Influence of Cu-doping on the structural and optical properties of CaTiO3 powders

Article abstract of DOI:10.1016/j.materresbull.2016.04.024

Here, we report on the effect of chemical substitution on the structural and optical properties of Cu-doped CaTiO₃ (CTO) polycrystalline powders synthesized by the polymeric precursor method. Our findings are discussed based on the structural order-disorder effects originating from the modification of the Ca_1-xCu_xTiO₃ microcrystal matrix. These results may elucidate the compositional modulation and methods of controlling the structural design, as well as reveal the changes in the optical behavior of this system at an atomic level.

Full text of DOI:10.1016/j.materresbull.2016.04.024

Accepted Manuscript
Title: Influence of Cu-doping on the structural and optical
properties of CaTiO3 powders
Author: L.H. Oliveira A.P. de Moura F.A. LA Porta I.C.
Nogueira E.C. Aguiar T. Sequinel I.L.V. Rosa E. Longo J.A.
Varela
PII:
S0025-5408(16)30182-9
DOI:
Reference:
http://dx.doi.org/doi:10.1016/j.materresbull.2016.04.024
MRB 8760
To appear in:
MRB
Received date:
Revised date:
Accepted date:
20-9-2015
20-4-2016
21-4-2016
Please cite this article as: L.H.Oliveira, A.P.de Moura, F.A.LA Porta, I.C.Nogueira,
E.C.Aguiar, T.Sequinel, I.L.V.Rosa, E.Longo, J.A.Varela, Influence of Cu-doping on
the structural and optical properties of CaTiO3 powders, Materials Research Bulletin
http://dx.doi.org/10.1016/j.materresbull.2016.04.024
This is a PDF file of an unedited manuscript that has been accepted for publication.
As a service to our customers we are providing this early version of the manuscript.
The manuscript will undergo copyediting, typesetting, and review of the resulting proof
before it is published in its final form. Please note that during the production process
errors may be discovered which could affect the content, and all legal disclaimers that
apply to the journal pertain.

Influence of Cu-doping on the structural and optical properties of CaTiO3
powders
L. H. Oliveira,1 A. P. de Moura,1 F. A. La Porta,2* I. C. Nogueira,3 E. C. Aguiar,1 T. Sequinel,1 I. L. V. Rosa,3 E.
Longo1and J. A. Varela1
1 UNESP-Universidade Estadual Paulista, P.O. Box 355, 14801-907, Araraquara, SP, Brazil
2Department of Chemistry, Federal Technological University of Parana (UTFPR), PO Box 3131, 86036–370,
Londrina/PR, Brazil.
3DQ-UFSCar-Universidade Federal de São Carlos, P.O. Box 676, 13565-905, São Carlos, SP, Brazil
* Corresponding author:
felipe_laporta@yahoo.com.br; felipelaporta@utfpr.edu.br
Phone: +55 43 3315-9437

Graphical Abstract

Highlights


Ca1-xCuxTiO3 powders were successfully synthesized via a polymeric precursor
method.
Effects of Cu incorporated on the Ca-site into the CaTiO3 lattice as host matrix
has been investigated.

The optical behavior reveals that the Ca1-xCuxTiO3 powders have potential
applications in emerging technologies.

Here, we report on the effect of chemical substitution on the structural and optical properties of Cu-doped
CaTiO3 (CTO) polycrystalline powders synthesized by the polymeric precursor method. Our findings are
discussed based on the structural order-disorder effects originating from the modification of the Ca1-
xCuxTiO3 microcrystal matrix. These results may elucidate the compositional modulation and methods of
controlling the structural design, as well as reveal the changes in the optical behavior of this system at an
atomic level.

1. Introduction
Recently, complex perovskite oxides have been the subject of many studies in several fields because of
their desirable properties and potential applications [1-5]. Among these promising materials, which have a
general formula of ABO3, calcium titanate (CaTiO3 (CTO)) is a chemically stable, n-type semiconductor with a
wide band-gap of ~3.56 eV. CTO has excellent electrical and optical properties, which render it a potential
candidate for applications such as electronic devices, computer memory, and capacitors [6,7].
Along with the development of novel complex perovskite oxide materials, a large number of synthetic
strategies have recently been developed to produce CTO materials of different sizes and morphologies [8-15]. In
this context, it is well known that the polymeric precursor method is a simple, clean, and efficient approach to
fabricate complex functional materials. It also offers excellent chemical homogeneity and precise stoichiometric
control during heat treatments at relatively low temperatures [9,15,16]. Hence, this synthetic strategy
substantially reduces the presence of segregated phases and does not require long annealing periods compared
with other synthetic methods [9,15,16]. Moreover, it is well known that doping can enhance the physical and
chemical properties of these materials. Hence, these chemical substitutions, especially the partial replacement of
the Ca2+ or Ti4+ atoms, allow the emergence of technologically interesting properties for new applications.
From the most recent studies, up to now, the use of doped CTO materials have attracted large attention
and proven to be an effective means as strategy to obtain of materials with a giant dielectric constant, new
luminescent materials and so on [17-19]. As is well-known, the Cu2+ ions can influence dielectric permittivity as
well as its microstructure, modifying the potential barriers, mainly when segregated into grain boundary [20-22].
From recent studies, it has been observed that a small amounts of Cu2+ ions added to CTO material (e.g., such as
2 to 10 mol%) is responsible for increased its dielectric permittivity in 40 times. Although permittivity values
obtained for these systems being smaller than those are reported for CaCu3Ti4O12 (CCTO) perovskite, at higher
concentrations of Cu2+ ions, i.e. above 40% mol, a tangent of loss (tan) value of approximately 0.05 at 1kHz
was observed in the same frequency range [23]. However, it is crucial to elucidate the order-disorder effects and
structural changes that occur when the Ca2+ ions in the material are replaced with Cu2+ ions. This is a key factor
for the design of novel complex functional materials with tunable properties, and it is of great technological
importance for many applications.
Here, we report a detailed study on copper doping and its effects on the structural and optical properties
of Ca1-xCuxTiO3 powders synthesized by the polymeric precursor method. The obtained precursor was
investigated using thermogravimetric analysis (TGA). These obtained powders were analyzed using X-ray
diffraction (XRD) with Rietveld refinement, micro-Raman (MR) spectroscopy, field-emission scanning electron
microscopy (FE-SEM), Brunauer–Emmett–Teller (BET) nitrogen adsorption/desorption isotherms, as well as
ultraviolet–visible (UV–vis) and photoluminescence (PL) measurements. These analyses can elucidate the
structural changes, which is critical for controlling the optical behavior of the Ca1-xCuxTiO3 powders, being an
important strategy in order to revealing its potential applications.
2. Experimental section

