Structural and electrical conduction behaviour of yttrium doped strontium titanate: anode material for SOFC application

Article abstract of DOI:10.1016/j.jallcom.2018.03.170

Yttrium (Y³⁺) doped strontium titanate (Y_xSr_1-xTiO_3-δ with x = 0, 0.03, 0.05, 0.08 and 0.10) was synthesized via citrate-nitrate auto-combustion route. The electrical and structural properties of the prepared compositions were investigated to check their suitability as anode material for SOFC application. The XRD patterns of synthesized specimens showed single phase having a crystal structure similar to that of the undoped SrTiO₃. Trace amounts of secondary phase were observed in compositions with x ≥ 0.05. The FESEM micrographs of all compositions revealed a non-monotonous behaviour of grain size in Y_xSr_1-xTiO₃ system with YST8 as most porous sample among all prepared compositions. The oxygen content in the specimens was also confirmed through XPS analysis. The conductivity values in air and H₂ atmosphere were obtained from Nyquist plots in the temperature range of 100–700 °C. From a systematic study, the conduction mechanism was explained on the basis of defect chemistry. Moreover, the chemical stability of samples was checked through XRD patterns with corresponding FESEM micrographs after the conductivity measurements in H₂ atmosphere. Results showed that amongst all the prepared samples, Y_0.08Sr_0.92TiO₃ may be proposed as a promising composition as anode material for SOFC application.

Full text of DOI:10.1016/j.jallcom.2018.03.170

Accepted Manuscript
Structural and electrical conduction behaviour of yttrium doped strontium titanate:
Anode material for SOFC application
Saurabh Singh, Priyanka A. Jha, Sabrina Presto, Massimo Viviani, A.S.K. Sinha, Salil
Varma, Prabhakar Singh
PII:
S0925-8388(18)31038-7
10.1016/j.jallcom.2018.03.170
JALCOM 45400
DOI:
Reference:
To appear in: Journal of Alloys and Compounds
Received Date: 10 January 2018
Revised Date: 12 March 2018
Accepted Date: 13 March 2018
Please cite this article as: S. Singh, P.A. Jha, S. Presto, M. Viviani, A.S.K. Sinha, S. Varma, P. Singh,
Structural and electrical conduction behaviour of yttrium doped strontium titanate: Anode material for
SOFC application, Journal of Alloys and Compounds (2018), doi: 10.1016/j.jallcom.2018.03.170.
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ACCEPTED MANUSCRIPT
Structural and electrical conduction behaviour of yttrium
doped strontium titanate: anode material for SOFC
application
Saurabh Singha, Priyanka A. Jhaa, Sabrina Prestob, Massimo Vivianib, A. S. K. Sinhac,
Salil Varmad and Prabhakar Singha*
aDepartment of Physics, Indian Institute of Technology (Banaras Hindu University) Varanasi-221005,
India
bCNR-ICMATE, c/o DICCA-UNIGE, Via all’Opera Pia 15, 16145, Genova, Italy
cDepartment of Chemical Engineering and Technology, Indian Institute of Technology (Banaras Hindu
University), Varanasi-221005, India
dChemistry Division, Mod. Labs., Bhabha Atomic Research Centre, Trombay, Mumbai, 400085, India
*E-mail: psingh.app@iitbhu.ac.in
Abstract
Yttrium (Y3+) doped strontium titanate (YxSr1-xTiO3-δ with x = 0, 0.03, 0.05,
0.08 and 0.10) was synthesized via citrate-nitrate auto-combustion route. The electrical
and structural properties of the prepared compositions were investigated to check their
suitability as anode material for SOFC application. The XRD patterns of synthesized
specimens showed single phase having a crystal structure similar to that of the undoped
SrTiO3. Trace amounts of secondary phase were observed in compositions with x ≥
0.05. The FESEM micrographs of all compositions revealed a non-monotonous
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behaviour of grain size in YxSr1-xTiO3 system with YST8 as most porous sample among
all prepared compositions. The oxygen content in the specimens was also confirmed
through XPS analysis. The conductivity values in air and H2 atmosphere were obtained
from Nyquist plots in the temperature range of 100 to 700 oC. From a systematic study,
the conduction mechanism was explained on the basis of defect chemistry. Moreover,
the chemical stability of samples was checked through XRD patterns after the
conductivity measurements in H2 atmosphere. Results showed that amongst all the
prepared samples, Y0.08Sr0.92TiO3 may be proposed as a promising composition as
anode material for SOFC application.
