# Structural, vibrational and electronic properties of Nb substituted
orthovanadates LaV1-xNbxO4

Ashok Kumar  Department of Applied Physics, Delhi Technological University,
Delhi-110042, India  Department of Physics, Atma Ram Sanatan Dharma College,
University of Delhi, New Delhi-110021, India  Anurag Sharma  Department of
Physics, Indian Institute of Technology Delhi, Hauz Khas, New Delhi-110016,
India  Madhav Sharma  Department of Physics, Indian Institute of Technology
Delhi, Hauz Khas, New Delhi-110016, India  Vinod Singh  Department of Applied
Physics, Delhi Technological University, Delhi-110042, India  Anita Dhaka
Department of Physics, Hindu College, University of Delhi, New Delhi-110007,
India  Rajendra S. Dhaka<EMAIL_ADDRESS>Department of Physics,
Indian Institute of Technology Delhi, Hauz Khas, New Delhi-110016, India

###### Abstract

We investigate the structural, vibrational, morphological, and electronic
properties of Nb substituted orthovanadate LaV1-xNbxO4 samples prepared by the
solid-state reaction method. The x-ray diffraction (XRD) analysis reveals the
presence of three crystal structures [monoclinic monazite ($m-m$) type for the
$x=$ 0, two-phase equilibrium of monoclinic monazite ($m-m$) and tetragonal
scheelite ($t-s$) type for the 0.2$\leq$$x$$\leq$0.8, and monoclinic
fergusonite ($m-f$) type for the $x=$ 1 samples] with an increase in Nb5+
concentration. The Raman spectroscopy and x-ray photoelectron spectroscopy
(XPS) were employed to study the vibrational and electronic properties of all
the samples, respectively. In order to choose an excitation wavelength that
does not cause undesirable fluorescence and has observable intensities of all
the vibrational modes, the Raman spectra are collected using 532 nm, 633 nm,
and 785 nm laser lines. With increasing the Nb5+ concentration, new Raman
modes associated with Nb-bonds are clearly visible and the intensity of
V-bonds assigned modes is decreasing. The XPS analysis shows the unchanged 3+
oxidation state of La ion where the intensity of the V 2$p$ core-level
decreases while the Nb 3$d$ core-level increases with $x$. The equal spin-
orbit energy splitting of the states is confirmed by the average energy
difference (across La core-level spectra for all the samples) for state I as
well as bonding and anti-bonding of state II. Interesting, the relative
intensity of La 3$d$ state I and state II show systematic change with Nb
doping altering the metal ligand overlap. We discuss and provide insight into
the evolution of the structural, morphological, and chemical features with Nb
substitution in LaV1-xNbxO4 samples.

##  I  Introduction

In various polycrystalline oxides, rare earth orthovanadates (RVO4; R-Rare
earth elements) are interesting because of their potential applications in
catalysis, polarizers, luminescent materials, and laser host materials [1, 2,
3]. Also, researchers have reported that complex oxide materials show
interesting structural, magnetic and electronic properties [4, 5, 6, 7], and
may be utilized for various applications such solid oxide fuel cells and as an
electrode material for Lithium-ion batteries due of their high specific
capacity and cycle stability [8]. It is interesting to note that the lanthanum
based orthovanadates LaVO4 shows the structural trend in rare-earth family, it
crystallizes in tetragonal–zircon ($t-z$) type polymorphs with space group
I41/amd and monoclinic–monazite ($m-m$) type polymorph with space group P21/n.
However, it thermally stabilizes in $m-m$ type, whereas the $t-z$ structure
remains in metastable state at room temperature, because of the largest ionic
radius of La3+ in lanthanide series, it has a higher oxygen coordination
number (9) in $m-m$ type structure as compared to 8 in $t-z$ type [9]. The
zircon structure contains a pattern of VO4 tetrahedra (having four identical
V-O bonds) [10] and RO8 dodecahedra (coordination no. 8), sharing their edges
alternatively and linked together in chains along the $c-$axis. In the
monazite structure, deformed VO4 tetrahedra with four different V-O bonds [11]
are connected to RO9 polyhedra (coordination no. 9) and sharing their edges.
The zircon type LaVO4 sample is difficult to prepare at ambient conditions by
conventional solid state reaction method but few reports say that it can be
synthesized and stabilized by hydrothermal and precipitation methods [12, 13,
14].

The structural and electronic properties of lanthanum orthovanadate with
pentavalent niobium substitution are vital to understand for their practical
use. Though the parent compound LaVO4 with substitution at the La site has
been extensively explored [15, 16], there are very few studies to understand
the effect of substitution at V site [18, 17]. As the niobium is located just
below vanadium in the periodic table and has many advantages like Vanadium
prices have recently risen to about 300% higher, niobium (Nb5+) is
biocompatible, isoelectronic to vanadium ion and has a larger ionic radius
(0.48 Å) with four coordination numbers in comparison to vanadium ion (0.36 Å)
[19]. The LaNbO4 is a rare-earth niobate and shows a well-known temperature
and composition/substitution-induced structural transformation. For example,
the LaNbO4 undergoes a thermally induced structural transition from monoclinic
fergusonite ($m-f$) with space group I2/a to tetragonal scheelite ($t-s$) with
space group I41/a) phase at $\sim$495$\degree$C [20]. Similarly, it undergoes
structural transformation by substituting Nb5+ at V5+ site [21]. It has been
reported that lanthanum niobate shows interesting properties and very useful
for technological applications such as proton conductivity [22, 23], good
dielectric, high energy emission using X-ray excitation [24] and its potential
for applications in a variety of fields, including sensors [25], contrast
agents, waveguides, ferroelectrics [26], phosphors [27], laser crystals [28],
luminophores, LEDs [29], etc.

In this paper, we study the structural, vibrational, morphological, and
electronic properties of LaV1-xNbxO4 using various experimental tools like
x-ray powder diffraction (XRD), scanning electron microscopy (SEM), high
resolution transmission electron microscopy (HR-TEM), selected area electron
diffraction (SAED), Raman spectroscopy, and x-ray photoelectron spectroscopy
(XPS). We find the phase purity and structural transition by performing the
Rietveld refinement of the measured XRD patterns at room temperature. The
Raman spectra of LaV1-xNbxO4 samples are measured with different excitation
wavelengths of 532 nm, 633 nm, and 785 nm, where we find significant intensity
of all the Raman active modes as well as interesting changes with Nb
substitution. The Raman spectra exhibit a pattern of maximum intensity peaks
that is compatible with Badger’s rule. The structural phase transition
observed in the XRD analysis of LaV1-xNbxO4 is also supported by the intensity
variation in the Raman mode observed in the samples with increasing Nb
concentration. Through the SEM micrographs, we identify that the samples
contain fine particles along with pores as well as changes in particle size
and shape can be seen in the surface images of the samples. The core-level
photoemission reveals the oxidation state and electronic structure of the
constitute elements in these samples. The intensity of the core-level spectra
of all the samples varied systematically with an increase in Nb5+
concentration, as shown by XPS analysis. The average energy difference (for
the La core-level spectra of all the samples) for state I, state II bonding,
and state II anti-bonding verified the equal spin-orbit energy splitting of
the states. Moreover, we find a systematic change in the relative intensity of
La 3$d$ state I and state II with Nb doping, which suggest an altering in the
metal ligand overlap.

