dhaou2021.pdf

Enhancement of dielectric responses and conduction
properties of Zn-doped TiO2 for energy storage
and photosensitivity applications
Mariem Dhaou1,2
, Elimame Elaloui1
, Kamel Khirouni3
, Hajer Guermazi2
, and Samir Guermazi2,
*
1
Research Unit: Materials Environment and Energy (UR14ES26), Science Faculty Sidi Ahmed Zaroug, 2112 Gafsa, Tunisia
2
Laboratory of Materials for Energy and Environment, and Modelling, Faculty of Science, University of Sfax, 3000 Sfax, Tunisia
3
Laboratory of Physics of Materials and Nanomaterials Applied to Environment, Faculty of Science, University of Gabes, Gabes,
Tunisia
Received: 12 February 2021
Accepted: 26 March 2021
Published online:
5 May 2021
© The Author(s), under exclu-
sive licence to Springer Science
+Business Media, LLC, part of
Springer Nature 2021
ABSTRACT
This work deals with the physical investigations on Zn-doped TiO2 synthesized
via a solvent-controlled non-aqueous sol–gel route. X-ray diffraction analysis
indicated that TiO2 particles crystallized in both tetragonal anatase and rutile
structures. FTIR analysis confirmed the insertion of Zn in TiO2 network and
revealed the presence of surface defects in the prepared powders, as proved by
their porous morphologies. UV–visible absorption was performed to provide an
insight into the band gap variation in the Zn-doped TiO2 nanopowders as a
function of Zn doping. On the other hand, the electrical properties were studied
using complex impedance spectroscopy in the frequency range from 40 Hz to
1 MHz at temperature range 480–600 K. The impedance plots are well fitted to a
(Rg//CPEg) − (Rgb//CPEgb) equivalent circuit. As well, the correlated barrier
hopping model (CBH) was proposed to describe the conduction mechanism.
Finally, Photocatalytic activities of the Zn-doped TiO2 nanoparticles were
evaluated via the degradation of the rhodamine B (RhB) dye under UV light
irradiation. The results showed enhanced performance by Zn doping, compared
to the undoped TiO2 nanoparticles.
1 Introduction
Titanium oxide TiO2 has been widely investigated
thanks to its good thermal, physical, and electrical
performances. It is a n-type semiconductor having a
wide band gap energy of 3.2 eV at room temperature
[1]. Also, it has a relatively high refractive index and
dielectric permittivity [2], making it promising for
various applications, such as positive electrode in
batteries, transparent conducting electrodes in pho-
tovoltaic devices, optoelectronic devices, gas sensors,
and photocatalytic degradation of organic dyes for
Address correspondence to E-mail: samir.guermazi@gmail.com
https://doi.org/10.1007/s10854-021-05858-x
J Mater Sci: Mater Electron (2021) 32:13187–13204

water and air purification, self-cleaning surfaces, and
antibacterial activity [3–7].
Currently, many complex architectures, especially
two-dimensional (2D) nanostructures become seri-
ously investigated. They usually exhibit a more
excellent optoelectronic and photovoltaic perfor-
mance due to their high charge mobility [8], which
promote transport of electrons as the sheets have
highly active surfaces. This, in turn, reduces the
recombination rate and improves the separation
capability, as mentioned in earlier reports [9,10]
Several physical and chemical methods have been
used to prepare TiO2 nanoparticles and thin films,
such as co-precipitation [11,12], polyol method [13],
RF sputtering [14], chemical vapor deposition (CVD)
[15], pulsed laser deposition [16], and spray pyrolysis
[17]. Moreover, many efforts have been conducted to
obtain nanopowders with desirable physical and/or
chemical properties. Among the preparation tech-
niques, sol–gel is cost-effective. It has many advan-
tages such as homogeneity, lower crystallization
temperature, a relative chemical purity of powders
and self-limiting surface reactions. Commonly, TiO2
contains different types of defects, due to oxygen
vacancies [18], oxygen interstitial [19], Ti vacancies
[20], Ti interstitial [21] and more complex defects [22].
Defects can also be monitored via appropriate sub-
stitution doping. Many works reported that small
concentrations of doping metals such as Ta, Al, Yb,
Nb, In, Sn, and Fe can significantly affect the opto-
electrical and dielectric properties of TiO2 [23–27].
This enhanced the charge storage mechanism by
introducing more oxygen in TiO2 network, which in
turn creates additional energy levels in the band gap
of TiO2. In addition, the surface morphology takes a
leading role. Thus, porous, interconnected, conduct-
ing materials and materials with high surface to
volume ratio such as nanocrystals can offer high
specific capacitance value. To date, many researches
have been focused on nanostructured TiO2 (thin
films, powders, nanocomposites, …) and it was
reported that their morphology plays a crucial role in
electric and dielectric properties due to their large
surface area [28,29].
In this context, the exploration of the dielectric
properties of the material is one of current research
topics, as it can predict whether the material is a
promising candidate for energy storage application.
As well, many studies on photoactivity of TiO2 oxide
have been reported since the early twentieth century.
Indeed, it was reported that UV absorption may
produce active oxygen species on TiO2 surface caus-
ing the photo-bleaching of dyes. Recently, ultrafine
nanoparticles of TiO2 have been synthesized. These
nanoparticles exhibit a noticeable photodegradation
efficiency in visible light of methylene blue (MB) dye
and complete photoreduction of Cr4?
ions [30]. In the
same line, Mn–Zn co-doped TiO2, synthesized by a
simple step sono-chemical route without annealing
treatment process, showed a relatively high photo-
catalytic degradation performance of rhodamineB
(RhB) dye under visible light [31].
Till now, the design of such nanoarchitecture
monolith materials has been considered a long-
standing challenge. So, by monitoring experimental
conditions, we reach a simple process to synthesize
dimensional nanomaterials with variable sizes, wide
surface area, porosity, and tunable physicochemical
characteristics such as light absorption, optoelec-
tronic enhancement, and dielectric properties with
the incorporation of Zn ions.
Although, to the best of our knowledge, those
properties of dimensional monoliths have not so far
been tackled, a growing interest is being devoted to
such prospects, more particularly with complicated
process by adding organic materials [32,33] or by
one- or two-step hydrothermal method [34,35],
coaxial electrospinning technique [36], and polymer-
ization method [37].
Among these various nanostructures currently
available, 2D nanosheets structures have attracted
considerable attentions, because of their massive
interior contact sites and large number of grain
boundaries. Despite intriguing, some questions about
the role of oxygen vacancies and microstructure on
optoelectrical properties and photoactivity still
remain unexplored. Herein, besides the synthetic
protocol and the physical investigations of Zn-doped
TiO2 nanopowders, a specific emphasis is put on
their photocatalytic properties against (RhB) dye
under UV illumination.
2 Experimental details
2.1 Material preparation
Undoped and 1 mol%, 2 mol%, and 3 mol% Zn-
doped TiO2 were synthesized via the sol–gel process.
For that, given amounts of Ti(OPr)4 and ethanol, in
13188 J Mater Sci: Mater Electron (2021) 32:13187–13204