Keywords: Perovskite; Defects; Electrical conductivity; SOFC
1. Introduction
Fuel cells, introduced to resolve the issues of environmental pollution and
energy crisis are receiving more attention in last two decades to make them efficient,
cost-effective and technologically viable clean energy source. Solid oxide fuel cell
(SOFC) is one of the special type of fuel cell that converts chemical energy directly into
electrical energy without involving the process of combustion [1, 2]. These properties
have made SOFC as an attractive emerging technology. The wide range of doped
SrTiO3 has been found as anode materials for SOFC [3-6]. In literature, several attempts
have been made to use Y-doped SrTiO3 with A-site deficient stoichiometry as anode in
a solid oxide fuel cell for e.g. Y0.07Sr0.895TiO3 has been utilized as anode in an
electrolyte-supported fuel cell [7]. Further, Y0.08Sr0.88TiO3 was used to impregnate a
backbone of YSZ electrolyte supported SOFC [8] and YST-SDC composite
impregnated by Ni to act as a catalyst [9].
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Perovskite-type oxides (ABO3) are promising candidates as anode for SOFCs
and sensor as well due to their mixed ionic/electronic (electrical) conductivity with good
catalytic activity [10-12]. Pure SrTiO3 perovskite cannot be used as anode due to its low
electrical conductivity, whereas donor doped SrTiO3 systems are proposed to be
suitable to anodic conditions due to their chemical and thermal stability, high electrical
conductivity and carbon/sulfur tolerance [13]. It is well known that in SrTiO3 based
systems, the oxidation of fuel occurs at triple phase boundary region and mixed
conductivity feature provides oxidizing reaction sites that decreases the polarization and
accelerate the anode reaction [6]. However, the combination of ionic and electronic
conductivity could not satisfy the requirements for the electrode materials, therefore,
many efforts have been made to enhance the conductivity of the various materials. It is
reported that the acceptor doping in the perovskite materials enhances the ionic
conductivity [14, 15] whereas, the donor doping increases the electronic conductivity
[6]. Yttrium acts as donor dopant for the SrTiO3, and suitable for substitution due to less
difference between ionic radii of Y3+ (1.22 Å, CN 12) and Sr2+ (1.44 Å, CN 12). As the
results of doping, two phenomena occur simultaneously, first maintain electro-neutrality
in lattice defect structure due to charge imbalance between Y3+ & Sr2+ and second
improve the electrical conductivity of SrTiO3 via the formation of oxygen-rich planes.
In addition to that, redox coupling between Ti4+ and Ti3+ that also occurs in the system
contributes to improving the electrical conductivity under reducing atmosphere and
shows an n-type semiconducting behaviour [16-19].
In the present work, a systematic investigation on Y3+ doped SrTiO3 (YxSr1-
xTiO3-δ with x = 0, 0.03, 0.05, 0.08 and 0.10) was performed and results obtained were
analysed. Samples were synthesized via citrate-nitrate auto-combustion route using
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citric acid as fuel agent. The effect of Yttrium (Y3+) doping on electrical conductivity in
SrTiO3 under air and H2 atmospheres was investigated through electrochemical
impedance measurements.
2. Experimental
2.1 Sample preparation
The samples of the system YxSr1-xTiO3-δ with x = 0, 0.03, 0.05, 0.08 and 0.10
abbreviated as YST0, YST3, YST5, YST8 and YST10, respectively were synthesized
via the citrate-nitrate auto-combustion route. The starting materials, Y2O3 (99.9%), SrO
(99.5%), C12H28O4Ti (97%) and citric acid were taken in appropriate amounts. The
weighed Y2O3 and SrO were converted individually into Yttrium nitrate and Strontium
nitrate respectively with the help of dilute nitric acid. Thereafter, all nitrates were
dissolved in distilled water individually and mixed together to form a clear and
homogeneous solution. The weighed citric acid was also added to the solution. Another
solution of titanium isopropoxide (C12H28O4Ti) in ethylene glycol was mixed at 100 ºC
to form a homogeneous solution. Then all the solutions were dissolved in distilled water
with a few drops of ammonium solution to maintain the pH value of 2. In the mixed
solution, a white precipitate of titanium hydroxide was formed which was dissolved by
adding a few drops of nitric acid under pH value of 2. The transparent solution was
o
slowly evaporated on a hot plate at temperature 250 C under continuous stirring via
magnetic stirrer. After a few hours, it is converted into a light yellow coloured gel
which upon further heating converted into light brownish colour ash. The as-prepared
powder was ground with mortar and pestle. To remove the organic binders and nitrates,
it was calcined in air at 1000 oC for 10 h in an alumina crucible. The calcined powders
were pelletized in 12 mm disk by a hydraulic press of 5 tons. Thereafter, pellets were
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sintered in air at 1200 oC for 12 h with a heating rate of 5 ºC/min and then cooled down
to room temperature.