The polymeric gel of Ca1-xCuxTiO3 with 0  x  0.10 moles of Cu2+ ions was synthesized by polymeric
precursor method. Initially, a 1.0 × 10-2 mol solution of titanium citrate was maintained at 60 °C. Subsequently,
(1- x) × 10-2 mol of Ca(NO3).4H2O was added to the solution. Following its complete dissolution, stoichiometric
quantities of 0.00, 0.02, 0.04, or 0.10 x 10-2 mol of Cu(NO3)2.2.5H2O were also added to the system. These
samples were denoted as CTO, x = 0.02, x = 0.04, and x = 0.10, respectively. Under constant stirring, the system
was maintained at 80 °C until its complete homogenization. Then, ethylene glycol was added to the citric acid
solution in a mass ratio of 60:40, respectively; thus producing a polymeric resin. Figure 1 shows the
TGA/differential thermal analysis (DTA) curve for thermal decomposition of the polymeric resin. The
TGA/DTA analysis was performed using NETZCH-STA-409 CELL equipment at a heating rate of 5 °C/min
under an air atmosphere.
Figure 1. TGA/DTA curves of amorphous powder obtained by the calcination of the polymeric gel.
Figure 1 presents the TGA/DTA curve for the thermal decomposition of the polymeric resin heat treated
from 25 to 1000oC. As can be seen, an exothermic peak was observed at 350 oC which is associated to a weight
loss of 6.88 %, corresponding to decomposition of the polymer chains involving the metallic cations producing
o
CO2 and H2O, which produced an amorphous powder (see Figure 2(a-d)). Above 600 C, none weight loss was
detected, indicating that a stable ceramic oxide is formed. According to this result, crystalline Ca1-xCuxTiO3
powders with 0  x  0.10 of Cu2+ ions are obtained after different stages of calcination. The first step, so-called
the pre-calcination, constitutes of the polymeric resin calcination to eliminate the larger amount of organic
material, citric acid and ethylene glycol, used in the synthesis. The second step, in particular consists of a heat
treating of these material for removal of residual organic materials as well as obtain a crystalline Ca1-xCuxTiO3
powders. In this work, the polymeric resin was heated in a conventional oven at 350 °C for 4 h using a heating
rate of 1 °C/min (i.e., the pre-calcination). This resulting material was manually de-agglomerated in the agate
mortar. Then, using a tubular oven under an O2 atmosphere, this material was heated at 800 °C for 2 h using a
heating rate of 5 °C/min (see Figure 2 (a-d)).