(2:3) ratio, were mixed in a glass container. After the
complete mixing, glacial acetic acid was slowly
added to the resultant homogeneous yellow solution
in order to promote the acid hydrolysis at pH 3. In all
obtained samples, a stoichiometric amount of com-
mercial Zinc acetate (ZnC4H6O4, Aldrich) was intro-
duced after 60 min of stirring to obtain Zn-doped
TiO2. Thus, obtained sols were stirred, capped, and
kept at ambient conditions for more than 24 h, in
order to complete the gelation step, after which
transparent xerogels (monoliths) were obtained. We
proceeded to heat treatment of these powders.
Firstly, they were dried at 110 °C for 12 h, then
annealed at 500 °C for 5 h to obtain a thermally
stable crystalline phase. The gelation time for
undoped TiO2 was 60 min, while for Zn-doped TiO2,
the gelation needed more much time (about 15 days).
2.2 Characterization techniques
First, X-ray diffraction (XRD) patterns were recorded
to characterize the phase and crystal structure of
nanoparticles using a Phillips powder diffractometer
operating with copper Kα radiation source (λ=
1.54056 Å) over a wide range of Bragg angles
(10°2h 60°). Also, the structural analysis was
carried out using the standard Rietveld method [38].
The morphology of the TiO2 powders was analyzed
by a Hitachi S4100-1scanning electron microscope
(SEM). Second, Fourier transformed infrared (FTIR)
analysis was carried out using Perkin Elmer spec-
trometer in the wavenumber range 400–4000 cm−1
.
The optical absorption measurements of the powders
were recorded using a Shimadzu UV-3101 PC spec-
trophotometer in the UV–visible range (from 200 to
2000 nm wavelength). On the other hand, the elec-
trical properties of TiO2 nanoparticles calcined at
Fig. 1 X-ray diffractograms of a undoped, b Zn1%-, c Zn2%- and
d Zn3%-doped TiO2 monolith correspond to black dot, red solid
line refers to calculated (Rietveld fitting) and sky-blue solid line
refers to the difference between observed and Rietveld fitting data
(Color figure online)
J Mater Sci: Mater Electron (2021) 32:13187–13204 13189

500 °C were investigated using the impedance spec-
troscopy [39]. Prior to measurements, the TiO2 pow-
ders were pressed into pellets of 8 mm diameter and
1 mm thickness to form a plate capacitor configura-
tion. Aluminum films of 6 mm diameter were
deposited on both sides of the pellet. Afterwards,
electrical impedances were measured using a Tega-
m3550ALF impedance analyzer operating over a
frequency range 40 Hz–1 MHz and at temperature
480–600 K. Finally, the photocatalytic degradation
experiments were carried out by monitoring the
evolution of the main absorption of RhB aqueous
solution (λmax=554 nm) after irradiation with UV
light for 10–140 min. The absorbance spectra were
recorded by means of Shimadzu UV3100 spec-
trophotometer, using a high-pressure mercury lamp
(300 W) as a light source. Before illumination, the
solution was magnetically stirred for 60 min (in
dark), in order to ensure an adsorption equilibrium
between the sample surface and the organic dye [40].
A detailed description of the photocatalytic reactor
was previously reported [41,42]. The RhB dye pho-
todegradation efficiency (η) was calculated for the
main absorption of RhB solution as follows [43]:
gð%Þ ¼ 1
C
C0

100; ð1Þ
where C is the RhB dye concentration at time t of the
reaction and C0 is the concentration of the initial dye
solution (C0=30 mg/l), deduced from C/C0=A/A0
(A is the absorbance of the RhB dye solution).
3 Results and discussions
3.1 Structural analysis
Figure 1 shows the XRD patterns of undoped and
zinc-doped titanium dioxide nanopowders as a
function of Zn content. The diffraction peaks in all
samples matched well with the typical anatase
(A) and rutile (R) phases of tetragonal TiO2. Besides,
in 3 mol% Zn-doped TiO2 sample, a small trace of
brookite phase was formed, confirmed by the char-
acteristic (200) peak. All the diffraction peaks agreed
with the JCPDS cards (153-0152 for anatase, 900-7532
for rutile and 900-4139 for brookite phases). No
additional peaks characteristic of secondary phases
such as zinc oxide were detected.
Moreover, specific surface area (Sa), was estimated
for all the samples using the following equation:
Sa ¼
6
Dq
; ð2Þ
where D is the average crystallite size estimated
according to Scherrer model. ρ is the density of
samples calculated by the following relation [44]:
Table 1 Structural parameters of the anatase (A) and rutile (R) phases
Samples 0.0 mol% 1 mol% 2 mol% 3 mol%
Phases A R A R A R A R
Lattice constant (Å) a=3.78(3) a=4.59(2) a=3.78(3) a=4.59(2) a=3.78(4) a=4.59(2) a=3.78(3) a=4.59(3)
c=9.50(8) c=2.95(7) c=9.50(3) c=2.95(7) c=9.50(3) c=2.95(8) c=9.50(7) c=2.95(7)
Unit cell volume (Å3
) 136.1 62.3 136.0 62.3 136.1 62.3 135.1 62.4
Average crystallite size D (nm) 23 50 19 20 19 22 20 21
Density, ρ (g/cm3
) 3.89 4.24 3.89 4.25 3.89 4.28 3.89 4.24
Speciﬁc surface area Sa (105
cm2
/g) 6.31 2.72 9.08 7.84 9.08 6.40 7.70 6.73
Fig. 2 Speciﬁc surface area of TiO2 nanopowder as a function of
Zn content

q ¼
nM
NV
; ð3Þ
where n=4 for anatase and n=2 for rutile phase, M is
molecular weight, N is Avogadro’s number and V is
the unit cell volume.
The obtained values are listed in Table 1. The
density (ρ) of all the samples was found to have
almost the same value (about 3.9 g/cm3
for anatase
phase and about 4.25 g/cm3
for rutile phase), while
the specific surface area increased with the increase
of doping concentration. This may be attributed to
the decrease of the crystallite size. In fact, from Fig. 2,
we noticed that Zn-doped samples have larger
specific surface area, which is promising for high
photocatalytic efficiency. This will be confirmed in
the following sections.
Fig. 3 The SEM images and histograms of a undoped and b 1 mol% Zn-doped TiO2
Fig. 4 FTIR spectra of undoped TiO2 (a) and Zn-doped TiO2 with
1 mol% (b), 2 mol% (c), and 3 mol% (d)