2.2 Characterization techniques
X-ray diffraction (XRD, Rigaku Miniflex II desktop) with Cu- Kα radiation (λ =
1.54098 Å) in 2θ range of 20o to 70o with a step size of ∆2θ = 0.02º was used to
characterize the phase structure of as sintered and post reduced YxSr1-xTiO3 samples
(x=0.00, 0.03, 0.05, 0.08 and 0.10) at room temperature. The relative density of each
sample was measured by Archimedes principle employing Denver SI-234 density
measurement kit. The microstructure of all fractured samples was studied by field
emission scanning electron microscope (FESEM, NOVA NANOSEM 450). Average
grain sizes of Y doped SrTiO3 samples were statistically calculated by ImageJ software.
The XPS spectra of all compositions were measured by a high performance analytical
instrument (KRATOS, Amicus) using Mg target under 1.0 × 10-6 Pa pressure. The XPS
peaks were analysed and fitted by using XPSPEAK 4.1 software. For electrical
measurements, the sintered pellets were polished and coated with Ag-paste on both
o
sides. The resultant coated samples were cured at 700 C for 20 minutes. Electrical
conductivity was measured in air between room temperature and 700 oC. Samples were
then reduced at 700 oC for 24 h under pure flowing hydrogen and electrical conductivity
was measured again down to room temperature. A four-probe test (Probostat, Norecs)
was employed and the impedance was measured by Frequency Response Analyzer and
Potentiostat (SOLARTRON 1250+1286, Schlumberger), in the range of frequency
between 1 Hz to 104 Hz.
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3. Results and discussion
3.1 Structural studies
The XRD patterns of the sintered samples of the system YxSr1-xTiO3 (i.e. YST0, YST3,
YST5, YST8 and YST10) are shown in Fig. 1 (a). To check the symmetry, Rietveld
refinement was carried out using FULL PROF Suite software package with Pseudo-
Voigt wave function. From the best fit, χ2 values are mentioned in Table 1. It is
ꢀ
observed that all the samples show cubic phase with space group Pm3m (JCPDS Card
Number: 86-0178). In the samples with x ≤ 0.03 no impurity peaks are detected, while
for x ≥ 0.05, secondary phase Y2Ti2O7 (JCPDS Card Number: 87-1849) is observed. It
is also observed that the diffraction peaks shift towards higher angle as the
concentration of Yttrium is increased as shown in Fig. 1 (b) which clearly indicates that
the lattice parameters are decreasing continuously. There may be two factors
responsible for decrease in lattice parameters. First, incorporation of Y3+ on Sr sites that
leads to decrease lattice parameters due to smaller ionic radii of Y3+ (1.22 Å, CN 12) as
compared to Sr2+ (1.44 Å, CN 12). Secondly, increasing amount of secondary phase
Y2Ti2O7 which may increase the defects such as oxygen vacancies, lattice disorders etc.
in the samples with x > 0.03 and which are removed on sintering leading to lattice
contraction [20]. Similar kind of decreasing behaviour in lattice parameter has been
reported by other authors [21].
However, more complex incorporation mechanisms might take place (see
below), with the same effect on lattice constant.