The Ca1-xCuxTiO3 powders with 0<x<0.10 mol of Cu2+ ions were characterized by XRD with Rietveld
refinement using a Rigaku-DMax/2500PC (Japan) with Cu-Kα radiation (λ = 1.5406 Å). The analysis was
conducted in the 2θ range from 10° to 130° with a scanning rate of 0.02°/min. The Rietveld refinement was
performed using the General Structure Analysis software (GSAS) program [24]. Rietveld refinement is a method
in which the position and intensity of peaks in a diffraction pattern obtained from step-scanning measurements
enables an approximate structural model of an actual structure to be produced [25]. Hence, refined parameters
for all the samples were obtained, such as the scale factor, background, shift lattice constants, profile of half-
width parameters (u, v, w), isotropic thermal parameters, strain anisotropy factor, occupancy, and atomic
functional positions. Furthermore, the background was corrected using a Chebyshev polynomial of the first kind.
The model described by Stephens [26] was used to account for the anisotropy in the half width of the reflections.
As implemented in the GSAS program, the profile peak function was modeled using a convolution of the
Thompson-Cox-Hastings pseudo-Voigt function [27] with the asymmetry function described by Finger et al.
[28].
Micro Raman (MR) spectra were obtained using the 488 nm line of a Nd:YAG laser, maintaining its
maximum output power at 25 mW to study the structural organization at short-range. The surface areas of the
Ca1-xCuxTiO3 powders were evaluated by the BET method using nitrogen adsorption/desorption analysis, which
was conducted on a Micromeritics ASAP 2010 surface area and porosity analyzer. The morphologies of the as-
prepared Ca1-xCuxTiO3 powders were thoroughly investigated using FE-SEM (JEOL, Model 7500F, JEOL Ltd.
Tokyo, Japan). The UV-vis absorption spectroscopy of the as-prepared materials was performed using Cary 5G
equipment, and the band-gap energy (Egap) values were calculated by the Kubelka-Munk method [29]. The PL
emission spectra were measured with a 139 Thermal Jarrel-Ash Monospec 27 monochromator and a Hamamatsu
R446 photomultiplier. A krypton ion laser (Coherent Innova) with an exciting wavelength of 350 nm was used,
maintaining a nominal output power of 200 mW. All the measurements were performed at room temperature.
3. Results and discussion
The crystalline structure of the as-prepared Ca1-xCuxTiO3 powders were evaluated by XRD with Rietveld
refinement analysis and MR spectroscopy. Figure 2 (a-d) shows the XRD patterns of the Ca1-xCuxTiO3 powders,
revealing a polycrystalline nature at long-range. An analysis of these results indicated that all the diffraction
peaks were perfectly indexed to the CTO orthorhombic structure with a space group of Pbnm, in accordance
with the Inorganic Crystal Structure Database (ICSD) card number 74-212. For x > 0.02 compositions, a CCTO
cubic secondary phase with a space group of Im-3 was identified in addition to the CTO phase. The secondary
phase was indexed according to ICSD card number 95-714. This suggests that the presence of the CCTO phase
was owed to the limited atomic diffusion of the Ca, Ti, and Cu atoms, which is promoted by the extended heat
treatments used to eliminate the organic material. Furthermore, these findings suggest that the amount of
precipitated CCTO phase increased as the Cu2+ ion concentration increased.

Figure 2. (a-d) XRD patterns of Ca1-xCuxTiO3 powders. Inset in (a-d) illustrates XRD patterns for the pre-
calcination step.
The Rietveld refinement measurements of the as-prepared Ca1-xCuxTiO3 powders were used to elucidate
the structural changes and to determine the percentage of the second phase, as shown in Figure 3 and Table 1. As
shown in Figure 3 (a-d), the Rietveld refinement results showed a good relationship between the observed XRD
patterns and the theoretical ones. The solubility limit was estimated as ~0.02 mol for the CTO orthorhombic
structure. In these analyses, the determined percentages of the secondary phase in x = 0.04 and x = 0.10 systems
were 4.34 wt.% and 9.80 wt.%, respectively, indicating that all the Cu atoms are crystallized in a CCTO
secondary phase.
In CTO unit cell, the Ca atoms are coordinated to twelve O atoms that form dodecahedral [CaO12]
clusters, while the Ti atoms are coordinated to six O atoms that form octahedral [TiO6] clusters. In this structure,
the adjacent [TiO6] clusters are displaced from its symmetric center. For Ca1-xCuxTiO3 powders, this effect can
be noted by the small variation of approximately 156 to 158° between [TiO6]−α−[TiO6] clusters. This behavior
can be attributed to the replacement by the Ca2+ ions to Cu2+ ions into A-site of the perovskite structure, i.e., due
to its polarization effect on [CaO12] e [TiO6] clusters in the host lattice, which is able to create distortions and
strains in Ca1-xCuxTiO3 powders and also have strong a influence on its optical properties [30]. However, it can
be noted that the lattice parameters and unit cell volumes obtained through the XRD refinement are very close to
those published in the literature for the CTO materials [30-32]. Moreover, changes in the crystallization, heating
rates, heat-treatment temperature, processing time, or crystal growth process can influence the structural
distortion degree of the octahedral TiO6 clusters within CTO matrix [33,34].

Figure 3. Rietveld refinements of the XRD patterns of as-prepared Ca1-xCuxTiO3 powders: (a) CTO, (b) x =
0.02, (c) x = 0.04, and (d) x = 0.10.
Most notably, there were no other significant changes in the lattice parameter values with the
introduction of Cu2+ ions into the A sites of the Ca1-xCuxTiO3 powders. This may be associated with the heat
treatment conditions used to obtain the Ca1-xCuxTiO3 powders. In the polymeric precursor method, the
amorphous powder has been subjected to a heat treatment in a conventional oven. During this process, the
energy is slowly transferred from the surface of the powder to its interior, promoting the growth of a CTO
crystalline structure with a greater degree of structural organization. Moreover, it should be noted that Ca2+ and
Cu2+ possess similar ionic radii (Cu2+ = 0.960 Å and Ca2+ = 0.969 Å). This corresponds with the lack of variation
in the lattice parameters and the cell volume of the copper-doped CTO, as shown in Table 1. Furthermore, the
peaks (212), (142) observed in the XRD patterns of CTO phase, also should be attribute to CCTO phase
simultaneously, and as a consequence of the increase in Cu composition is observed a large increase in the
relative intensity of these peaks, which is in good agreement with the Rietveld refinement results.