3.2 SEM observations
The surface morphology of TiO2 powders was ana-
lyzed through SEM investigations. Figure 3a, b dis-
plays SEM images for undoped and 1 mol% Zn-
doped TiO2 NPs as an example. One can clearly
notice changes in shape and size of TiO2 nanoparti-
cles with doping. The SEM micrographs revealed the
presence of a small irregular spherical and slightly
elongated morphology of the undoped TiO2 particles
with an average particle size in the range of 140 nm.
However, Zn-doped nanopowders exhibit randomly
distributed 2D nanosheets with a nanometric thick-
ness. Furthermore, the porous texture of these sam-
ples was not uniform and it is obviously doping
dependent. It is clear that both macropores and
mesopores were included and interconnected. This
porous morphology would lead to an enhancement
in photocatalytic activities of Zn-doped TiO2
nanopowders.
3.3 FTIR spectroscopy
FTIR spectra of all samples are displayed in Fig. 4.
The broad band located around 3450 cm−1
was
attributed to the stretching vibration mode (ν(OH))
originated from the adsorbed water [45]. The later
was confirmed by the band around 1622 cm−1
which
attributed to the bending vibrations of H–O–H
(δ(OH)) [46]. Indeed, the porous surface morphology
proved by SEM images facilitate the water adsorption
at the surface of TiO2 nanoparticles. Furthermore, the
bands spread in the frequency domain between 500
and 700 cm−1
and are related to the presence of
metal–oxygen bonds (Ti–O–Ti), (Ti–O, Zn–), and (Ti–
O–Zn) [47]. In addition, we noticed a change in the
vibrational band shape with doping mainly for the
sample with 3% Zn doping. This can be related to the
disorder created by the incorporation of Zn atoms in
TiO2 network.
3.4 UV–visible analysis
In order to investigate the changes in optical transi-
tions of TiO2 NPs due to Zn doping, UV–visible
absorption measurement was performed. The
absorption spectra of the samples (Fig. 5a) were
computed from the diffuse reflectance spectra using
the Kubelka–Munk equation. The absorption at
around 285 nm corresponds to the electronic transi-
tion from 2p state of oxygen atoms in the valence
band, to the 3d state of Ti in the conduction band [48].
Moreover, from this figure, it is noted that the UV
light absorption increases with increasing Zn content,
which may be due to sp-d exchange interactions
between the band electrons and the localized d-elec-
trons of Zn2?
ions.
The band gap energies Eg of the Zn-doped TiO2-
NPs were estimated from the Tauc’s plots (Fig. 5b)
according to the following formula [49]:
ah# ¼ B h# Eg
n
; ð4Þ
where α is the absorption coefficient given by α=
2.303/A.d, with A is the absorbance and d is the
thickness of the sample, B is a constant and hm is the
photon energy. The exponent n value depends on the
nature of the electronic transition. It is equal to 1/2
Fig. 5 a Absorbance spectra of all TiO2 samples and b the
corresponding Tauc’s plot, in the insert: the evolution of the gap
energy Eg (eV) with Zn content in doped TiO2 NPs

for TiO2 as it is a direct band gap semiconductor [50].
The optical band gap was obtained by extrapolating
the linear part of the curve to the hm axis (Fig. 5b).
Results showed a slight red-shift of the bandgap of
Zn-doped TiO2 samples. It shifted from 3.37 eV for
undoped sample to 3.32 eV for 3 mol% Zn-doped one
(insert of Fig. 5b). The shrinkage of the band gap with
increasing Zn doping can be explained by the intro-
duction of new defect states in the band gap of TiO2.
Indeed, doping induces a new band distribution rel-
ative to Zn2?
and oxygen vacancies and consequently
can narrow the bandgap by forming a shallow level
states below the conduction band [51].
3.5 Impedance analysis
The Nyquist plots (− Z″ vs Z′) of TiO2 samples as a
function of temperature are shown in Fig. 6a–d. At
selected temperatures ranging from 480 to 600 K, the
figures exhibit deformed semicircles. They indicated
the existence of non-Debye relaxation phenomenon
with distribution of relaxation times. Non-Debye-like
behavior of individual electro-active regions in the
material can be modeled by a variety of circuit ele-
ments. The more common model is formed by com-
binations of resistance (R) and constant phase
element (CPE). The CPE component was proposed by
Abram et al. [52]. Its impedance is expressed as
follows:
ZCPE ¼ ½QðjxÞa

1
; ð5Þ
where j is the imaginary unit (j2
=−1), Q is a constant,
and α is a dimensionless parameter (0α1). It
characterizes the degree of deviation from ideal
Debye-like illustration.
Fig. 6 Nyquist diagrams of a undoped, b Zn1%-doped TiO2, c Zn2%-doped TiO2, and d Zn3%-doped TiO2, as a function of temperature