The Reference Intensity Ratio (RIR) method was used to calculate the concentration of
secondary phase [19]. RIR relates the concentration of phases (Xi) in a sample with the
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intensity of the highest peak (Ii) of each phase, obtained by XRD through tabulated RIRi
coefficient defined as:
⁄
ꢁꢂꢁꢃ = ꢁꢃ = ꢂꢃ ꢂꢄ
(1)
where Ic is the corundum intensity in 50:50 mixture with phase i. Therefore, the
concentration of secondary phase, Xs , can be calculated by equation
⁄
⁄
⁄
ꢅꢆ ꢅꢇ = ꢂꢆ ꢂꢇ ∙ ꢁꢇ ꢁꢆ ,
(2)
where p and s are referred, respectively to main and secondary phase and Xp is the
concentration of the main phase. In this case, Rp and Rs are equal to 7.32 and 7.89,
respectively.
Xs values, as function of x ≥ 0.03 in YxSr1-xTiO3-δ are reported in Fig. 2 (a). It is
observed that Xs increases by increasing amount of yttrium doping.
The lattice parameters of the studied systems are also plotted as a function of tolerance
factor (Fig. 2 (b)). From Table 1, it is observed that the unit cell volume decreases with
increase in Y content, which may be attributed to the replacement of the Sr cation by
incorporation of Y3+ within the perovskite structure [22].
Fig. 2 (c) shows the variation of intensity of XRD peak with 2θ-δ (o) where δ is
the angle at which maxima occurs showing diffuseness of XRD peak (110) with the
compositions and observed that the diffuseness of XRD peak increases towards left with
the Y3+ substitution in SrTiO3. The increase in diffuseness is due to the formation of
polarons in the matrix [23]. Microstrain and crystallite size of the samples were
calculated from XRD results by using Williamson-Hall model given by [24]:
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βcosθ = 0.9 ꢈꢉ + 4ꢊꢋꢌꢍꢎ
(3)
where t is average crystallite size,
ε is microstrain, β is full width half maxima
(FWHM) at Bragg’s angle (2θ) and λ being the X-ray wavelength of CuKα (λ = 1.54098
Å). The slope of βcosθ vs 4sinθ curve (graphs not shown here) gives the value of
microstrain and intercept gives the value of crystallite size. Fig. 2 (d) shows the
variation of microstrain and crystallite size with the compositions. It is observed that
value of microstrain and crystallite size showed the similar trend i.e. they are maximum
for YST5, whereas, the microstrain and crystallite size decrease for YST3 and YST8.
FESEM images in Fig. 3, reveal that grains and grain boundary are clearly
formed in YST5, whereas, its formation in other compositions are relatively weak.
Maximum microstrain and crystallite size in YST5 could be due to the enhanced grain-
to-grain interactions as shown in FESEM images. It may be caused by the generation of
high amount of stress concentration and local distortions under high sintering
temperature [25].
3.2 Microstructural analysis
o
FESEM images of fractured samples, sintered at 1200 C in air, are shown in
Fig. 3 (a)-(e). YST8 sample appears the most porous among all compositions, and that
finding was confirmed with the value of density measured by Archimedes method as
reported in Table 1. The average grain sizes of compositions YST0, YST3, YST5,
YST8 and YST10 were found to be 2.19 µm, 1.15 µm, 3.67 µm, 1.55 µm, and 3.87 µm,
respectively. Along with the formation of dense grains, water-mark and lamellar
features are also observed in dense micrographs as illustrated in Fig. 3 (c and e). This
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non-monotonous behavior of grain size and density is coherent with the crystallite size
and microstrain behavior, reported in Fig. 2 (d).
3.3 XPS analysis
The oxidation states of constituent elements are analysed using X-ray
Photoelectron Spectroscopy (XPS) technique. Fig. 4 (a) shows the XPS wide spectra of
Y doped SrTiO3 samples. The entire XPS spectrum consists of core and satellite binding
energy peaks of constituent elements including C. The binding energy (BE) peaks of the
constituents are marked for different doping concentrations Y (0.0 ≤ x ≤ 0.10). The C 1s
peak at ∼286 eV originates from external contamination or surface-absorbed CO2. The
binding energies were further calculated after calibration with respect to the C1s
reference peak. The spectra revealed binding energy peaks of Sr 3d at ∼134 eV, Y 3d
at∼157 eV, Sr 3p at ∼ 269.4 eV, Ti 2p at ∼ 459.4 eV, and O 1s at ∼ 530.6 eV.
Accordingly, the Sr2+ state is confirmed by the above mentioned peak positions. With
the increase in Y3+ doping content, BE values shifted to higher energy. In Fig. 4 (a), a
satellite peak of O (KLL) is also observed at ∼741.3 eV [26]. Also, O 1s spectral
regions are used to obtain the information regarding the presence of oxygen vacancies
in the samples.
