resulting in higher conductivity. Further, according to the Jonscher's universal power law, σ ac can be written as 6,41,43,44 : ac dc s where, A is the temperature dependent constant of the material, ω is the angular frequency and s is the frequency and temperature dependent exponent that lies between 0-1. DC conductivity (σ dc (T)) is calculated to extrapolate AC conductivity at the lowest frequency ( Fig. 9) and given in Table 2. The activation energy (E a ) of the conduction is calculated using the Arrhenius plot (Fig. 10) between lnσ dc and reciprocal of temperature (1000/T) by the following equation. www.nature.com/scientificreports www.nature.com/scientificreports/ where, σ o and k B are the pre-exponential factor, and Boltzmann's constant, respectively. The values of σ ac and σ dc show an increasing trend with an increase in lithium content (7.5-15 mol%) as shown in Fig. 11. AC conductivity lies in the range of 0.11-0.29 Scm −1 at 450 °C and 10 7 Hz whereas DC conductivity lies in the range of 0.08-0.12 Scm −1 at 450 °C. In vanadate materials, conductivity depends on the small polarons hopping (SPH) between V 4+ site to the neighboring V 5+ site 6,7,39 . SPH directly affects by the mobility and number of mobile charge carriers. On the replacement of V 5+ by Li + , structure of the systems becomes more open due to transition from VO 4 tetrahedral to VO 5 trigonal bipyramid as discussed in physical parameters and structural properties. This structural change is responsible for a significant increase in oxygen vacancies with an increase in hopping of ions which leads to an increase in conductivity 4,5,42 . In addition, decrement in density, inter-ionic distance, polaron radius and optical band gap of the samples also support this increasing trend of conductivity with the addition of Li 2 O  It is observed that the transport mechanism in lithium vanadate systems is attributed to small polaron hopping (SPH) between V 4+ to V 5+ states, which is increased by the addition of lithium into vanadium. The present samples exhibit good DC and AC conductivity. Therefore, these materials can be considered as electrolytes for batteries/solid oxide fuel cells applications. . The polaron radius was calculated by π = × r R ( /2) ( /6)

Materials and Methods
. The X-ray diffraction (XRD) patterns of crushed samples were recorded by PANalytical X'Perts Pro MPD diffractometer with Cu-Kα radiations (λ = 1.54 Å). The scan rate and scan range were 2°min −1 and 10-80°, respectively. Fourier transform infra-red (FTIR) spectra of as-quenched samples were recorded in the wavenumber range of 200-4000 cm −1 at room temperature using Perkin Elmer-Spectrum-RF-1 FTIR spectrometer. The finely ground powder of as-quenched sample and KBr powder are mixed together to palletized for the FTIR measurement. The Raman spectra of the powder samples were recorded over the spectral range of 100-5000 cm −1 using Features STR 500 Airix Raman system equipped with Ar laser (532 nm) as the excitation source.
A double beam UV-Vis spectrophotometer (Model: Hitachi 3900H) was used to record diffused reflectance spectra (DRS) at room temperature between the wavelength 200-800 nm. Thermal analysis was carried out using TG/DTA (Exstar TG/DTA 6300) instrument of as-quenched powder samples in N 2 medium with a heating rate of 10 °C/min over the range of room temperature to 1000 °C. The dielectric and conductivity analysis carried out using SI-1260 Solartron analytical LCR impedance analyzer. Both sides of the as-quenched samples are coated with Pt using 3000-FC auto fine coater from JEOL to serve as an electrode. The variation in impedance with frequency and temperature were studied from 10Hz-1MHz and 50-400 °C, respectively.