Table 1. Lattice parameters, unit cell volume, and correlation parameters obtained through the Rietveld
refinement method for the Ca1-xCuxTiO3 powders with 0  x  0.10 mol of Cu2+ ions prepared by polymeric
precursor method.
LatticeParameters
V (Å3)
α[Ti-O-Ti]
RBragg(%) Rwp(%) χ2
a (Å)
b (Å)
c (Å)
5.3864(8)
5.3863(9)
5.3868(16)
5.3861(12)
5.4374 (8)
5.4382(10)
5.4374(18)
5.4369(13)
7.6464(11)
7.6457(13)
7.6464(23)
7.6449(17)
223.9 (6) 156.3º
223.9(7) 157.2º
223.9(11) 157.8º
223.8(8) 157.9º
3.14
3.46
3.57
4.39
8.44
6.89
5.96
5.39
1.313
1.33
1.27
1.34
CTO
x= 0.02
x= 0.04
x= 0.10
For x ≥ 0.02 compositions, we observed the formation of a CCTO/CTO polycrystalline system [11]. In
the CCTO structure, the Cu atoms are coordinated with four oxygen atoms of the square-planar [CuO4] clusters.
Moreover, the distortion of the [TiO6] clusters in the CCTO matrix, owed to changes in the Cu-O bond length,

could increase the number of [TiO5VO ] [TiO6] clusters in the structure. The refinement results reveal a square
planar distorted octahedral Cu2+ complexes clusters, as a consequence of Jahn-Teller effects [35]. In this case,

x
the conductivity of CCTO would be owed to charge transference from [TiO5VO ] to [CuO4]´, [TiO5Vo ] to
[CuO4]x, or [CuO4]´ to [CuO4]x. Moreover, this process leads to the production of a Cu+ species ([CuO4]´) in the
lattice [36-41]. Thus, MR spectroscopy can be used to determine the level of structural defects or asymmetry–
symmetry in the Ca1-xCuxTiO3 systems.
Figure 4 shows the MR spectra of Ca1-xCuxTiO3 powders, confirming the structural order at short-range.
The structural behavior of the CTO powders resulted in nine Raman-active modes at 156, 180, 224, 246, 285,
335, 470, 491, and 631 cm-1, which are assigned to the orthorhombic structure of the CTO. In particular, the
peaks located at around 156, 180, 224, 246, 285, and 335 cm-1 are associated with the [O-Ti-O] bending
vibration modes, while the peaks at 470 and 491 cm-1 are assigned to distortions in the [TiO6] octahedral
clusters. The peak at 631 cm-1 is assigned to the [Ti-O-Ti] anti-symmetric stretching mode of the [TiO6]
octahedral clusters [9,10]. For the Ca1-xCuxTiO3 powders with x = 0.0 and x = 0.02, the horizontal arrows
indicate the broadening of the 631 cm-1 Raman-active mode as a function of the Cu2+ ion concentration in the
CTO lattice. In our previous work [30], significant changes in this Raman mode could be related to the
symmetry breaking of the [CaO12] clusters caused by the increase in Cu composition in the A-site of CTO host
matrix. This suggests the existence of an asymmetric structure due to the modification of the rotational modes in
the [TiO6] octahedral clusters (see Figure 4); specifically, a high electron density region (excess of electrons) and
a low electron density region (excess of holes) are formed.
In this case, the addition of the Cu2+ ions into the CTO matrix completely changes the MR spectra at
short-range. However, besides the Raman modes ascribed to the CTO orthorhombic phase for the x = 0.04 and x
= 0.10 systems, Raman-active modes were also verified at 447 cm-1(•) and 503cm-1(•), which were ascribed to
the rotation-like modes of the [TiO6] linkages related to the CCTO secondary phase [42-44]. Moreover, the
Raman mode at 503 cm-1 was only observed for the x = 0.10 composition. We consider that the addition of the
Cu atoms to the Ca1-xCuxTiO3 powders (x = 0.04 and x = 0.10) causes changes in the structural polarization of

the CCTO phase. This can be attributed to structural defects (oxygen vacancies and/or distortions) owed to the
different bond lengths of the Cu-O linkages [11,30,33]. This behavior can be indicative of a higher Ca2+/Cu2+
ratio, and thus depends on the spin orientation. This appears as the secondary phase in the Ca1-xCuxTiO3 samples,
which is owed to the structural rearrangement from a lower-symmetry CTO orthorhombic system to a higher-
symmetry CCTO cubic system. This confirms the formation of a polycrystalline system, and correlates with the
Rietveld data. Therefore, the phase transformation and cluster rotation induce depolarization via the replacement
of the Ca2+ ions by the Cu2+ ions.
Figure 4. MR spectra of as-prepared Ca1-xCuxTiO3 powders.
Figure 5 shows the FE-SEM micrographs of the as-prepared Ca1-xCuxTiO3 powders. The inset in Figure
5 shows the specific surface area (S) values of the Ca1-xCuxTiO3 powders. By analyzing the obtained results, as
shown in Figure 5, it can be determined that the Ca1-xCuxTiO3 powders prepared by the polymeric precursor
method consist of aggregated, irregular, and lamellar particles, which explains the low surface area values
obtained for these systems. As can be seen in Figure 5(a), there is little contact between the particles in the CTO
powder. In this sample, the particles are less reactive (surface area of 3.61m2/ g), and will not have sufficient
energy to initiate the sintering process (neck forming); thus, the particles are merely aggregated.
When the Cu2+ concentration is increased from x = 0 to 0.02 (see Figure 5 (b) and (c)), the bonding
between the particles is enhanced because the surface area of these particles is increased from 3.61 to 3.75 m2g-1,
respectively. Subsequently, during the heat treatment, contact between the particles will be maintained by van
der Waals forces. Sintering is initiated between the contacting particles (neck forming) and consequently large
aggregates will form [45]. Moreover, as the Cu2+ concentration is increased above 0.02, the particles become
more reactive, and larger agglomerates can be observed. In general, the FE-SEM micrographs reveal that the