The O-1s peaks, observed for all the compositions in the spectrum, are broader
and asymmetric in nature. As the peaks are broad in nature, therefore, they may be split
into two peaks. These peaks can be attributed to two different types of oxygen species:
lattice oxygen and adsorbed oxygen [27]. Fig. 4 (b) reveals that the sample YST5 has a
higher percentage of adsorbed oxygen in comparison to the lattice oxygen as illustrated
in Table 1. Whereas, in all other investigated samples, the percentage of lattice oxygen
is more than the adsorbed oxygen.
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The oxidation states of all constituent atoms of the samples have been
investigated by XPS spectra. The XPS spectra for Ti 2p of YST5 sample is shown in
Fig. 4 (c). The binding energies of Ti 2p3/2 and Ti 2p1/2 for the YST5 sample are 458.3
and 464.2 eV respectively as shown in Fig. 4 (c). These ions exhibit n-type
semiconducting behaviour due to redox coupling of Ti4+/Ti3+ and increase the electrical
conductivity as reported by Fuentes et. al 2015 [28].
3.4 Impedance analysis
From spectra obtained by impedance spectroscopy, in both air and hydrogen
atmosphere, the total resistance (Rt) was obtained as the intercept of depressed semi-
circular arc on the real axis at low frequency. The total conductivity of the samples was
calculated by applying the formula:
⁄
⁄
ꢏ = 1 ꢁꢉ ∙ ꢐ ꢑ
(4)
where σ is the conductivity, l is the thickness of pellet and S is the surface area of the
pellet. To analyse the mechanism of conductivity, activation energy (Ea) was calculated
for each regime in air as well as in reducing atmosphere by using the relation:
⁄
⁄
ꢏ = ꢏꢒ ꢓ ∙ exp(− ꢀꢁ ꢂꢃ)
(5)
where σo is a pre-exponential factor, Ea is the activation energy for conduction, k is the
Boltzmann constant and T is the absolute temperature.
In Fig. 5 (a and b), the Arrhenius plots of log σT vs. 1000/T are reported for all
compositions in air and hydrogen atmosphere, respectively.
It is observed that the conductivity of each composition increases in hydrogen in
comparison to that of the air. YST5 showed the highest conductivity in air whereas
YST8 in H2 atmosphere. The total conductivity of all reduced samples mostly reflects
the electronic conductivity because the oxygen vacancies have the lower mobility as
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compared to electrons in the Y-doped SrTiO3 system. However, the Arrhenius plots, i.e.
graph between log (σ.T) and 1000/T exhibit the two slopes in the two temperature
ranges (Fig. 5 (a) and (b)). At the lower temperature range i.e. 200 to 400 ºC, the linear
trend is observed implying the n-type semiconductor like conduction behaviour, while
change of slope at higher temperature range i.e. 400 to 700 ºC indicates the polaron
hopping conduction mechanism along with oxygen ion conductivity [18, 23].
The activation energy calculated for all compositions is listed in Table 1. It can
be noted that in air, Ea values are low for samples YST5 and YST10, suggesting
electronic nature of conductivity, especially in the high temperature range.
The activation energy is further lowered after reduction in H2 atmosphere. For
all the samples, the activation energy has been found to be in the range of 0.1 to 0.2 eV
except for the YST3 sample where Ea exceeds 0.5 eV.
Above reported experimental observations can be explained by considering the
incorporation mechanisms of dopant into the perovskite structure of SrTiO3. The
average value of ionic radius for Y3+ in 6-fold coordination (0.90 Å) and in 12-fold
coordination (1.22 Å) [29], three types of incorporation can be considered and
expressed by the following equations written in the Kröger-Vink notation:
ꢉꢊꢋꢌꢍ
ꢎ
ꢄ ꢆꢇ + 2ꢃꢈꢆꢅ ꢏꢐꢐꢐꢑ 2ꢄ∙ + 2ꢃꢈꢒ + 6ꢆꢓ + ꢅꢔ ꢆꢅ
(6a)
ꢅ
ꢉꢊ
ꢋꢌ
ꢍ
ꢉꢊꢋꢌꢍ
ꢎ
ꢄ ꢆꢇ ꢏꢐꢐꢐꢑ ꢄ∙ + ꢄꢒ + 3ꢆꢓ
(6b)
(6c)
ꢅ
ꢉꢊ
ꢋꢌ
ꢍ
ꢘꢙꢚꢃꢛ
ꢜ
ꢔ ꢖꢗ ꢝꢞꢞꢞꢟ 2ꢔꢠ + ꢡꢛ∙∙ + 3ꢖꢛꢢ
ꢕ
ꢚꢃ
Eq. (6a) accounts for the incorporation of Y3+ at the Sr-site, with formation of positively
charged defects, compensated by electrons or, equivalently, by partial reduction of Ti4+
to Ti3+.