The experimental data of Fig. 6a–d are fitted using
the Z-view software. So, the best fit of experimental
data was obtained with an equivalent circuit that
consists of two parallel combinations of resistance
and CPE elements (Rg//CPEg in serial with Rgb//
CPEgb) as presented in the inset of this figure. The
equivalent circuit parameters estimated from the best
fit of experimental data are gathered in Table 2a–d.
These parallel combinations described the volume
(grains) and interfacial (grain boundaries) effects.
Therefore, the overlapped semicircles confirm the
dispersal nature of the relaxations and the strong
heterogenic nature of materials [53]. This result can
be related to the multiphase nature of TiO2 NPS
(anatase and rutile phases), as detailed in XRD anal-
ysis section.
The a values are close to the unit, which implies the
capacitive behavior of the Zn-doped TiO2.
When increasing temperature, the low-frequency
intersection point of the diagrams with the x-axis
shifts towards the origin of the diagram indicating
the decrease in the resistivity of the samples, assisted
by the increase of the mobility of charge carriers and
thus the conduction process [54]. In addition, the
presence of free charge carriers and impurities at the
grain boundaries can influence the electrical con-
ductivity. At high frequencies, the second semicircle
is very weak, which indicates the dominance of the
grain boundary contributions in conductivity [55].
Figure 7a–d shows the variation of real part (Z′) of
impedance as the function of frequency at different
temperatures for the all TiO2 samples. It is found that
the magnitude of Z′ decreases smoothly with the
Table 2 Equivalent circuit
parameters of the undoped and
Zn-doped TiO2 samples
T (K) Rg (Ω) CPEg-T (F) a Rgb (Ω) CPEgb-T (F) a
(a) The undoped TiO2 sample
480 1.6972E6 1.265E−10 0.941 238870 6.159E−11 1
500 3.0275E6 6.327E−11 0.966 69487 4.396E−9 0.886
520 1.6976E6 4.172E−11 1 138140 1.096E−10 0.943
540 1.2237E6 1.467E−11 1 112070 8.898E−11 0.949
560 1.1185E6 5.700E−11 1 111990 1.224E−10 0.939
580 2.1056E6 6.092E−11 0.977 16045 4.956E−10 0.942
600 9.0219E7 3.150E−11 1 89756 1.000E−10 0.946
(b) 1 mol% Zn-TiO2 sample
480 1.87136E6 2.8371E−10 0.927 67638.4 1.7094E−8 0.906
500 1.30297E6 2.0195E−10 0.968 74203.4 1.3860E−5 0.965
520 644067 2.5560E−10 0.940 6714.3 3.3200E−10 0.878
540 530456 1.9528E−10 0.972 2320.1 9.6051E−10 0.793
560 381832 2.7 E−19 0.984 11155.6 1.038E−9 0.808
580 142547.2 1.18 E−18 1 34157.8 1.894E−9 0.852
600 131131.6 1.6290E−10 0.988 114.9 1.385E−7 0.916
(c) 2 mol% Zn-doped TiO2 sample
480 367.5 2.8E−19 0.836 3.7585E6 3.359 E−11 0.989
500 444.4 1.85 E−18 0.714 1.8173 E6 3.419E−11 0.988
520 438.1 3.07 E−18 0.694 1.2107 E6 3.697 E−11 0.985
540 422 1.77 E−18 0.718 683270 3.520 E−11 0.989
560 437.1 2.7 E−19 0.821 3.86640 3.078 E−11 0.999
580 482.2 1.18 E−18 0.731 220600 3.248 E−11 0.995
600 489.9 1.12 E−18 0.729 144910 3.244 E−11 0.996
(d) 3 mol% Zn-doped TiO2 sample
480 2.8465E6 1.6274E−10 0.956 4.55183E21 1.2968E−9 0.966
500 1.47378E6 1.38769E−10 0.971 3073.4 1.4021E−9 0.973
520 0.891419E6 1.59415E−10 0.964 5.45342E6 1.94355E−5 0.975
540 0.558022 E6 9.43033E−10 0.858 1.32544E6 1.0359E−6 0.955
560 0.276128 E6 1.97192E−10 0.964 50375.8 1.7021E−9 0.842
580 0.189890 E6 3.91571E−10 0.904 23646.2 4.0972E−9 0.846
600 0.108993 E6 1.03916E−10 1 26441 2.2612E−10 0.942

increase in both frequency and temperature (i.e.,
NTCR behavior). The decrease of Z′ values with the
increase of frequency for the selected temperatures
may be due to release of space charges. The coinci-
dence of the Z′ values at high-frequency side at all
temperatures confirms the concept of space charge
effect [56]. Furthermore, the effect of doping on the
resistivity can be highlighted through the analysis of
the values of Z′ at low frequencies at a given tem-
perature. It is obvious that it increases by factor of
about ten in Zn-doped samples compared to
undoped TiO2.
In order to understand the charge motion mecha-
nism and the relaxation process as a function of
temperature, the frequency dependence of imaginary
part (Z″) of the complex impedance was plotted at
different temperatures (Fig. 8a–d). The (Z″) spectra
clearly show the relaxation peak which the frequency
and intensity are strongly temperature dependent.
Indeed, the relaxation peak shifts towards higher
frequencies with increasing temperature, whereas (Z
″)max value decreases. This indicates a thermally
activated dielectric relaxation process.
3.6 Dielectric analysis
The dielectric losses can be deduced from impedance
components as follows:
e00
¼
Z0
xC0 Z02
þ Z002
ð6Þ
Figure 9a–d displays the frequency dependence of
the imaginary part of dielectric constant. The dielec-
tric losses rise sharply at low frequencies, reach high
values at high temperatures, and decreases as fre-
quency increases. The decrease in high frequencies is
attributed to ionic inertia hindering the dipole to
follow the frequency response. Thus, the dielectric
Fig. 7 Frequency–temperature dependence of Z′ of a undoped, b 1 mol% Zn, c 2 mol% Zn, and d 3 mol% Zn-doped TiO2

losses approach a zero value. Furthermore, the
charge carriers are trapped by defects in grain
boundaries, forming dipole moments. Hence, the
hopping of charge carriers between defect centers
(Titanium interstitials and oxygen vacancies) at the
grain boundaries constitutes the dielectric polariza-
tion at lower frequencies.
In the explored frequency range, ε″(f) plots do not
show any loss peaks, due to the conduction phe-
nomena or/and electrode polarization which may
obscure a dielectric relaxation [57–59]. To overcome
this difficulty, we used the electric modulus formal-
ism, in which the electric modulus is defined as
follows:
M
¼
1
e
: ð7Þ
3.7 Electrical modulus analysis
The complex electric modulus representation M* (ω)
developed by Provenzano et al. [60] is one of the
methods to analyze the space charge relaxation phe-
nomena which is not clearly detected, when the
complex permittivity formalism is used. The electric
modulus M* is given by the following equation:
M
¼
1
e
¼
e0
e02
þ e002
þ j
e00
e02
þ e002
¼ M0
þ jM00
; ð8Þ
where M′, M″ are the real and the imaginary parts of
the electric modulus and e00
is expressed as follows:
e00
¼
ðes e1ÞðxsÞ1a
sin 1a
ð Þp
2