addition of Cu2+ ions promotes the sintering process at lower temperatures, and also results in structural changes
and the relaxation of the material.
Figure 5. FE-SEM images of CaTiO3(a), x = 0.02 (b;c), x = 0.04 (d;e), and x = 0.10 (f;g) powders prepared by
the polymeric precursor method.
Figure 6 displays the UV-Vis spectra of the as-prepared Ca1-xCuxTiO3 powders. According to previous
studies [10,46], both the CTO and CCTO exhibit an optical absorption spectrum governed by direct electronic
transitions. For the Ca1-xCuxTiO3 powders, the UV-vis absorption measurements showed Egap values varying
between 3.51 and 2.10 eV, which is associated with specific structural disordering of the lattice [10-12].
Therefore, this behavior can be associated with the formation of acceptor and donor levels between the
conduction and valence bands that may have been introduced by the oxygen vacancies of the particles, which are
x
•
in three different charge states (i.e., VO , VO , and VO••). Two different band-gap values were determined for both
of the Ca1-xCuxTiO3 (x=0.04 and x=0.10) powders with excess Cu atoms (see Figure 6). These values are
associated with the CCTO cubic phase (2.45 eV and 2.05eV) and the CTO orthorhombic phase (3.2eV and
2.28eV). This behavior suggests that a low concentration of energy states exists within the valence and
conduction bands, which can be ascribed to the increased Cu content within the host matrix. This plays an
important role in the optical transitions of such systems. This stoichiometric change leads to a greater
contribution from the 3d orbital states of the Cu atoms in the CCTO secondary phase and a lower contribution
from the Ti 3d orbital states of the CTO phase [47,48]. So in principle, the cluster-like elucidation have played a
significant role on the understanding of PL performance is supported and strengthened by different extrinsic (at
surface) and intrinsic (at bulk) defect level distributions, which are responsible for tunable the optical behavior in
such systems. The defect structure and density variation surface and/or bulk might be responsible for the
different PL behavior of CCTO and CTO. Effective separation (electron/hole) requires the presence of a cluster-
to-cluster transfer process of electrons or holes from [CaO12]o/d or [CuO12]o/d or [TiO6]o/d, where o= ordered and

d= disordered. One way to boost PL efficiency is to exchange ordered complex clusters to disordered complex
clusters.
Figure 6. UV-vis absorbance spectra of as-prepared Ca1-xCuxTiO3 powders: (a) CTO, (b) x = 0.02, (c) x = 0.04,
and (d) x = 0.10.
PL measurements were performed to elucidate the influence of the various Cu atomic concentrations on
the optical properties of the Ca1-xCuxTiO3 powders. This is a very simple, convenient method for the analysis of
structural defects at medium-range [49-51]. Figure 7 shows the PL emission spectra of the as-prepared Ca1-
xCuxTiO3 powders. Several PL emission peaks can be observed for the Ca1-xCuxTiO3 powders, which exhibit
typical broad-band behavior. Specifically, the relaxations involve several energy states within the band gap,
which are in accordance with the optically measured gap. In particular, the individual contribution of each de-
convoluted PL profile was obtained using the Voigt function of the PL emission spectra, and analyzed using the
PeakFit program (version 4.05) [52]. The peak positions and corresponding areas, shown in Figure 7, were
included in the analysis.

Figure 7. Deconvolution of PL emission spectra of the as-prepared Ca1-xCuxTiO3 powders: (a) CTO, (b) x =
0.02, (c) x = 0.04, and (d) x = 0.10.
We compare the optical behavior of the as-prepared Ca1-xCuxTiO3 powders, which is a crucial step in
order to revealing its potential applications. The Ca1-xCuxTiO3 powders (x = 0.02 and 0.04) presented six
components; one component was situated at the violet region (390–418 nm), blue region (460–490 nm), orange
region (617 nm), and red region (670 nm), respectively, and two were situated at the green region (500–565 nm).
The CTO and x = 0.10 samples presented just two components, one at the blue region and one at the orange or
red region. As can be seen, CTO system presented a PL emission mainly centered at the blue region (54.05%),
which is associated to a higher concentration of deep defects. As the Cu2+ concentration was increased to x =
0.02, the area percentage of the orange component decreased to 1.9%. Moreover, it was possible to observed a
red component situated at 670 nm (2.2%). The x = 0.04 composition presented a decrease in the area percentage
of the blue (39%) and green (15%) components, while an increase in the area percentage of the orange (17%)
and red (7%) components was observed. This behavior is also accompanied by the displacement of the PL
emission to a higher wavelength. Finally, the PL emission of the x = 0.10 composition was situated in just the
blue region (460 nm). From these results, it can be determined that the increased Cu2+ concentration is
responsible for promoting a reduction in the structural polarization of the Ca1-xCuxTiO3 powders. This
polarization favors the elimination of deep defects and promotes a redistribution of the intermediary states in the
optical band gap. This fact suggests that the PL behavior of the Ca1-xCuxTiO3 powders is directly influenced by
the structural disorders that yield discrete levels in the forbidden band gap, which were confirmed by the UV-vis
measurements (see Figure 6). Therefore, the analysis of the PL data of the as-prepared Ca1-xCuxTiO3 powders