##  II  Experimental

We use solid-state reaction method to prepare LaV1-xNbxO4 ($x=$ 0 to 1)
samples by mixing V2O5 (99.6$\%$, Sigma), Nb2O5 (99.99$\%$, Sigma), and La2O3
(99.99$\%$, Sigma) as precursors in the stoichiometric proportions. The La2O3
was pre-dried for 6 hrs at 900$\degree$C to remove the moisture. After that
the mixture was ground evenly for 8 hours, then heated for 17 hrs at
1000$\degree$C. The mixture was then reground and sintered at 1250$\degree$C
for 13 hrs to improve the crystallinity of the samples. The phase purity and
structural parameters of LaV1-xNbxO4 were determined using Panalytical XPert3
powder x-ray diffractometer at room temperature using the Cu source of
K$\alpha$ radiation ($\lambda$ = 1.5406 Å ). We use the step size of
0.033$\degree$ for each XRD scan taken in the 2$\theta$ range from 10$\degree$
to 90$\degree$. The lattice parameters are extracted by the Rietveld
refinement of XRD patterns using FullProf software, where linear interpolation
is used to fit the background. We use Jeol JSM-7800F Prime field emission
scanning electron microscope (FE-SEM) with LN2 free SDD X-max 80 EDS detector
in high vacuum mode to produce the scanning electron microscope (SEM)
micrographs of the materials’ surfaces. The analysis of particle size and
change in morphology of LaV1-xNbxO4 was done using ImageJ software by
analyzing SEM micrographs at the surface of the pellet samples. In order to
execute FE-SEM, the non-conducting LaV1-xNbxO4 pellets were turned into
conducting by coating the surface with a thin layer of Au using a sputter
coater. We use the JEOL/JEM-F200 microscope, equipped with thermal electron
field emission and OneView CMOS camera (4k $\times$ 4k pixels), to collect HR-
TEM data by operating the system at an acceleration voltage of 200 keV.

The Raman spectra were recorded at room temperature with the Renishaw inVia
confocal Raman microscope using 2400 lines/mm grating, 10X objective, and
three different wavelengths; (i) 532 nm, gas laser with a power of 1 mW, (ii)
633 nm, where the semiconductor diode laser with a power of 1 mW, (iii) 785 nm
semiconductor diode laser with a power of 0.1 mW. The samples can be
identified by their particular Raman fingerprint, and their structural and
chemical information can be discovered through the examination of several
Raman active modes in LaV1-xNbxO4. The x-ray photo emission spectroscopy (XPS)
measurements are done using AXIS Supra instrument (Kratos Analytical Ltd). The
survey spectra and core level spectra (La 3$d$, Nb 3$d$, V 2$p$, and O 1$s$,
for each sample), were recorded at room temperature using the monochromatic
X-ray source: Al K$\alpha$-1486.6 eV(step size 1 eV for the survey and 0.1 eV
for core level spectra), with a charge neutralizer, is used to offset the
charging impact in these insulating materials. The pass energy of the analyzer
was 160 eV and 20 eV for the survey and core-level spectra, respectively. For
all the wide scans and core-level spectra, the C 1$s$ peak is fitted to obtain
the peak binding energy (BE) and the calibration for charge correction was
done using the C 1$s$ BE reference at 284.6 eV for each sample. We utilize the
Igor Pro 9 software to analyze the observed Raman spectra and fitted the modes
using the Lorentzian peak function as well as XPS spectra using the Voigt
function.

##  III  Results and Discussion

Figure 1: (a–f) The Rietveld refined x-ray diffraction patterns of LaV1-xNbxO4
($x=$ 0–1) samples. The experimental, simulated, and difference between the
experimental and simulated spectra are shown by open red circles, black solid
lines, and blue solid lines, respectively. The Bragg positions corresponding
to their respective space groups are shown by green vertical markers. In the
side of each panel (a1–f1), we show partial amplification for clarity between
2$\theta$ = 25–35$\degree$ for all the samples.

The Rietveld refined room-temperature x-ray diffraction (XRD) patterns of the
polycrystalline LaV1-xNbxO4 ($x=$ 0–1) samples are displayed in Fig. 1 and
lattice parameters of the samples are summarised in Table 1, where we can see
that angle $\beta$ is increasing in the $m-m$ type phase of LaV1-xNbxO4
samples with Nb5+ substitution due to higher ionic size of Nb5+ as compared to
V5+. The crystallization of LaV1-xNbxO4 is clearly observed in three different
phases depending on the substitution of Nb5+ at the site of V5+, and also been
reported by Aldred et al. in ref. [21]. We observe that the structure changes
from $m-m$ to $m-f$ with increase in the Nb5+ concentration from 0 to 100%.
For the $x=$ 0 and 1, a pure monoclinic phase is obtained with no impurity
peaks. In between $x=$ 0.2 and 0.8, a monoclinic monazite ($m-m$) and a
tetragonal scheelite ($t-s$) type phases coexist. Moreover, all the Bragg
reflections of LaVO4 and LaNbO4 can easily be indexed to the $m-m$ and $m-f$
phases with the space group P21/n and I2/a for the $x=$ 0 and 1 samples,
respectively. We find that the contribution of space group I41/a is increasing
from the $x=$ 0.2 to 0.8 samples (see Table 1) due to the increase of $t-s$
phase with the substitution of Nb5+ at the site of V5+ in LaV1-xNbxO4 samples.
So, it can clearly be seen that the LaV1-xNbxO4 samples crystallize in
monoclinic monazite ($m-m$) type ($x=$ 0), coexistence of monoclinic monazite
($m-m$) and tetragonal scheelite ($t-s$) type (0.2$\leq$$x$$\leq$0.8), and
monoclinic fergusonite ($m-f$) type ($x=$ 1) [21, 30].