1 þ 2ðxsÞ1a
cos 1a
ð Þp
2

þ ðxsÞ2ð1aÞ
þ
r0
e0x
ð9Þ
Fig. 8 Frequency–temperature dependence of − Z″ of a undoped, b 1 mol% Zn, c 2 mol% Zn, and d 3 mol% Zn-doped TiO2

r0 is the static conductivity, α is a parameter between
0 and 1, εs and ε∞ are, respectively, the static and
infinite frequency dielectric constants, ω is the
angular frequency (ω=2πf), τ is the relaxation time,
and ε0=8.8510–12
F/m is the permittivity of vacuum.
As it can be seen, the electric modulus can greatly
decline the electrode polarization and conduction
effects, which appear to obscure relaxation in the
permittivity presentation [61].
Figure 10 shows the variations of imaginary part of
electric modulus M″ with frequency at different
temperatures. It exhibits a relaxation peak that shift
toward the higher frequency side with an increase in
temperature. This indicates a thermally activated
behavior of the relaxation.
The peak is assigned to the transition from long
range to short range mobility of charge carriers with
the increase in frequency [62]. In fact, it is associated
to the dipolar relaxation of the TiO2 grains [63,64].
This peak is usually related to oxygen interstitial Oi,
oxygen vacancy VO, and titanium interstitial Tii.
These observations are in good agreement with the
results obtained by Ben Taher et al. [65]. Besides, the
asymmetry in the relaxation peak indicates distribu-
tion of the relaxation time, and hence the relaxation in
the material is of non-Debye type. Furthermore, it is
obvious that Zn-doped TiO2 exhibit more dispersive
relaxation behavior vs temperature, compared to the
undoped TiO2 (Fig. 10), due to structural disorder
induced by insertion of Zn ions into TiO2 network.
3.8 Conductivity study
The ac conductivity (σ) of the material can be calcu-
lated using the following relation:
Fig. 9 Variation of the imaginary part of dielectric constant of a undoped, b 1 mol% Zn, c 2 mol% Zn, and d 3 mol% Zn-doped TiO2

r ¼ xe0e00
ð10Þ
The (σ) plots of undoped and Zn-doped TiO2
samples are displayed in Fig. 11a–d. The conductivity
is almost constant at low frequencies, then increases
with the increase of frequency. At low frequencies,
the frequency-independent values of conductivity
correspond to the dc conductivity. Moreover, an
increase in σdc is noticed when increasing tempera-
ture. In fact, at low frequencies, few charge carriers
can tunnel through the potential barrier at the grain
boundaries, related with low conductivity. More
charge carriers tunneling let to enhance the conduc-
tivity at higher temperatures and frequencies.
Beyond a definite value of frequency, the charge
carriers get sufficient energy to overcome the poten-
tial barrier and hence a rapid increase in conductivity
occurs at higher frequencies [66]. Furthermore, it is
clearly seen from Fig. 11e that dc conductivity shar-
ply decreases in the doped samples from about 5.10–
9
Ω−1
cm−1
in the undoped TiO2 to 7.10–10
Ω−1
cm−1
in
3% Zn-doped TiO2. This can be related to Zn2?
incorporated into a TiO2 lattice. Indeed, as their ionic
radius is almost equal (rTi
4?
=0.068 nm and rZn
2?
=
0.074 nm), Zn can easily substitute Ti ions. Therefore,
Zn ions cause defects like titanium interstitials and
oxygen vacancies in the host TiO2 system, which act
as deep charge traps [67], thus reducing the charge
mobility and consequently decreasing the conduc-
tivity. Furthermore, several reports [68,69] approved
that the dopants confined in interstitial sites can act
as trapping or recombination centers for excited
electrons and holes [70].
The ac conductivity can be described by the fol-
lowing equation:
Fig. 10 Frequency dependence of the imaginary part M″ of a undoped, b 1 mol% Zn, c 2 mol% Zn, and d 3 mol% Zn-doped TiO2

rðxÞ ¼ rdc þ Axs
ð11Þ
known as Jonscher’s universal power law [71], where
A is a constant and s is a temperature and frequency
dependent exponent (0s1). The exponent “s”
characterizes the degree of interactions between
charge carriers and their environment. The second
term of Eq. (11) (Aωs
) describes the frequency
dependence of ac conductivity and characterizes all
dispersion phenomena; besides “s” gives information
about charge transport. At low frequencies, the
temperature dependence of dc conductivity is given
as follows:
rdcðTÞ ¼ r0expðEa=KBTÞ; ð12Þ
where σ0 is a pre-exponential factor which includes
the charge carrier mobility and density of states, Ea is
the activation energy associated with dc conductivity,
KB is Boltzmann constant, and T is the absolute
temperature.
Figure 12 displays the Arrhenius plots of the dc
conductivity (Eq. (12)), in which the slopes permit the
evaluation of the activation energy for the all TiO2
Fig. 11 Frequency dependence of the AC conductivity for: a undoped, b 1% Zn-doped TiO2, c 2% Zn-doped TiO2, and d 3% Zn-doped
TiO2, as a function of temperature, and e the dc conductivity at 480 K, as well as its activation energy vs Zn doping

samples. So, the activation energy was found to
increase from 0.37 eV for undoped sample to 1.09 eV
for 3% Zn-doped TiO2, as seen in Fig. 11e, in accor-
dance with the decrease of the dc conductivity. This
can be explained by a reduction in the charge car-
rier’s mobility in doped TiO2 due to the creation of
deep level defects induced by the insertion of Zn in
TiO2 structure. It is reported that the major defects
that arise in TiO2 are titanium interstitials (Tii) and
oxygen vacancies (VO) [72]. According to the litera-
ture [73], oxygen vacancies form two energy levels at
0.75 eV and 1.18 eV below the conduction band,
corresponding, respectively, to singly and doubly
ionized oxygen vacancies (VO). Correspondingly, the
calculated activation energy of 0.37 eV, for the
undoped TiO2 sample, nearly matches that of the first
ionization energy of oxygen vacancies (VO
−
). While in
Zn-doped TiO2 the Ea values were in accordance with
the doubly ionized oxygen vacancies (VO
=
). So, as
previously mentioned, the obtained Ea values of the
undoped and Zn-doped samples can be attributed to
singly and doubly ionized oxygen vacancies (VO)
respectively. Thus, for doped TiO2, the increase of the
activation energy originated from Zn2?
state charges
incorporated into TiO2 matrix. Indeed, Zinc atoms act
as donors by supplying two free electrons, when the
Zn2?
ions occupy Ti4?
ion site.
Generally, the variation of the exponent “s”
(Eq. 11) is related to the conduction mechanism. This
exponent “s” was calculated from the linear parts of
(σ=f(f)) at high frequencies (Fig. 12). The decrease of
“s” with temperature (Fig. 13) suggests that the cor-
related barrier hopping (CBH) mechanism governs
the ac conduction [74–76]. This reflects the hopping of
charge carriers between two sites over a barrier sep-
arating them.
3.9 Photocatalytic tests
The photocatalytic activities of undoped and Zn-
doped TiO2 were evaluated using the degradation of
the standard organic RhB dye, under UV irradiation.
For that the absorption of RhB dye solution with TiO2
NPs was measured and typical absorption spectra are
plotted over the time in Fig. 14a. It is obvious that
TiO2 NPs react well leading to a decrease of the main
absorption band of RhB solution over the time.
Indeed, RhB is a phenothiazine dye in which the
chromophore part contains amino and thiocarbonyl
bonds and exhibits a strong absorbance in the visible
region (554 nm). However, the vanishing of this band
reveals that the dye is degraded. Also, the photocat-
alytic activity exhibits a monotonous improvement
by increasing Zn content and reaches the highest
efficiency with 3% Zn doping (Fig. 14b). Moreover,
Fig. 14c clearly shows that all TiO2 samples have high
photocatalytic performances against RhB after
140 min under UV irradiation, which was clearly
improved by 3% Zn doping. In fact, for this sample
the photodegradation efficiency reached 95% while it
was of 82% for the undoped TiO2. This fact may be
related to the small average crystallite size (Table 1)
and a large specific area (Fig. 2) of the 3% Zn-doped
Fig. 12 Ln ((σdc T) versus reciprocal temperature for undoped
and Zn-doped TiO2 samples
Fig. 13 Variation of s for undoped and Zn-doped TiO2 samples