suggest that such materials are promising to fabrication of visible lighting lasers devices in the blue, green and
red regions of the electromagnetic spectrum, i.e., depending on the Cu composition in the host matrix.
In previous studies [30], using as a strategy of synthesis the microwave-assisted hydrothermal method,
in particular, the PL profile reported for the Cu-doped CTO powders with x = 0.02 is subtly different from the
one in this study. In general, the optical behavior is strongly dependent on the synthetic methods which provide a
different structural organization level for such materials, i.e., especially at the medium range. To further clarify
the PL mechanism based on structural changes induced by the dopants in the host matrix, the formation of
x
•
oxygen vacancies (i.e., VO , VO , and VO••) has a crucial role in the elucidation of this microscopically observed
behavior. The oxygen vacancies within the disordered structure may provide a possible explanation for the
changes observed in the PL emission of the as-prepared Ca1-xCuxTiO3 powders. Based on the structural order-
disorder effects, we used the Kröger–Vink notation to explain and represent the oxygen vacancy formation (i.e.,
x
•
VO , VO , and VO••) in the Ca1-xCuxTiO3 systems, which are summarized in the following equations:
x

[TiO6]x + [TiO5Vo ]   [TiO6]´+ [TiO5Vo ]
(1)
(2)


[TiO6]x + [TiO5Vo ]  [TiO6]´ + [TiO5Vo
]
´
Here, [TiO6 ] is a donor, [TiO5VO ] is a donor–acceptor, and [TiO5VO] is an acceptor. This model can be
extended for the lattice modifiers; in this case, the charge will transfer from [TiO6 ] to [CuO ]x , or
´
12
from[CaO VOx ] to [CaO ]'
.
12
11
[CuO12]x + [TiO6]´
 [CuO12]´+[TiO6]x
(3)
(4)
[CaO12]x + [CaO11Vo ]  [CaO12]´+ [CaO11Vo ]
x



[CaO12]x + [CaO11Vo ]  [CaO12]´+ [CaO11Vo
]
(5)
´
´
[CuO ]x [CaO ] [CuO ] [CaO ]x
(6)
12
12
12
12
Following the above logic, the following equations can summarize the charge transfer among the clusters for the
CCTO secondary phase:
[CuO4]x + [TiO6]´
 [CuO4]´+[TiO6]x
(7)
[CuO4]x + [CaO12]´  [CuO4]´ + [CaO12]x
(8)
e'
[CuO4]´  [CuO4]x where, CuII and CuI
(9)
In this case, the charge will transfer from [TiO6 ]´ to [CuO4 ]x or from [CaO VOx ] to [CaO ]' . An analysis
11
12
of the complex clusters based on the Kröger–Vink notation for these systems revealed that chemical substitution

promotes the formation of oxygen vacancies and can explain, in principle, the higher densification observed in
the CTO phase and hence are consistent with our previous reported [11], in which the analysis of the X-ray
absorption near edge structure (XANES) spectroscopy that reveals different coordination environments for these
materials.
The intentional structural changes generated by the compositional modulation of the system can be
reasonably understood by equations 1–9, which summarize the modification of the local matrix responsible for
the creation of new luminescent centers in such materials. For these systems, the charge transfer that occurs
during the formation of these species is strongly related to the structural distortions of the [TiO6]-α-[TiO6]
clusters and the formation of defects inside the band gap, as discussed earlier. Thus, it is highly likely that the PL
emission of the Ca1-xCuxTiO3 powders can be explained by the identification of these key factors. Moreover, the
proportion of the secondary phase also affects the PL behavior of the complex clusters in Ca1-xCuxTiO3 systems,
i.e., promotes a redistribution of the intermediary states inside the optical band gap. All these effects can
generate strain and distortion in the CTO and CCTO crystalline lattice by modifying the density of structural
defects. This promotes control of the PL behavior, and hence provides a rational method to analyze the effects of
order-disorder on the electronic structure of these systems.
4. Conclusions
This work reported on the optical properties and structural control of Ca1-xCuxTiO3 powders synthesized
via the polymeric precursor method. XRD with Rietveld refinements, MR, FE-SEM, BET, UV-vis, and PL
measurements were extensively employed to investigate the relationship between the structure and properties of
these complex materials based on Ca1-xCuxTiO3 systems. The XRD and Rietveld results suggest that a
CTO/CCTO polycrystalline system forms in powders with a Cu2+ ion concentration are greater than 0.02. For the
powders with Cu concentrations of x = 0.04 and 0.10, the MR spectroscopy also presented a different Raman-
active mode, which was ascribed to the rotation-like modes of the [TiO6] linkages related to the CCTO phase.
This can be ascribed to a structural rearrangement in the Ca1-xCuxTiO3 samples, from a lower-symmetry CTO
orthorhombic system to a higher-symmetry CCTO cubic system. These distortions (i.e., on short-, medium-, and
long-range length scales) in the structural arrangement of the Ca1-xCuxTiO3 system are responsible for the
rational control of their physical and chemical properties. The optical behavior of the Ca1-xCuxTiO3 powders are
directly influenced by structural disorders that yield discrete levels in the forbidden band gap, which suggests
that the samples have varying concentrations of deep and shallow defects. Our results demonstrate the influence
of the structural changes originating from the modification of the matrix (i.e., the effect of replacing Ca+2 with
Cu+2 in Ca1-xCuxTiO3 powders), which is crucial for the future development of devices that will use these
complex systems.
Acknowlegments
The authors appreciate the support of the following Brazilian research financing institutions CAPES, CNPq and
FAPESP (Proc. 2009/00367-6 and Proc. 2013/14647-6)