Table 1: The Rietveld refinement parameters of polycrystalline LaV1-xNbxO4 ($x=$ 0–1) samples with Nb substitution induced metastable tetragonal–scheelite phase for the $x=$ 0.2 to 0.8 samples, determined using the FullProf software. $~{}~{}~{}x~{}~{}~{}$ |  $\chi^{2}$ |  Space Group |  a (Å) |  b (Å) |  c (Å) | $\beta(\degree)$ |  Volume $(\rm\AA^{3})$
---|---|---|---|---|---|---|---
0 | 1.09 | P21/n |  7.042(3) |  7.276(4) | 6.724(7) |  104.88 (6) | 333.033(5)
0.2 | 2.63 | P21/n - 84$\%$ | 7.046(1) | 7.278(2) | 6.733(3) | 104.91(1) | 333.685(4)
|  | I41/a - 16$\%$ | 5.336(1) | 5.336(1) | 11.731(2) | 90 | 334.042(4)
0.4 | 2.46 | P21/n - 78$\%$ | 7.043(0) | 7.276(3) | 6.732(4) | 104.91(2) | 333.397(3)
|  | I41/a - 22$\%$ | 5.332(3) | 5.332(3) |  11.735(2) | 90 | 333.509(6)
0.6 | 3.70 | P21/n - 45$\%$ | 6.818(9) | 7.596(5) | 8.030(0) | 105.21(7) | 401.383(6)
|  | I41/a - 55$\%$ | 5.329(8) | 5.329(8) | 11.714(8) | 90 | 332.787(0)
0.8 | 4.31 |  P21/n - 4$\%$ | 6.878(5) | 7.459(4) | 7.679(8) | 105.61(1) | 379.517(2)
|  | I41/a - 96$\%$ | 5.375(4) | 5.375(42) | 11.624(0) | 90 | 335.869(7)
1 | 4.91 | I2/a |  5.558(5) |  11.529(1) | 5.201(8) |  93.99(2) | 332.546(3)

Moreover, for the $x=$ 0 sample, the $m-m$ type crystal structure shows high
intensity diffraction peaks corresponding to (200) as well as (120) crystal
planes at 26.17$\degree$ and 27.78$\degree$, respectively. However, the $t-s$
type structure contains a peak corresponding to (112) plane at 28.08$\degree$,
and the $m-f$ type structure shows high intensity peaks for the
($\overline{1}$21) and (121) planes at 27.5o and 28.9o respectively. In the
measured XRD pattern for the $x=$ 0.2 to 0.8 samples, the diffraction peaks
for the (200), (120) and (112) planes are present, which clearly indicate the
co–existence of both the $m-m$ or $t-s$ type structures. The presence of (110)
plane at 17.65$\degree$ for the $x=$ 0.2 and 0.4 samples is due to the
dominance of $m-m$ type structure in the LaV1-xNbxO4. The (200) and (120)
peaks are also present in these samples; however, their intensity decreasing
with higher concentration of Nb substitution and become negligible for the
$x\geq$ 0.6 sample. As the Nb5+ concentration becomes more than V5+
concentration the $t-s$ type structure dominates, which results in the
reduction/absence of diffraction peaks corresponding to (200) and (120)
planes. The variation in peak intensity corresponding to (200) and (120)
crystal planes and the presence of (112) plane indicate the co–existence of
$t-s$ and $m-m$ type structures for the $x=$ 0.2 to 0.8 samples. This also
validates that the $m-m$ type structure (P21/n) is decreasing and $t-s$ type
structure (I41/a) is increasing with increasing the Nb concentration, i.e.,
from the $x=$ 0.2 to 0.8 samples. The determined $\%$ of phases by Rietveld
refinement of XRD data is presented in Table 1. For the $x=$ 1 sample, the
presence of ($\overline{1}$21) and (121) peaks further confirms the $m-f$ type
structure of LaNbO4 and consistent with literature [31].

Note that pure $m-m$ and $m-f$ phases are observed for the $x=$ 0 and 1
samples, respectively. However, for the $x=$ 0.2–0.8 samples, both the
monoclinic and scheelite-tetragonal phases coexist in a certain ratio. These
results reveal that LaV1-xNbxO4 samples undergo three phase transformation;
monoclinic monazite ($m-m$) type (for the $x=$ 0), two-phase equilibrium of
monoclinic monazite ($m-m$) and tetragonal scheelite ($t-s$) type
(0.2$\leq$$x$$\leq$0.8), and monoclinic fergusonite ($m-f$) type (for the $x=$
1) with increased substitution of Nb5+ at the V5+ site. It is quite
interesting to note that small amount of Nb5+ substitution can transform LaVO4
from $m-m$ phase to mix of $m-m$ and $t-s$ phases. It has also been observed
that LaNbO4 shows structural transition from monoclinic to a tetragonal phase
at $\sim$495$\degree$C. This structural transformation is very important in
governing the protonic conductivity of LaNbO4 [32]. For some compositions of
LaV1-xNbxO4, this transition temperature shifts near room temperature. The
reported temperature-dependent XRD measurements also suggest that at $x=$ 0.75
(25$\%$ substitution of V5+ at Nb+5 sites in LaNbO4) [21], it possess a
tetragonal structure at room temperature as its transition temperature is 250
K. The XRD pattern below 250 K shows some residual intensity (broadened lines)
of tetragonal structures because of precursor effects. Similarly, we can see
broad peaks in XRD patterns for the $x=$ 0.8 sample due to the above mentioned
effect [21]. As we increase the Nb concentration, we find some new peaks
appearing in the $x=$ 0.2 sample at 33.56o, 52.68o, 56.69o, and 58.06o. All
these peaks are the symbols of t–s structure that belongs to (020), (116),
(312), and (224) planes, respectively [33]. These peaks maintained up to the
$x=$ 0.8 sample, which confirms presence of some $t-s$ phase, and also
indicate the substitution-induced phase transformation. This is an important
finding that LaNbO4 can possess a tetragonal structure at room temperature by
just 20$\%$ replacement of Nb5+ sites by the V5+ sites. This result opens the
possibility for a wide range of applications of LaNbO4 at room temperature.
All these patterns discussed above suggest that substitution of larger Nb5+
($r=$ 0.48 Å) ion for V5+ ($r=$ 0.36 Å) affects the lattice constant of
LaV1-xNbxO4 and confirms the transformation of 3 different phases with
increasing concentrations of Nb5+.

Figure 2: The scanning electron microscope images of the LaV1-xNbxO4 ($x=$
0–1) samples.

The scanning electron microscope images of the LaV1-xNbxO4 for the $x=$ 0–1
are shown in Fig. 2, which depict the closed packed surface morphology in all
the samples and some variation in the particle size is clearly visible. The
pores are clearly visible from the top view of the surface. We can see that
with the increase in Nb5+ concentration, the particle size slightly decreases
from $x=$ 0 to $x=$ 0.4 sample, then increases and becomes maximum at $x=$ 0.8
and again decreases for the $x=$ 1 sample.