sample. The latest can help diffusion of the reactive
molecules to the active sites, as the samples have
porous morphology, which leads to an increase in the
charge of the dye molecules and thus activates dye
photodegradation.
Furthermore, from Fig. 14b, one can notice a dif-
ference in photodegradation kinetics of RhB using
undoped and Zn-doped TiO2 nanopowders. Indeed,
the kinetic rate is an important parameter in pho-
todegradation studies because it can calculate the rate
at which pollutant is removed from the aqueous
solution. So, it is reasonable to quantify the
photodegradation kinetics of the RhB dye. In all
cases, the degradation kinetics are well adjusted to
pseudo-first-order process, which can be evaluated
using Langmuir and Hishelwood formula [77]:
LnðC0=CÞ ¼ Kapp:t; ð13Þ
where Kapp is the apparent rate constant of the first-
order reaction (min−1
).
A plot of (Eq. (13)) is given in Fig. 14d. It exhibits a
straight line in which the slope equals the apparent
first-order rate constant Kapp. The obtained values are
plotted versus Zn doping content and presented in
the insert of Fig. 14d. As shown, the 3% Zn-doped
Fig. 14 a Absorbance spectra of RhB under catalyst and UV
irradiation, b photodegradation efﬁciencies (C/C0) vs. irradiation
time, c photodegradation efﬁciency after 140 min under UV
irradiation, for undoped and Zn-doped TiO2 d Ln(C0/C) vs.
irradiation time, and e the mechanism of dye degradation

TiO2 sample has the highest rate constant, which
degrades the RhB dye faster. This proves the
enhancement of photocatalytic performance of TiO2
by Zn doping.
Following the above discussion, the proposed
mechanism of dye degradation is shown in Fig. 14e.
When irradiated by UV light, valence electrons of
TiO2 were excited to the conduction band (CB),
resulting on generation of electron–hole pairs. The
electrons in the CB react with dissolved oxygen and
produce reactive oxygen species, which react with
water molecules to form hydroxyl radicals (·OH). The
dye is then degraded by the hydroxyl groups. While
holes at the valence band of TiO2 react with absorbed
water or hydroxyl groups to form surface hydroxyl
radicals which then degrade dye. The holes can oxi-
dize the dye molecules directly. The main reactions
are shown below [78]:
TiO2 ! e
þ hþ
ð14Þ
TiO2 e
ð Þ þ O2 ! TiO2 þ O
2 ð15Þ
TiO2 hþ
ð Þ þ H2O=OH
! TiO2 þ OH
ð16Þ
OH
þ RhB ! degradation product ð17Þ
the electron in the CB from reacting with molecular.
4 Conclusion
In summary, undoped and Zn-doped TiO2 has been
successfully prepared by sol–gel method. All syn-
thesized nanopowders are formed by a mixture of
rutile and anatase crystalline phases with tetragonal
structure. Zn doping highly modified the powder
morphology and leads to the formation of nanosh-
eets. The specific surface area increased with the
increase of doping concentration. The optical band
gap decreases with the increase of Zn content, due to
defects creation and the decrease of the average
crystallite size. Dielectric properties were analyzed
using complex impedance for the sample at various
temperatures. The impedance plots are well fitted to
a (Rg//CPEg) − (Rgb//CPEgb) equivalent electrical
circuit. The equivalent circuit parameters estimated
from theoretical fit proved the enhancement of
capacitive behavior of the Zn-doped TiO2. The fre-
quency dependence of the conductivity is interpreted
in terms of Jonscher’s law. As well, the conduction
mechanism was dominated by the correlated barrier
hopping model (CBH). Finally, photocatalytic
activities of the Zn-doped TiO2 nanoparticles were
evaluated via the degradation of RhB dye aqueous
solution under UV light irradiation. Also, the results
showed that Zn-doped TiO2 nanopowders exhibit a
noticeable improvement of photodegradation of RhB
dye typically for 3% mol Zn doping, compared to the
undoped TiO2 nanoparticles. Its photodegradation
efficiency reaches 95% after irradiation time of
140 min. The kinetic study revealed that 3% mol Zn-
doped TiO2 sample is a promising photocatalyst for
the degradation of RhB pollutant. Further works are
planned to test these samples in other sensitivity
applications such as gas and bio-sensors.
Acknowledgements
Authors gratefully thank the financial support of the
Tunisian Ministry of High Education and Scientific
Research.
Author contributions
MD: Conceptualization, Methodology, Software, and
Writing—Original draft preparation. KK: Dielectric
measurements. HG: Visualization, Investigation, and
Writing–Reviewing. EE: Visualization, Validation,
Supervision, and Writing–Reviewing. SG: Writing–
Reviewing and Editing, Validation, and Supervision.
Funding
This research was supported by the Tunisian Minis-
try of High Education and Scientific Research.
Declarations
Conflict of interest The authors declare that they
have no known competing financial interest or per-
sonal relationships that could have appeared to
influence the work reported in this manuscript.
References
1. N.S.P. Bhuvanesh, J. Gopalkrishnan, J. Mater. Chem. 7, 2297
(1997)
2. T. Fuyuki, H. Matsunami, Jpn. J. Appl. Phys. 25, 1288 (1986)