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In the second case, Eq. (6b), the incorporation takes place both at Sr- and Ti-sites, with
self-compensation of charges. Finally, Eq. (6c) is relevant to the case of incorporation at
the Ti-site with formation of negatively charged defects, compensated by oxygen
vacancies.
As already reported in case of BaTiO3 [30], the most energetically favourable
incorporation mechanism for Y3+ and similar rare-earths like Dy, Er and Gd, is the self-
compensation.
However, other factors can also determine the actual incorporation mechanism,
the most important being the amount of dopant, the cationic ratio and the temperature.
Lattice parameters and strain are determined by the incorporation mechanism so that,
minimal local distortion (microstrain) can be expected when Y3+ is replacing both
(larger) Sr2+ and (smaller) Ti4+ atoms. The formation of the pyrochlore phase also
introduces point defects, namely Ti vacancies, which affect the lattice size and strain
state. Therefore, results indicate that for YST3 self-compensation (Eq. 6a) is mainly
active, causing limited microstrain and negligible amount of electrons leading to
conduction of ionic nature due to oxygen vacancies produced either intrinsically
(Schottky disorder) or through the mechanism of Eq. (6c).
Similar considerations hold for YST8 sample, where the small microstrain indicates
simultaneous substitution at both cationic sites, which also reflect lower conductivity
and high Ea in air. In this sample, the presence of a secondary phase i.e. Ti vacancies in
the main phase may account for observed lowering of the unit cell volume.
Mechanism of Eq. (6a) is mainly active in samples YST5 and YST10 as demonstrated
by higher conductivity, lower Ea and higher microstrain. XPS analyses are in agreement
as a large amount of Ti3+ was detected on the surface of YST5 sample.
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Under reducing atmosphere, lattice oxygen may be lost and thus oxygen vacancies will
be generated as per the following equation
2ꢃꢈꢆꢅ ꢏꢐꢐꢑ 2ꢃꢈꢋꢒ ꢌ + ꢕꢍ∙∙ + 3ꢆꢍꢓ + ꢅꢔ ꢆꢅ
(7)
in order to conserve the electro-neutrality, the following equation must be satisfied.
•
••
′
(8)
[TiTi ] = [YSr ]+2[VO ]
Therefore, both extrinsic defects introduced by Y incorporation at Sr-site and intrinsic
oxygen vacancies, which are formed under reducing conditions, contribute to the
formation of electrons and increase of conductivity. In case of YST8 sample, adsorbed
oxygen (i.e. 19.5%) exhibit a minimum value as shown in Table 1, indicates the
presence of high lattice oxygen in the sample. To maintain the electro-neutrality, high
amount of defects may generate and they could be compensated by Sr and oxygen
vacancies under reducing conditions leading to improvement in electrical conductivity
[17]. The conductivity of all samples is actually increased after exposure to hydrogen
atmosphere; the higher value reported for sample YST8 can also be attributed to the
large porosity (Fig. 3 (d)) which allowed for a more efficient exchange between the gas
phase and solid phase. This also suggests that the reducing stage performed (24 h at
700 °C) was not sufficient to get equilibrium in dense samples. In addition, the porosity
can be further increased by incorporation of A-site deficiency like (Y0.08Sr0.92)1-xTiO3-δ
and B-site deficiency like Y0.08Sr0.92Ti1-xO3-δ as reported by several authors [17, 18].