Figure 3: The HR-TEM images of LaV1-xNbxO4 for the (a) $x=$ 0.2 and (b) $x=$
0.8 samples. The magnified view of HR-TEM images in (c, d) for the $x=$ 0.2
sample, and (e, f) for the $x=$ 0.8 sample. (g, h) The SAED patterns for the
$x=$ 0.2 and 0.8 samples, respectively.

An average particle size (D) of LaV1-xNbxO4 is 5.14 $\micro$m for the $x=$ 0,
4.22 $\micro$m for the $x=$ 0.2, 3.56 $\micro$m for the $x=$ 0.4, 8.73
$\micro$m for the $x=$ 0.6, 11.31 $\micro$m for the $x=$ 0.8, and 5.70
$\micro$m for the $x=$ 1 samples. It is found that the change in crystal
surface morphology of LaV1-xNbxO4 samples with increasing Nb5+ concentration
causes variation in the particle size and shape.

Further, in Figs. 3 (a, b) we display the HR-TEM images indicating distinct
sets of planes with characteristic spacing for the $x=$ 0.2 and 0.8 samples.
The images in Figs. 3(c, d) and (e, f) for the samples $x=$ 0.2 and $x=$ 0.8,
respectively, show these plane sets in magnified view. The spacing between the
planes is determined using ImageJ software, and we find the $d-$spacings of
0.43 and 0.32 nm for the (-1,1,1) and (1,2,0) planes in the $P2_{1}/n$ phase
for the $x=$ 0.2 sample, and 0.28 and 0.31 nm for the (0,0,4) and (1,1,2)
planes in the $I4_{1}/a$ phase for the $x=$ 0.8 sample. However, these planes
only correspond to the dominating phase of the mixed-phase samples. The
selected area electron diffraction (SAED) patterns in Figs. 3(g, h) indicate
the contributions from both the phases. The indexed ($h,k,l$) planes that
relate to $P2_{1}/n$ are coloured in white and yellow colour is designated to
the $I4_{1}/a$ space group, as marked in Figs. 3(g, h). We find that the
analysis of HR-TEM and SAED results is consistent with the XRD refinement data
for these samples, as presented in Figs. 1(b, e).

Figure 4: The room temperature Raman spectra of LaV1-xNbxO4 ($x=$ 0 to 1) samples using (a) 532 nm, (b) 633 nm, and (c) 785 nm excitation wavelengths. The dotted blue lines represent the Lorentzian line shape to deconvolute the individual modes. Table 2: The experimentally observed frequencies $\omega_{obs}$ of the individual Raman modes in LaV1-xNbxO4 ($x=$ 0–1) samples measured at room temperature. $x$ |  0 |  0.2 |  0.4 |  0.6 |  0.8 |  1
---|---|---|---|---|---|---
Peak |  $\omega_{obs}$ |  $\omega_{obs}$ |  $\omega_{obs}$ |  $\omega_{obs}$ |  $\omega_{obs}$ |  $\omega_{obs}$
S0 |  Bg(127.24) |  |  |  |  |  Bg(121.86)
S1 |  Ag(141.99) |  Ag(140.66) |  Ag(140.66) |  Ag(140.66) |  |
S2 |  Bg(154.05) |  Bg(155.39) |  Bg(152.71) |  |  |
S3 |  Ag(187.44) |  Ag(187.44) |  Ag(186.11) |  |  |
S4 |  Bg(206.07) |  Bg(206.07) |  Bg(206.079) |  Bg(207.40) |  Bg(211.39) |  Ag(219.35)
S5 |  Ag(235.26) |  Ag(233.94) |  Ag(233.94) |  |  |
S6 |  Ag(245.84) |  Ag(244.52) |  Ag(244.52) |  |  |
S7 |  Bg(305.11) |  Bg(305.11) |  |  |  |
S8 |  Ag(326.07) |  Ag(326.07) |  Ag(326.07) |  Ag(324.76) |  Ag(328.69) |
S9 |  Ag(344.36) |  Ag(344.36) |  Ag(344.36) |  |  |
S10 |  Ag(370.41) |  Ag(370.41) |  Ag(369.11) |  Ag(369.11) |  |
S11 |  Bg(393.78) |  Bg(393.78) |  Bg(391.19) |  Ag(391.19) |  Ag(388.60) |  Bg(396.38)
S12 |  Ag(420.96) |  Ag(418.37) |  Ag(418.37) |  |  |
S13 |  Bg(436.44) |  Bg(435.15) |  Bg(433.86) |  Bg(436.44) |  |
S14 |  Ag(765.31) |  Ag(765.31) |  Ag(766.54) |  Ag(766.54) |  |
S15 |  Bg(788.68) |  Bg(788.68) |  Bg(788.68) |  Bg(788.68) |  |
S16 |  Ag(816.86) |  Ag(815.63) |  Ag(814.41) |  Ag(811.96) |  Ag(807.07) |  Ag(803.39)
S17 |  Ag(840.06) |  Ag(840.06) |  Ag(840.06) |  Ag(840.06) |  |
S18 |  Bg(855.88) |  Bg(855.88) |  Bg(854.67) |  Bg(854.67) |  |
S19 |  Bg(874.10) |  Bg(874.10) |  Bg(874.10) |  |  |
S20 |  |  |  |  Ag(170.10) |  Ag(168.76) |  Ag(174.10)
S21 |  |  |  |  |  |
S22 |  |  |  |  |  Ag(108.41) |  Ag(105.72)
S23 |  |  |  |  |  |  Bg(164.75)
S24 |  |  |  |  |  |  Bg(198.09)
S25 |  |  |  |  |  |  Bg(282.77)
S26 |  |  |  |  |  Bg(316.91) |  Ag(322.14)
S27 |  |  |  |  |  |  Ag(331.30)
S28 |  |  |  |  |  Ag(345.67) |  Bg(343.06)
S29 |  |  |  |  |  |  Bg(405.44)
S30 |  |  |  |  |  |  Ag(422.25)
S31 |  |  |  |  |  |  Bg(623.49)
S32 |  |  |  |  |  |  Ag(648.58)
S33 |  |  |  |  Bg(669.83) |  Bg(661.08) |  Bg(661.08)
Table 3: Summary of all the 34 Raman active modes and their assignments with the help of literature (cited in the last coloum of the table) for the LaVO4 and LaNbO4 samples. |  LaVO4 |  LaNbO4 |
---|---|---|---
Peak | $\omega_{th}$ | Assignments | $\omega_{th}$ | Assignments | Refs.
S0 |  Bg(127) |  Translation mode of La atoms in monoclinic phase |  Bg(125.1) |  Coupled translation-rotational mode of La atoms in monoclinic phase and NbO${}_{4}^{3-}$ around an axis perpendicular to b-axis | [43, 19, 45]
S1 |  Ag(143) |  Translation mode of La–O bonds |  |  | [47, 45]
S2 |  Bg(158) |  Translation mode of La–O bonds |  |  | [47]
S3 |  Ag(188) |  Translation mode of La–O bonds |  |  | [47, 45]
S4 |  Bg(204) |  Translation mode of La–O bonds |  Ag(222.3) |  Translational mode of La–O bonds along b-axis | [43, 48]
S5 |  Ag(230) |  Translation mode of La–O bonds |  |  | [47]
S6 |  Ag(252) |  Translation mode of La–O bonds |  |  | [47, 45]
S7 |  Bg(316) |  Bending vibration of O–V–O bonds |  |  | [49, 45]
S8 |  Ag(336) |  Bending vibration of O–V–O bonds |  |  | [49, 45]
S9 |  Ag(355) |  Bending vibration of O–V–O bonds |  |  | [49, 45]
S10 |  Ag(380) |  Bending vibration of O–V–O bonds |  |  | [49, 45]
S11 |  Bg(389) |  Bending vibration of O–V–O bonds |  Bg(398.1) |  Triply degenerate deformation mode (Rocking mode of NbO${}_{4}^{3-}$) | [43, 49, 45]
S12 |  Ag(423) |  Bending vibration of O–V–O bonds |  |  | [49, 45]
S13 |  Bg(427) |  Bending vibration of O–V–O bonds |  |  | [49, 45]
S14 |  Ag(784) |  Stretching vibration of V–O bonds |  |  | [50]
S15 |  Bg(799) |  Stretching vibration of V–O bonds |  |  | [50]
S16 |  Ag(806) |  Stretching vibration of V–O bonds |  Ag(805.2) |  Non degenerate Stretching mode of Nb–O bonds | [50, 43, 51, 52]
S17 |  Ag(836) |  Stretching vibration of V–O bonds |  |  | [50]
S18 |  Ag(861) |  Non degenerate stretching mode of shortest V–O bonds |  |  | [50, 45]
S19 |  Bg(892) |  Stretching vibration of O–V–O bonds |  |  | [51, 52, 45]
S20 |  |  |  Ag(177.1) |  Translational mode along b axis | [43]
S21 |  |  |  Bg(114) |  Rotational mode of NbO${}_{4}^{3-}$ along an axis perpendicular to b-axis | [43]
S22 |  |  |  Ag(108.6) |  Rotational mode of NbO${}_{4}^{3-}$ along b-axis | [43, 48]
S23 |  |  |  Bg(170) |  Translational mode parallel to ac-plane | [43]
S24 |  |  |  Bg(200.2) |  Translational mode parallel to ac-plane | [43]
S25 |  |  |  Bg(284.9) |  Translational mode parallel to ac-plane | [43]
S26 |  |  |  Ag(321.7) |  Doubly degenerate scissors mode of NbO${}_{4}^{3-}$) | [43, 48]
S27 |  |  |  Ag(326.9) |  Doubly degenerate scissors mode of NbO${}_{4}^{3-}$) | [43, 48]
S28 |  |  |  Bg(344) |  Translational mode parallel to ac-plane | [43]
S29 |  |  |  Bg(404.9) |  Triply degenerate deformation mode (Rocking mode of NbO${}_{4}^{3-}$) | [43]
S30 |  |  |  Ag(425.8) |  Triply degenerate deformation mode (Twist mode of NbO${}_{4}^{3-}$) | [43]
S31 |  |  |  Bg(625.5) |  One of triply degenerate Stretching mode of Nb–O bonds | [43, 51, 48]
S32 |  |  |  Ag(649) |  One of triply degenerate Stretching mode of Nb–O bonds | [43, 51, 48]
S33 |  |  |  Bg(664.9) |  One of triply degenerate Stretching mode of Nb-O bonds | [43, 51, 48]