3. C. Guo, M. Ge, L. Liu, G. Gao, Y. Feng, Y. Wang, Environ.
Sci. Technol. 44, 419 (2009)
4. X. Li, J. He, A.C.S. Appl, Mater. Interfaces. 5, 5282 (2013)
5. Z. Wang, X. Peng, C. Huang, X. Chen, W. Dai, X. Fu, Appl.
Catal. B 219, 379 (2017)
6. J. Reszczyńska, T. Grzyb, J.W. Sobczak, W. Lisowski, M.
Gazda, B. Ohtani, A. Zaleska, Appl. Surf. Sci. 307, 333–345
(2014)
7. Z. Wei, M. Endo, K. Wang, E. Charbit, A. Markowska-
Szczupak, B. Ohtani, E. Kowalska, Chem. Eng J. 318, 121
(2017)
8. Y. Chen, Z. Lao, B. Sun, X. Feng, S.A.T. Redfern, H. Liu, J.
Lv, H. Wang, Z. Chen, Identifying the ground-state NP sheet
through a global structure search in two-dimensional space
and its promising high efficiency photovoltaic properties.
ACS Mater. Lett. 1, 375–382 (2019)
9. L. Sheng, T. Liao, L. Kou, Z. Sun, Mater. Today Energy 3, 32
(2017)
10. Y. Wang, J. Zhai, Y. Song, Phys. Chem. Chem. Phys. 17,
5051 (2015)
11. J. Ben Naceur, M. Gaidi, F. Bousbih, R. Mechiakh, R.
Chtourou, Curr. Appl. Phys. 12, 422 (2012)
12. Y. Shang-Wei, K. Horng-Huey, C. Hsiu-Mei, C. Yen-Ling, L.
Jian-Hong, W. Chiu-Ming, W. Moo-Chin, J. Alloys Compd.
613, 107 (2014)
13. R. Bargougui, N. Bouazizi, W. Ben Soltan, A. Gadri, A.
Azzouz, S. Ammar, Appl. Phys. A (2016). https://doi.org/10.
1007/s00339-016-9837-9
14. H. Huang, W. Yang, Y. Xie, X. Chen, Z. Wu, IEEE Electron
Device Lett. 31, 588 (2010)
15. A.M. Alotaibi, S. Sathasivam, B.A.D. Williamson, A. Kafi-
zas, C. Sotelo-Vazquez, A. Taylor, D.O. Scanlon, I.P. Parkin,
Chem. Mater. 30, 1353 (2018)
16. H. Lu, S. Fang, J. Hu, B. Chen, R. Zhao, H. Li, C.M. Li, J.
Ye, ACS Omega (2020). https://doi.org/10.1021/acsomega.0c
02838
17. I. Dundar, M. Krichevskaya, A. Katerski, I.O. Acik, R. Soc,
Open Sci. 6, 181578 (2019)
18. X. Pan, M.Q. Yang, X. Fu, N. Zhang, Y.J. Xu, Nanoscale 5,
3601–3614 (2013)
19. K.M. Pangan-Okimoto, P. Gorai, A.G. Hollister, E.G. See-
bauer, J. Phys. Chem. C 119, 9955 (2015)
20. Z.-Y. Shen, Q.-G. Hu, Y.-M. Li, Z.-M. Wang, W.-Q. Luo, R.-
H. Liao, J. Mater. Sci. Mater. Electron. 24, 3089 (2013)
21. B. Santara, P.K. Giri, K. Imakita, M. Fujii, J. Phys. Chem.
117, 23402 (2013)
22. Y. Shi, H. Sun, M.C. Nguyen, C. Wang, K. Ho, W.A. Saidi, J.
Zhao, NLM 9, 11553 (2017)
23. W. Tuichai, S. Danwittayakul, N. Chanlek, P. Thongbai, J.
Alloys Compd. 725, 310 (2017)
24. H. Peng, B. Shang, X. Wang, Z. Peng, X. Chao, P. Liang, Z.
Yang, Ceram. Int. 44, 5768 (2018)
25. X. Zhao, P. Liu, J. Alloys Compd. 715, 170 (2017)
26. N. Al Saqri, A. Mondal, J.F. Felix, Y.G. Gobato, V.O. Gordo,
H. Albalawi, D. Jameel, H. Alghamdi, F. Al Mashary, D.
Taylor, M.S. Abd El-sadek, M. Henini, J. Alloys Compd. 698,
883 (2017)
27. Q. Cai, Y. Zhang, C. Liang, P. Li, H. Gu, X. Liu, J. Wang, Z.
Shentu, J. Fan, G. Shao, Electrochim. Acta 261, 227 (2018)
28. K. Elghniji, A. Atyaoui, S. Livraghi, L. Bousselmi, E. Gia-
mello, M. Ksibi, J. Alloys Compd. 541, 421 (2012)
29. Y. Yu, F. Yang, S. Mao, S. Zhu, Y. Jia, L. Yuan, M. Salmen,
B. Sun, Effect of anodic oxidation time on resistive switching
memory behavior based on amorphous TiO2 thin films
device. Chem. Phys. Lett. 706, 477–482 (2018)
30. I. Elmehasseb, S. Kandil, K. Elgendy, Optik (2020). https://d
oi.org/10.1016/j.ijleo.2020.164654
31. C. Wattanawikkam, W. Pecharapa, Radiat. Phys. Chem.
(2020). https://doi.org/10.1016/j.radphyschem.2020.108714
32. R. Sui, C.E. Deering, R. Prinsloo, C.B. Lavery, N. Chou, R.A.
Marriott, Faraday Discuss. (2020). https://doi.org/10.1039/
C9FD00108E
33. N.N. Rabin, S. Ida, M.R. Karim, Md. Saidul Islam, R. Ohtani,
M. Nakamura, M. Koinuma, L.F. Lindoy, S. Hayami, A.C.S
Omega 3, 2074–2083 (2018)
34. L. Lai, E. Lei, C. Hu, D. Zhao, W. Zhao, Z. Guo, D. Huang,
Mater. Res. Express 7, 015053 (2020)
35. L. Cao, R. Wang, D. Wang, X. Li, H. Jia, Mater. Lett. 160,
286 (2015)
36. Z. Pan, L. Yao, J. Zhai, K. Yang, B. Shen, H. Wang, ACS
Sustain. Chem. Eng. 5(6), 4707 (2017)
37. F. Wen, H. Lou, J. Ye, W. Bai, L. Wang, L. Li, W. Wu, Z. Xu,
G. Wang, Z. Zhang, L. Zhang, Compos. Sci. Technol. 182,
107759 (2019). https://doi.org/10.1016/j.compscitech.2019.
107759
38. H.M. Rietveld, J. Appl. Crystallogr. 2, 65 (1969)
39. E. Barsoukov, J.R. Macdonald, 2nd edn. in Impedance
Spectroscopy Theory, Experiment, and Applications (Wiley,
Hoboken, 2005), pp. 1–20
40. H. Bel Hadjltaief, S. Ben Ameur, P. Da Costa, M. Ben Zina,
M.E. Galvez, Appl. Clay Sci. 152, 148 (2018)
41. S. Ben Ameur, H. Bel hadjltaief, A. Barhoumi, B. Duponchel,
G. Leroy, M. Amlouk, H. Guermazi, Vacuum 155, 546 (2018)
42. S. Ben Ameur, A. Barhoumi, H. Belhadjltaief, R. Mimouni,
B. Duponchel, G. Leroy, M. Amlouk, H. Guermazi, Mater.
Sci. Semicond. Process. 61, 17 (2017)
43. S. Ben Ameur, H. BelHadjltaief, B. Duponchel, G. Leroy, M.
Amlouk, H. Guermazi, S. Guermazi, Heliyon 5(6), e01912
(2019)