3.5 Chemical stability
To check the chemical stability of the YST samples, XRD measurements were
repeated after the reduction. Results are reported in Fig. 6, together with the patterns
obtained after sintering in air as discussed in section 3.1. It is observed that the XRD
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patterns of YST5 to YST10 (Fig. 6 (b to d)) overlap with each other showing no signs
of chemical and structural degradations. Whereas, the sample YST3 (Fig. 6 (a)) showed
the formation of additional secondary phases after reduction. In addition, corresponding
FESEM micrographs of reduced samples (inset in Fig. 6) also confirmed that the sample
YST3 strongly degraded with the formation of Ti-Sr mixed oxides. Also, for all the
compositions, the peaks have been shifted towards the left in the reducing atmosphere
showing lattice expansion. The lattice expansion might be attributed to increasing the
oxygen vacancies concentration under reducing atmosphere as per Eq. (6c).
Among all the investigated samples, the most promising candidate for use as
anodic material in SOFC is Y0.08Sr0.92TiO3 due to its good chemical stability coupled
with high conductivity and low activation energy under reducing condition.
4. Conclusions
In the present work, Y-doped SrTiO3 system was studied. By changing the
concentration of doping element, different incorporation mechanisms are activated,
particularly simultaneous incorporation of Y3+ at Sr- and Ti-sites or incorporation of Y3+
at Sr-site with electron compensation. XPS revealed the confirmation of electron
compensation mechanism. Upon reduction, the conductivity increased and the
activation energy decreased to nearly half in all the samples. Also, the sample YST8
was found to be chemically stable with 34% porosity having the highest electrical
conductivity under reducing conduction. While, the sample YST3 is strongly degraded.
Furthermore, HRTEM measurement may be incorporated to identify the beginning of
degradation of the samples as a future work. Hence, YST8 may be considered as a
promising anode material for solid oxide fuel cells.
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Acknowledgements
The authors acknowledge the BRNS (BARC) India under the Project Grant No.
34/14/15/2015/BRNS for financial support. One of the authors Saurabh Singh is
thankful to MHRD for providing Institute TA-ship.
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Table 1: Activation energy, relative density and adsorbed oxygen of YST samples
Relative Adsorbed
Ea (eV)
Ea (eV)
χ2 Volume Density Oxygen
O2 atmosphere
H2 atmosphere
Sample
(Å3)
(%)
68
(%)
200-400 ºC
400-700 ºC
1.465
200-400 ºC
400-700 ºC
0.668
YST3
YST5
YST8
YST10
7.69 56.015
26.4
1.043
0.222
0.689
0.800
0.511
0.124
0.100
0.197
7.78 55.883
7.78 54.742
7.81 54.556
92
66
95
65.9
19.5
41.3
0.139
1.355
0.497
0.332
0.259
0.279

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Fig. 1: (a) XRD patterns of the sintered samples of the system YxSr1-xTiO3 (i.e, YST0, YST3,
YST5, YST8, and YST10) (b) Angle shift of (110) peak with the Y doping.

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Fig. 2: (a) Secondary phase concentration, xs, as function of x ≥ 0.03 in YxSr1-xTiO3-δ (b)
Variation of lattice parameter with tolerance factor, (c) The variation of intensity of XRD
peak (110) via 2θ-δ (o) with compositions (x) for YxSr1-xTiO3-δ system, and (d) The variation
of microstrain and crystallite size with compositions (x).

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Fig. 3: FESEM images of fractured samples sintered at 1200 oC in air for (a) YST0 (b) YST3
(c) YST5 (d) YST8 and (e) YST10.

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Fig. 4: (a) The XPS survey spectra of Y doped SrTiO3 (b) The O-1s core level spectra (c)
The XPS spectra of Ti 2p of YST5 sample.

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Fig. 5: The Arrhenius plot of log σT vs. 1000/T (a) in O2 atmosphere (b) in H2 atmosphere.

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Fig. 6: XRD patterns of Y-doped SrTiO3 samples before and after reducing atmosphere and
inset FESEM micrographs of reduced samples.

ACCEPTED MANUSCRIPT
Highlights
The manuscript contains the following highlights:
ꢀ The mechanism for conductivity was predicted by Kröger-Vink notation with the help of
electron generation. This prediction was confirmed through structural measurements.
ꢀ The conductivity spectra were analysed through Arrhenius relation in air and H2
atmosphere.
ꢀ The chemical stability of the reduced Y-doped samples was checked by XRD
measurement.
ꢀ The conductivity data were analysed through impedance spectroscopy and concluded that
8 mol% Y-doped SrTiO3 is suitable as anode material for SOFC.