The Raman spectra of LaV1-xNbxO4 measured at three different excitation
wavelengths, 532 nm, 633 nm, and 785 nm are presented in Fig. 4 for all the
samples ($x=$ 0–1). Three different excitation wavelengths are used to
distinguish the fluorescence effect on the Raman signal and to avoid the
background effects from the sample. We use Lorentzian line shape function to
deconvolute and fit the observed individual Raman peaks, as marked in Table 1.
We find that all the specific Raman peak positions (Raman shift) are
independent of excitation wavelength for a sample, which confirm their
inherent characteristic of that particular sample, as shown in Fig. 4. The
intensity of the modes may vary due to several reasons like the polarizability
of the molecule, excitation wavelength of the laser source, and the
concentration of the active group [34]. Though there are minor changes in the
intensity variation of Raman modes measured with different excitation
wavelengths, we can see that Raman active peaks are changing systematically
for all the measured samples in Fig. 4.

In the measured spectra we see 20 peaks corresponding to LaVO4 and 17 peaks
for LaNbO4. According to the Group theory calculations LaVO4 contains 72
vibrational modes and out of them, 36 modes are Raman active modes (18Ag \+
18Bg) [36, 37, 35] (here, A and B denote symmetric and antisymmetric
vibrations about the principal axis of symmetry and subscripts $g$ indicates
that the vibrations are symmetric relative to a symmetry center,
respectively). All the 20 Raman peaks for the $x=$ 0 sample are represented
from S0 to S19, as shown in the Table 2. The theoretical approach predicted
that the 8Ag+10Bg modes are for $m-f$ structure, and 13 Raman-active modes in
$t-s$ structure (as observed in the $x=$ 0.6 sample), which are summarized in
Table 2. The reason for the absence of some of the peaks could be due to the
overlap of several Ag and Bg modes and their low Raman scattering cross-
section. All the assignments related to each Raman peak in LaV1-xNbxO4 are
summarised in Table 3. We can see in Table 2 that the S0 mode (127.24 cm-1) is
present only in LaVO4 and LaNbO4 samples and absent for rest of the
intermediate samples. The reason for origin of S0 mode is translational motion
of La atoms in the monoclinic phase. All the concentrations from $x=$ 0.2 to
0.8 in LaV1-xNbxO4 results in the formation of $t-s$ type structure or $m-m$
and $t-s$ in equilibrium type structure. So, the formation of mixed phase may
result in the disappearance of S0 mode. The S18 is the most intense mode for
the LaVO4 sample, which decreases with Nb substitution. Whereas, we find that
the intensity of S16 mode increases with Nb substitution and the same becomes
the most intense mode for the LaNbO4 sample, as can be seen in Fig. 4(a). For
the $x=$ 0.8 sample, the S18 mode completely disappears, which indicates that
crystal phase transformation from a mixed phase of $m-m$ and $t-s$ in
equilibrium to turning into an approximately pure (96%) $t-s$ phase [38]. This
behaviour of S0, S16 and S18 corroborate with the structural phase
transformation with Nb+5 substitution, as has been observed in XRD analysis.
Furthermore, the presence of S8, S9, S10, S13, S14, S15, S17 and S18 modes in
the $x=$ 0 sample confirms the existence of VO${}_{4}^{3-}$ ions since none of
these modes are visible in LaNbO4 [39, 40, 41, 42].