44. A. Maurya, P. Chauhan, S.K. Mishra, R.K. Srivastava, J.
Alloys Compd. 509, 8433 (2011)
45. A. León, P. Reuquen, C. Garín, R. Segura, P. Vargas, P.
Zapata, P.A. Orihuela, Appl. Sci. 7, 49 (2017)
46. P. Singla, O.P. Pandey, K. Singh, Int. J. Environ. Sci. Technol.
13, 849 (2016)
47. Z. Zheng, B. Huang, X. Meng, J. Wang, S. Wang, Z. Lou, Y.
Dai, Chem. Commun. 49, 868 (2013)
48. V.R. Akshay, B. Arun, S. Dash, A.K. Patra, G. Mandal, G.R.
Mutta, A. Chanda, M. Vasundhara, R.S.C. Adv. 8, 41994
(2018)
49. R.B. Bylsma, M. Becker, J. Kossut, U. Debska, Phys. Rev. B
33, 8207 (1986)
50. M. Ismael, New. J. Chem. 43, 9596 (2019)
51. M. Bärtsch, M. Niederberger, ChemPlusChem 82, 42 (2017).
https://doi.org/10.1002/cplu.201600519
52. E.J. Abram, D.C. Sinclair, A.R. West, J. Electroceram. 10,
165 (2003)
53. N. Kılınc, S. Ozturk, L. Arda, A. Altındal, Z.Z. Ozturk, J.
Alloys Compd. 536, 138 (2012)
54. J. Plocharski, W. Wieczoreck, Solid State Ion. 28, 979 (1988)
55. D.C. Sinclair, A.R. West, J. Mater. Sci. 29, 6061 (1994)
56. W. Dong, W. Hu, A. Berlie, K. Lau, H. Chen, R.L. Withers,
Y. Liu, ACS Appl. Mater. Interfaces 7, 25321 (2015)
57. J.B. Jorcin, M.E. Orazem, N. Pebere, B. Tribollet, Elec-
trochim. Acta 51(8–9), 1473–1479 (2006)
58. P. Pissis, A. Kyritsis, Solid State Ion. 97, 105 (1997)
59. M.H. Khan, S. Pal, E. Bose, Appl. Phys. A 118, 907 (2015)
60. V. Provenzano, L.P. Boesch, V. Volterra, C.T. Moynihan, P.B.
Macedo, J. Am. Ceram. Soc. 55, 492 (1972)
61. M.H. Abdullah, A.N. Yusoff, J. Mater. Sci. 32, 5817 (1997)
62. R. Ranjan, R. Kumar, N. Kumar, B. Behera, R.N.P. Choud-
hary, J. Alloys Compd. 509, 6388 (2011)
63. I.M. Hodge, C.A. Angell, J. Chem. Phys. 67, 1647 (1977)
64. A. Ray, A. Roy, S. De, S. Chatterjee, S. Das, J. Appl. Phys.
123, 104102 (2018)
65. Y. Ben Taher, N. Moutia, A. Oueslati, M. Gargouri, RSC Adv.
6, 39750 (2016)
66. M. Ashokkumar, S. Muthukumaran, J. Lumin. 162, 97 (2015)
67. X. Li, X. Cao, L. Xu, L. Liu, Y. Wang, C. Meng, Z. Wang, J.
Alloys Compd. 657, 90 (2016)
68. Wu. Zhao-Jin, W. Huang, K.-K. Cui, Z.-F. Gao, P. Wang, J.
Hazard. Mater. 278, 91 (2014)
69. X. Qiu, L. Li, J. Zheng, J. Liu, X. Sun, G. Li, J. Phys. Chem.
C 112, 12242 (2008)
70. B.K. Das, T. Das, K. Parashar, S. Parashar, R. Kumar, A.
Anupama, B. Sahoo, Electron. Mater. Lett. 16, 255 (2020)
71. C. Leon, A. Rivera, A. Varez, J. Sanz, J. Santamaria, K.L.
Ngai, Phys. Rev. Lett. 86, 1279 (2001)
72. N. Ait-Younes, F. Millot, P. Gerdanian, Solid State Ion. 12,
431 (1984)
73. D.K. Lee, H.I. Yoo, Solid State Ion. 177, 1 (2006)
74. L. Essaleh, S. Amhil, S.M. Wasim, G. Marín, E. Choukri, L.
Hajji, Physica E Low Dimens. Syst. Nanostruct. 99, 37
(2018)
75. O. Ajili, B. Louati, K. Guidara, J. Mater. Sci. Mater. Electron.
29, 8649 (2018)
76. J. Mohanty, P. Behera, S.R. Mishra, T. Badapanda, S. Anwar,
IOP Conf. Ser. Mater. Sci. Eng. 178, 012014 (2017)
77. H. Zeghioud, N. Khellaf, A. Amrane, H. Djelal, W. Elfalleh,
A.A. Assadi, S. Rtimi, J. Photochem. Photobiol. A 346, 493
(2017)
78. Md.S.A.S. Shah, A.R. Park, K. Zhang, J.H. Park, P.J. Yoo,
ACS Appl. Mater. Interfaces 4, 3893 (2012)
Publisher’s Note Springer Nature remains neutral with
regard to jurisdictional claims in published maps and
institutional affiliations.

dhaou2021.pdf

Recommended

Recommended

More Related Content

Similar to dhaou2021.pdf

Similar to dhaou2021.pdf (20)

More from mohamedafqir1

More from mohamedafqir1 (6)

Recently uploaded

Recently uploaded (20)

dhaou2021.pdf