References
[1] J. Jumpatam, B. Putasaeng, T. Yamwong, P. Thongbai, S. Maensiri, J. Eur. Ceram. Soc., 34 (2014) 2941-
2950.
[2] J. Maul, A. Erba, I. M. Santos, J. R. Sambrano, R. Dovesi, J. Chem. Phys., 142 (2015) 014505.
[3] G. F. Teixeira, T. R. Wright, D. C. Manfroi, E. Longo, J. A. Varela, M. A. Zaghete, Mater. Lett., 139 (2015)
443-446.
[4] S. Thirumalairajan, K. Girija, V.R. Mastelaro, N. Ponpandian, ACS Appl. Mater. Interfaces, 6 (2014) 13917-
13927.
[5] Y. Yamada, Y. Kanemitsu, J. Lumin., 133 (2013) 30-34.
[6] T. Kimijima, K. Kanie, M. Nakaya, A. Muramatsu, Appl. Catal., B: Environmental, 144 (2014) 462-467.
[7] T. Kimijima, K. Kanie, M. Nakaya, A. Muramatsu, CrystEngComm, 16 (2014) 5591-5597.
[8] C. R. Foschini, R. Tararam, A. Z. Simões, M. Cilense, E. Longo, J. A. Varela, J. Alloys Compd., 574 (2013)
604-608.
[9] T. M. Mazzo, L. M. R. Oliveira, L. R. Macario, W. Avansi Jr, R. D. S. André, I. L. V. Rosa, J. A. Varela, E.
Longo, Mater. Chem. Phys., 145 (2014) 141-150.
[10] M. L. Moreira, E. C. Paris, G. S. Nascimento, V. M. Longo, J. R. Sambrano, V. R. Mastelaro, M. I. B.
Bernardi, J. Andrés, J. A. Varela, E. Longo, Acta Mater., 57 (2009) 5174-5185.
[11] L. H. Oliveira, E. C. Paris, W. Avansi, M. A. Ramirez, V. R. Mastelaro, E. Longo, J. A. Varela, X. M.
Chen, J. Am. Ceram. Soc., 96 (2013) 209-217.
[12] R. Parra, R. Savu, L.A. Ramajo, M. A. Ponce, J. A. Varela, M. S. Castro, P. R. Bueno, E. Joanni, J. Solid
State Chem., 183 (2010) 1209-1214.
[13] L. Ramajo, R. Parra, J. A. Varela, M. M. Reboredo, M. A. Ramírez, M. S. Castro, J. Alloys Compd., 497
(2010) 349-353.
[14] T. Sequinel, I. G. Garcia, S. M. Tebcherani, E. T. Kubaski, L. H. Oliveira, M. S. Li, E. Longo, J.A. Varela,
J. Alloys Compd., 583 (2014) 488-491.
[15] A. O. Turky, M. M. Rashad, Z. I. Zaki, I. A. Ibrahim, M. Bechelany, RSC Advances, 5 (2015) 18767-
18772.
[16] M. M. Nakata, T. M. Mazzo, G. P. Casali, F. A. La Porta, E. Longo, Chem. Phys. Lett., 622 (2015) 9-14.
[17] F. Moura, A. C. Cabral, L. S. R. Rocha, E. C. Aguiar, A. Z. Simões, E. Longo, Ceram. Int. 42 (2016) 4837-
4844.
[18] T. Wanjun, C. Donghua, Mater. Res. Bull, 44 (2009) 836-839.
[19] J. Jumpatam, B. Putasaeng, T. Yamwong, P. Thongbai, S. Maensiri, Mater. Res. Bull, 77 (2016) 178–184.
[20] C. K. Yeoh, M. F. Ahmad, Z. A. Ahmad, J. Alloys Compd. 443 (2007) 155-160.
[21] T. T. Fang, L. T. Mei, J. Am. Ceram. Soc. 90 (2007) 638-640.
[22] B. Shri Prakash, K. B. R. Varma, J. Mater. Sci. 42 (2007) 7467-7477.
[23] B. Cheng, Y.-H. Lin, H. Yang, J. Lan, C.-W. Nan, X. Xiao, J. He, J. Am. Ceram. Soc. 92 (2009) 2776-
2779.
[24] A. C. Larson, R. B. V. Dreele, in: L.A.N. Laboratory (Ed.), Los Alamos National Laboratory, Los Alamos,
NM, 2001, pp. 230.
[25] H. Rietveld, J. Appl. Crystallogr., 2 (1969) 65-71.