All the Raman peaks arise due to different vibrational modes, i.e., bonds
between different constituent elements, i.e., the La3+, V5+, Nb5+ and O2-. The
comparison of experimentally observed peak positions of distinct Raman modes,
fitted using Lorentzian function, with the reported data [45, 44, 43, 9, 35],
shows a high degree of similarity, as presented in Table 2. In the $m-m$
structured LaVO4 crystal, nine O2- atoms are linked to La3+ whereas four O2-
atoms and V5+ are joined in a tetrahedral shape. There are four different O2-
locations and it is bound in a 3-coordinate geometry to two equivalent La3+
and one equivalent V5+ atoms at the first site. Also, it is bound to two
comparable La3+ and one equivalent V5+ atoms in a deformed single-bond
geometry in the second site. Three comparable La3+ and one equivalent V5+
atoms are linked to O2- in a 3-coordinate geometry at the third O2- site and
it is bound in a deformed single-bond geometry to three equivalent La3+ and
one equivalent V5+ atom in the fourth O2- site [46]. In the $m-f$ structured
LaNbO4 crystal, the La3+ is joined to eight O2- atoms in an 8-coordinate
geometry and six O2- atoms are bound to Nb5+ to create the deformed, edge-
sharing NbO6 tetrahedra. There are two different sites for O2- and it is
linked in a 4-coordinate geometry to two equivalent La3+ and two equivalent
Nb5+ atoms at the first O2- site. Also, it is bound in a 3-coordinate geometry
to two equivalent La3+ and one Nb5+ atoms at the second O2- site. In the
analysis of vibrational modes, it has been assumed that the LaNbO4 crystal is
made up of La3+ cations and NbO${}_{4}^{3-}$ molecular anions [43, 46]. It is
revealed experimentally that on the addition of Nb5+, it replaces V5+ from its
site and distorts LaVO4’s unit cell [30]. The mode of vibrations for LaVO4 is
categorised as follows: (I) the zone of high wavenumber (765–874 cm-1)
resulting from O-V-O bond’s stretching vibration (II) the intermediate
(305–436 cm-1) region resulting from O-V-O bond’s bending vibration, and (III)
the zone of low wavenumber ($<$ 285 cm-1) resulting from La atom’s translation
modes as the La atoms have high mass [9, 17], and the results are presented in
Table 3. Similarly, the vibrational modes of LaNbO4 are also categorized as
follows: (I) high wavenumber zone (623–803 cm-1) for stretching modes of Nb-O
bonds, (II) intermediate zone (322–422 cm-1) for deformation/scissor modes of
NbO${}_{4}^{3-}$, and (III) low wavenumber zone (121–282 cm-1) for rotational
modes of NbO${}_{4}^{3-}$ and translational lattice modes that include the
relative translations of anions and cations [43].

Figure 5: The room temperature XPS survey spectra of LaV1-xNbxO4 ($x=$ 0 to 1)
samples. Figure 6: (a) The La 3$d$ core-level spectra of LaV1-xNbxO4, $x=$ 0–1
samples. (b) The intensity ratio I2/I0, and (c) energy separation I2 \- I0 as
a function of doping level $x$. The fitted spin-orbit split components are
also shown for each sample. Figure 7: Tha Nb $3d$ core level spectra of
LaV1-xNbxO4, $x=$ 0–1 samples. The fitted spin-orbit split components are also
shown for each sample.

The LaNbO4 contains a total of three different modes; rotational modes of
NbO${}_{4}^{3-}$, vibrational modes of NbO${}_{4}^{3-}$ and translation modes
of La–O and O–La–O bonds. The S0, and S22 peaks are visible corresponding to
the combined translation-rotational (Bg) and rotational mode (Ag),
respectively, while the third rotational Bg mode (S21) is absent in the
observed experimental Raman spectra. The vibrational modes can be categorized
into (I) doubly degenerated scissor modes, (II) triply degenerated deformation
mode, which further splits into a pair of degenerated rocking mode and one
twist mode, (III) stretching mode, a non-degenerate and a triply degenerated
with increasing order of wave number [43]. The remaining modes are all
translational modes. From the Table 3, we can easily identify that LaNbO4
Raman modes are matching well with the reported one. Two NbO${}_{4}^{3-}$
scissor modes with almost degenerated wave numbers are projected to be seen in
the Ag spectrum. Out of all, the most obvious choices are S26 and S27 because
the remaining Ag band’s wavenumbers are too low to allocate them. In the
LaNbO4, as already discussed, the deformation modes are believed to be divided
into two almost degenerated rocking modes (S11 and S29) with Bg symmetry and a
twist mode (S30) with Ag symmetry. These modes are also present in the region
of intermediate wave numbers. The stretching modes are high energy vibrations
and here they are recognised as S16, S31, S32 and S33 peaks. As the non-
degenerate symmetric mode is expected to provide the strongest band, band S16
is allocated to it. The remaining S31, S32, and S33 peaks are assigned to
other three degenerate stretching modes. The invariance of the S4, S11 and S16
peak positions, all through, from the $x=$ 0 to 1 samples indicates no effect
on translational mode along the $b-$axis, the Bg frequency of VO${}_{4}^{3-}$
and NbO${}_{4}^{3-}$ for rocking and stretching modes. The S8 peak disappears
only in LaNbO4 spectrum because of absence of O-V-O bending vibrations [43,
17]. Interestingly, the S2, S3, S5, S6, S9, S12 and S19 peaks vanished just
before Nb concentration exceeds V (at $x=$ 0.4) and also S1, S10, S13, S14,
S15, S16, S17 and S18 peaks vanished just after Nb concentration became more
than V. It is quite possible that the low concentration of Nb in the sample
results in the weakening and then disappearance of some of the spectral peaks.
Due to the same reason some new peaks (S20 and S33) appeared in $x=$ 0.6–1
samples. Furthermore, the S20 peak arises due to translational mode along the
$b-$axis, and S33 peak appeared due to one of three triply degenerate
stretching modes of NbO${}_{4}^{3-}$ in the sample.

The most intense peak in LaNbO4 (S16) and LaVO4 (S18) at higher wavenumber is
due to the stretching of Nb-Ot and V-Ot bonds, where Ot represents the oxygen
atoms in the terminal position [53]. The terminal position of oxygen is that
where it connects the LaO8 dodecahedra and NbO6 octahedra in case of LaNbO4
and LaO9 muffin [54] and VO4 tetrahedra in case of LaVO4 [53]. Since the VO4
tetrahedra appears to be intrinsic in the peak broadening, it has been found
that this broadening in the Raman peaks spreads along samples with
intermediate Nb and V compositions. However, in certain samples, Nb5+ and V5+
cation-related variables may also play an important role in the increasing the
broadening of the peak. The broad peaks are made up of multiple modes which
are normally difficult to distinguish from one another [53]. The strongest
peak of LaVO4 (S18) is in the high wavenumber region which lies approximatly
52.5 cm-1 on higher side of the spectrum with respect to the strongest peak of
LaNbO4 (S16). This difference in wavenumber ($\Delta$) is related to average
bond length ($d$) of the atoms by $\Delta\propto 1/d^{3/2}$, as stated by the
Badger rule [55]. The bond lengths V–Ot in LaVO4 and Nb–Ot in LaNbO4 are
$\sim$1.72 Å [9] and $\sim$1.90 Å [53], respectively. The changes observed in
the Raman sepctra of the samples is quite consistent with the Badger’s rule.

Figure 8: Tha O $1s$ core level spectra of LaV1-xNbxO4, $x=$ 0–1 samples. Each
spectrum is shifted vertically for the clarity. Figure 9: Tha V 2$p$ core
level spectra of LaV1-xNbxO4, $x=$ 0–1 samples. The fitted spin-orbit split
components are also shown for each sample.

Finally we use x-ray photoemission spectroscopy (XPS) to investigate the
electronic structure by measuring the survey scan and particular elemental
core-level spectra of all the prepared samples. The identified peaks in the
survey spectra according to their binding energies are labeled and are in
agreement with reported values, as shown in Fig. 5. The characteristic La 3$d$
peaks cluster (830–870 eV), 4$d$ (4$d_{5/2}$ at 101 eV and 4$d_{3/2}$ at 104
eV), and 4$p$ (centered around 195 eV) [56]. The presence of these peaks of La
are clearly visible for every synthesised sample, and they are all remarkably
comparable. A consistent rise in Nb 3$d$ (discussed later) and Nb 3$p$
(3$p_{3/2}$ at 364 eV and 3$p_{1/2}$ at 379 eV) is observed along with an
increase in Nb doping and this feature of Nb is absent in the $x=$ 0 sample
[57]. For the V 2$p$ (2$p_{3/2}$ at 517 eV and 2$p_{1/2}$ at 525 eV) and V
2$s$ (630 eV) core-level peaks, the reverse behavior is anticipated and it is
quite evident as clearly visible in Fig. 5 [58]. The Voigt function has been
used to fit the core level spectra of the constituent elements. The fitted La
3$d$ core-levels are shown in Fig. 6(a). The spin-orbit splitting peaks
present in all the samples, have been de-convoluted at binding energies
834.3$\pm$0.2 eV, 836.0$\pm$0.3 eV, 838.7$\pm$0.1 eV, 847.9$\pm$0.1 eV,
851.1$\pm$0.2 eV, 853.0$\pm$0.3 eV, 855.6$\pm$0.1 eV, and 863.4$\pm$0.2 eV
(average B.E. of all the samples $\pm$ $\Delta$B.E., calculated for the $x=$
0–1 samples). The broad diffusive satellite peaks at 847.7 eV and 863 eV in
the locality of La 3$d$ core-level are coming from plasmons. Because of the
two final states I and II, and the subsequent spin-orbit splitting between
each state, making the structure complex. The primary strong peaks (3$d_{5/2}$
at 834.3 eV and 3$d_{3/2}$ at 851.1 eV, respectively) are associated with the
final state I (La4+ 3d94f0, L), which involves electron transfer to the
continuum from the 3$d$ core-level. The peaks at higher binding energies are
features of final state II (La3+ 3d94f1, L, -e) and this feature is
experimentally unresolved which indicates multiplet structure, as has been
suggested by Mullica et al. [56]. This corresponds to the electron transfer
from ligand (L, O2p in our case) valance band to the empty 4$f$ orbitals of La
[56, 59]. This multiplet structure of state II is composed of two bonding and
anti-bonding states. The prominent signals at higher binding energies
(3$d_{5/2}$ at 838.7 eV and 3$d_{3/2}$ at 855.6 eV) are due to bonding of
state II and the weak signals at lower binding energies (3$d_{5/2}$ at 836.0
eV and 3$d_{3/2}$ at 853 eV) are because of anti-bonding. The average energy
difference (over La core-level spectra of all the samples) between these three
pairs of peaks is nearly the same ( 16.9 eV) for the state I, state II
bonding, and state II anti-bonding, respectively. This verifies the unaltered
spin-orbit energy splitting of the states of La on Nb substitution [60].
Interestingly, we find a significant and systematic change in the intensity
variation of the peak at 838.7 eV (I2) relative to the primary peak at 834.3
eV (I0) with Nb doping. The metal-ligand orbital overlaps are reported to be
accountable for such doping-induced intensity variations [61, 62] where strong
ligands are found to populate the (La3+ 3d94f1, L, -e) state, intensifying I2
[63]. The intensity ratio I2/I0 is shown in Fig. 6(b), which shows a
consistent decrease as a function of doping $x$. This signifies that with Nb
substitution, the extent of overlapping between La(4f)-O(2p) orbitals
decreases monotonically. This conclusion can also be drawn from the trend in
the energy separation between I2 and I0 as a function of $x$, as shown in Fig.
6(c). However, the separation is minute in the subsequent samples, but for the
$x=$ 0 and $x=$ 1 the energy difference (I2 \- I0) is found to be of the order
of 0.3 eV. The value of I2 \- I0 was found to be varying for a variety of La-
containing compounds majorly because of the crystal structure, like 3.8 eV for
La0.5Sr0.5Co1-xNbxO3 and 5.3 eV La1.85Ba0.15CuO4 [60, 62]. Notably, this
energy separation could be related to the ease of electron transfer between
the ligand and the more ionic state of La, therefore having an opposite trend
with the tendency of ligand’s overlapping with the La $4f$ orbitals [62].