[26] P. Stephens, J. Appl. Crystallogr., 32 (1999) 281-289.
[27] P. Thompson, D. E. Cox, J. B. Hastings, J. Appl. Crystallogr., 20 (1987) 79-83.
[28] L. W. Finger, D. E. Cox, A. P. Jephcoat, J. Appl. Crystallogr., 27 (1994) 892-900.
[29] P. Kubelka, F. Munk, Aussig, Für. Technol. Phys., 12 (1931) 593-601.
[30] L. H. Oliveira, A. P. de Moura, T. M. Mazzo, M. A. Ramírez, L. S. Cavalcante, S. G. Antonio, W. Avansi,
V. R. Mastelaro, E. Longo, J. A. Varela, Mater. Chem. Phys. 136 (2012) 130-139.
[31] L. S. Cavalcante, V. S. Marques, J. C. Sczancoski, M. T. Escote, M. R. Joya, J. A. Varela, M. R. M. C.
Santos, P. S. Pizani, E. Longo, Chem. Eng. J. 143 (2008) 299-307.
[32] R. F. Gonçalves, A. R. F. Lima, M. J. Godinho, A. P. Moura, J. Espinosa, E. Longo, A. P. A. Marques,
Ceram. Int. 41 (2015) 12841-12848.
[33] V. S. Marques, L. S. Cavalcante, J. C. Sczancoski, D. P. Volanti, J. W. M. Espinosa, M. R. Joya, M. R. M.
C. Santos, P. S. Pizani, J. A. Varela, E. Longo, Solid State Sci., 10 (2008) 1056-1061.
[34] E. S. Junior, F. A. La Porta, M. S. Li, J. Andrés, J. A. Varela, E. Longo, Dalton Trans. 44 (2015) 3159-
3175.
[35] J. H. Clark, M. S. Dyer, R. G. Palgrave, C. P. Ireland, J. R. Darwent, J. B. Claridge, M. J. Rosseinsky, J.
Am. Chem. Soc., 133 (2011) 1016-1032.
[36] D. Capsoni, M. Bini, V. Massarotti, G. Chiodelli, M. C. Mozzatic, C. B. Azzoni, J. Solid State Chem., 177
(2004) 4494-4500.
[37] B. Cheng, Y.-H. Lin, H. Yang, J. Lan, C.-W. Nan, X. Xiao, J. He, J. Am. Ceram. Soc., 92 (2009) 2776-
2779.
[38] O. F. Schirmer, Ferroelectrics, 303 (2004) 131-135.
[39] M. H. Whangbo, M. A. Subramanian, Chem. Mater., 18 (2006) 3257-3260.
[40] L. Wu, Y. Zhu, S. Park, S. Shapiro, G. Shirane, J. Tafto, Phys. Rev. B, 71 (2005) 014118.
[41] A. F. L. Almeida, R. S. D. Oliveira, J. C. Goes, J. M. Sasaki, A. G. Souza Filho, J. Mendes Filho, A. S. B.
Sombra, Mater. Sci. Eng., B, 96 (2002) 275 − 283.
[42] N. Kolev, R. P. Bontchev, A. J. Jacobson, V. N. Popov, V. G. Hadjiev, A. P. Litvinchuk, M. N. Iliev, Phys.
Rev. B, 66 (2002) 132102.
[43] D. Valim, A. G. Souza Filho, P. T. C. Freire, S. B. Fagan, A. P. Ayala, J. Mendes Filho, A.F.L. Almeida, P.
B. A. Fechine, A. S. B. Sombra, J. Staun Olsen, L. Gerward, Phys. Rev. B, 70 (2004) 132103.
[44] R. M. German, Fundamentals of Sintering, in: S.J. Schneider (Ed.) Engineered Materials Handbook:
Ceramics and glasses, ASM International, 1991.
[45] R. F. Gonçalves, M. J. Godinho, E. R. Leite, A. P. Maciel, E. Longo, J. A. Varela, J. Mater. Sci., 42 (2007)
2222-2225.
[46] S. B. Fagan, A. G. Souza Filho, A. P. Ayala, J. M. Filho, Phys. Rev. B, 72 (2005) 014106.
[47] M. Matos, L. Walmsley, J. Phys.: Condens. Matter, 18 (2006) 1793.
[48] J.-C. Zheng, A. I. Frenkel, L. Wu, J. Hanson, W. Ku, E. S. Božin, S. J. L. Billinge, Y. Zhu, Phys. Rev. B, 81
(2010) 144203.
[49] F. M. C. Batista, F. A. La Porta, L. Gracia, E. Cerdeiras, L. Mestres, M. S. Li, N. C. Batista, J. Andrés, E.
Longo, L. S. Cavalcante, J. Mol. Struct., 1081 (2015) 381-388.
[50] L. S. Cavalcante, V. M. Longo, J.C. Sczancoski, M. A. P. Almeida, A. A. Batista, J. A. Varela, M. O.
Orlandi, E. Longo, M. S. Li, CrystEngComm, 14 (2012) 853-868.

[51] F. A. La Porta, M. M. Ferrer, Y. V. B. de Santana, C. W. Raubach, V. M. Longo, J. R. Sambrano, E. Longo,
J. Andrés, M. S. Li, J. A. Varela, J. Alloys Compd., 556 (2013) 153-159.
[52] PeakFit, version 4.05, S.S. Inc, Copyright © 2015.