The Nb 3$d$ core level spectra are shown in Fig. 7 where the spin-orbit
doublet of Nb 3$d$ core levels are fitted with a single peak for each
component and the calculated peak positions for the Nb-doped samples are found
to be 3$d_{5/2}$ at 206.2$\pm$0.2 eV and 3d3/2 at 209.0$\pm$0.2 eV [64, 65].
This confirms the presence of prevailed 5$+$ oxidation state of Nb atom [57]
in all the samples. However, for the $x=$ 1 sample the Nb 3$d_{5/2}$ is at a
higher binding energy as compared to the other Nb-containing samples, which
could be due to the charging effects and the change in chemical environments.
Therefore, Atuchin et al. characterized the Nb state by using energy
difference $\Delta$ (Nb 3$d_{5/2}$ – O 1$s$) instead of solely relying on Nb
3d5/2 binding energy position [66]. The evaluated $\Delta$ (Nb 3$d_{5/2}$ – O
1s) values are found to be around 323.5 eV. The calculated energy difference
with respect to O 1$s$ is independent of the carbon correction. The obtained
binding energy difference $\approx$323.5 eV is reported to be a fairly highest
value for the 5+ oxidation state of Nb. We can also see that the error in the
value of binding energy position $\Delta$ is only 0.1 eV in this case, while
for the Nb 3d5/2, and O 1$s$, it is 0.3, and 0.2 eV, respectively. In Fig. 8
we can also see that the O 1$s$ peak is shifting to higher binding energy for
the $x=$ 1 sample as compared to the $x=$ 0\. Similarly, for the Nb 3$d_{5/2}$
core-level, it is shifting to higher binding energy. The $\Delta$(Nb 3d5/2 \-
O 1s) value for the $x=$ 1 sample is quite consistent for all the samples,
which strongly supports the electronic characterization using energy
difference with respect to O 1$s$ instead of absolute peak positions.

In Fig. 9, we present the V $2p$ core level spectra for all the samples, which
shows spin orbit components of 2$p_{3/2}$, and 2$p_{1/2}$ at 516.9 and 524.8
eV, respectively indicating V in ${5+}$ state. Interestingly, an unusual
broadening in the V $2p_{1/2}$ component is observed for all the samples,
whereas no such additional component is evident in the V $2p_{3/2}$ peak at
516.9 eV. More importantly, the deconvolution of the V $2p_{1/2}$ component
reveals that the FWHM of the higher energy feature (denoted by I) (1.2 eV) is
nearly the same as that of $2p_{3/2}$ component (1.1 eV). In contrast, the
lower energy feature (II) is significantly broader (2.8 eV). Moreover, the
area ratio of the combined I and II with $2p_{3/2}$ is close to 1/2, which
clearly indicates the intrinsic origin of these two features from the
vanadium. In contrast to metallic V $2p$ core-level, the vanadium based
compounds have often been reported to exhibit an anomalous V $2p_{1/2}$ width
as a consequence of Coster-Kronig (C-K) transitions [68]. The C-K type
transition is a class of Auger transition in which an electron from a higher
sub-shell of the same shell fills the core hole [67]. In the present case, the
filling of $2p_{1/2}$ core hole by an electron from $2p_{3/2}$ may give rise
to the C-K transitions and that can result in an additional feature in the
$2p_{1/2}$ component. Therefore, it is likely that the component I is
attributed to the core-hole recombination with the screening electrons,
analogous to the 2p3/2, whereas an additional L2-L3 (C-K) relaxation process
gives rise to the feature II in 2p1/2 peak [69]. No significant change in
these components has been observed with the Nb substitution, indicating the
robust nature of the underlined system. Further, the approach of O $1s$ energy
difference is also implemented in this case as for vanadium oxides, the energy
difference $\Delta$(V 2p3/2 \- O 1s) is an advantageous reference [70]. The
average $\Delta$(V 2p3/2 \- O 1s) magnitude is 12.8$\mp$0.1 eV, which is in
good agreement with the literature for V5+ oxidation state [71].

##  Conclusions

In conclusion, the solid state reaction method was used to prepare LaV1-xNbxO4
samples successfully with regular variable Nb5+ concentration. The XRD
measurements established that the substitution of larger Nb5+ ion for V5+
affects the lattice constant of LaV1-xNbxO4 and goes through three different
phase transformations [monoclinic monazite ($m-m$) type ($x=$ 0), two-phase
equilibrium of monoclinic monazite ($m-m$) and tetragonal scheelite ($t-s$)
type (0.2$\leq$$x$$\leq$0.8) and monoclinic fergusonite ($m-f$) type ($x=$
1)]. The SEM micrographs helped in analyzing that the particle size and shape
altered due to the change in crystal phases of these samples with increasing
Nb5+ concentration. The analysis of HR-TEM and SAED data found to be
consistent with the XRD refinement data. The Raman spectra of LaV1-xNbxO4 were
studied using 532 nm, 633 nm, and 785 nm excitation wavelengths. All the Raman
assignments were found to have well-ordered enhancement/diminution with the
increase in Nb5+ doping. The variation in the intensity as well as
appearance/disappearance of the Raman mode with Nb concentration are
coinciding with the change in the structural phases, as observed in XRD
analysis. This further confirms that the phase transformation in LaV1-xNbxO4
agrees with the maximum intensity peak patterns shown in the Raman spectra of
these samples and are consistent with the Badger’s rule. The XPS analysis
reveal the changes in Nb 3$d$ and V 2$p$ core-level spectral intensities of
the samples with the increase in Nb5+ concentration. The equal spin-orbit
energy splitting of the states was confirmed by the average energy difference
(over La core spectra of all samples) for state I, state II bonding, and state
II anti-bonding and the observed changes in their relative intensities with Nb
substitution are due to the metal ligand orbitals overlap. These findings
provide valuable insights into the structural and electronic properties of
LaV1-xNbxO4 samples and their potential use in different fields of practical
applications.

##  Acknowledgment

AS and MS thank MHRD and CSIR, respectively for the fellowship. The authors
acknowledge IIT Delhi’s FIST (DST, Govt. of India) UFO scheme for providing
the physics department with the Raman facility. We thank the Physics
Department at IIT Delhi for the XRD and the Central Research Facility (CRF)
for the FESEM, EDX, and XPS. We also thank Ambuj Mishra for providing HR-TEM
facility at IUAC, New Delhi. The preparation of the samples was done in a
high-temperature furnace (from Nabertherm GmbH, Germany), funded by BRNS
through the DAE Young Scientist Research Award (Project Sanction No.
34/20/12/2015/BRNS). RSD acknowledges SERB–DST for the financial support
through a core research grant (project reference no. CRG/2020/003436).

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