Enhanced photovoltaic effects in A-site 
samarium doped BiFeO 3 ceramics: The 
roles of domain structure and electronic state
Authors: C.- S. Tu, C.- S. Chen, P.- Y. Chen, H.- H. 
Wei, V. Hugo Schmidt, C.- Y. Lin, J. Anthoniappen, 
and J.- M. Lee
NOTICE: this is the author’s version of a work that was accepted for publication in Journal of the 
European Ceramic Society. Changes resulting from the publishing process, such as peer review, 
editing, corrections, structural formatting, and other quality control mechanisms may not be 
reflected in this document. Changes may have been made to this work since it was submitted for 
publication. A definitive version was subsequently published in Journal of the European Ceramic 
Society, VOL# 36, ISSUE# 5, (2015), DOI: 10.1016/j.jeurceramsoc.2015.12.019.
C.- S. Tu, C.- S. Chen, P.- Y. Chen, H.- H. Wei, V. Hugo Schmidt, C.- Y. Lin, J. Anthoniappen, 
and J.- M. Lee, “Enhanced photovoltaic effects in A-site samarium doped BiFeO 3 ceramics: The 
roles of domain structure and electronic state,” Journal of the European Ceramic Society 36, 
1149-1157 (2015). doi: 10.1016/j.jeurceramsoc.2015.12.019.
Made available through Montana State University’s ScholarWorks 
scholarworks.montana.edu 
Enhanced photovoltaic effects in A-site samarium doped BiFeO3
ceramics: The roles of domain structure and electronic state
Chi-Shun Tua,∗, Cheng-Sao Chenb, Pin-Yi Chenc, Hsiu-Hsuan Weia, V.H. Schmidtd,
Chun-Yen Lina, J. Anthoniappene, Jenn-Min Leef
a Department of Physics, Fu Jen Catholic University, New Taipei City 24205, Taiwan
b Department of Mechanical Engineering, Hwa Hsia University of Technology, New Taipei City 23567, Taiwan
c Department of Mechanical Engineering, Ming Chi University of Technology, New Taipei City 24301, Taiwan
d Department of Physics, Montana State University, Bozeman, MT 59717, USA
e Department of Physics, University of San Carlos, Cebu 6000, Philippines
f National Synchrotron Radiation Research Center, Hsinchu 30076, Taiwan
a b s t r a c t
This work reports enhanced photovoltaic (PV) responses of (Bi1− xSmx)FeO3 (x=0.0, 0.05, 0.10) ceramics
(BFO100xSm) with ITO film under near-ultraviolet irradiation (=405nm). The ceramics were charac-
terized by micro-Raman scattering, high-resolution transmission electron microscopy, and synchrotron
X-ray absorption spectroscopy (XAS). A rhombohedral R3c symmetrywith tilted FeO6 octahedra has been
confirmed. The FeK-edge absorption spectra reveal a slight shift toward higher energy as A-site Sm3+ sub-
stitution increases. The oxygen K-edge XAS reveals an enhancement of hybridization between the O 2p
and unoccupied Fe 3d states due to Smdoping. The optical band gaps are in the range of 2.15–2.24 eV. The
maximal PV power-conversion and external quantum efficiencies respectively reach 0.37% and 4.1% in
the ITO/BFO5Sm/Auheterostructure. The PV responses can be described quantitatively by a p-n-junction-
like model. The domain structures and hybridization between the O 2p and Fe 3d states play important
roles for the PV responses.
1. Introduction
Among perovskite multiferroic materials, the bismuth ferrite
BiFeO3 related systems are the most studied candidates because of
their room-temperature antiferromagnetic behavior (Néel temper-
ature, TN∼630K) and ferroelectric properties (Curie temperature,
TC ∼1100K) [1]. The ferroelectric polarization can be attributed
to the hybridization of Bi 6p and O 2p states, which causes an
off-center displacement of Bi3+ toward O2− [2]. It was reported
that the Curie temperature TC in rare-earth-doped compositions
of (Bi1− xREx)FeO3 (RE= La, Nd, Sm, Gd) decreases with decreasing
average A-site ionic polarizability and tolerance factor [3]. Rhom-
bohedral R3c BiFeO3 displays a G-type antiferromagnetic order, in
which each spin in the Fe3+ ions is surrounded by six anti-parallel
neighbor spins [1,4].However, leakage currents andweakmagnetic
behavior in BFO are drawbacks for applications [1]. To increase the
ferroelectricity and ferromagnetism, many studies have focused
on rare-earth substitution in the A-site Bi3+ positions of the per-
ovskite unit cell, resulting in improvements in the magnetic and
ferroelectric properties [5,6].
Early studies of rare-earth-doped BFO materials focused on
structural and magnetic characterization [3–16]. The structures
of (Bi1-xREx)FeO3 (RE= La, Nd, Sm, Gd) for x≤0.1 correspond
to the perovskite rhombohedral R3c phase of BiFeO3 [3]. The
(Bi1− xSmx)FeO3 solid solutions show a ferroelectric rhombohe-
dral phase for x=0−12%, a coexistence of ferroelectric triclinic
and orthorhombic phases for x=12.5−20.0%, and a non-polar
orthorhombic phase for x=25% at room temperature [6]. Fer-
roelectric polarization hysteresis loops appear in the range
of 0≤ x≤17.5% in (Bi1-xSmx)FeO3 ceramics [6]. Temperature-
dependent dielectric permittivity results suggest that themagnetic
Néel temperature (TN) decreases with increasing Sm3+ concen-
tration for x=1−8% [7]. (Bi1− xSmx)FeO3 ceramics and thin films
show enhanced magnetization with increasing Sm ratio [12–16].
(Bi1− xSmx)FeO3 ceramics show a magnetic structure transition
fromaspincycloid toaG-typeantiferromagneticbehavior at x∼14%
[15,16].
Fig. 1. (a-c) SEM grain morphologies of thermally etched ceramics and (d) XRD spectra of as-sintered samples at room temperature. D¯ and A¯O are average grain size and
average oxygen atomic ratio, respectively.
Perovskite ferroelectric/piezoelectricoxideshavebeenexplored
for photovoltaic (PV) applications for several decades because
of easy electric-field-driven polarization [17–20], but a severe
drawback for ferroelectric photovoltaic devices has been the
low photovoltaic current under illumination. In recent photo-
voltaic studies [21–41], BFO thin films and crystals with various
electrodes have shown properties with potential for PV applica-
tions. Pt/BFO/SRO and Pt/Sm:BFO/SRO thin films heterostructures
showed significantly higher PV current densities in the low-
resistance state under white-light illumination [30,31]. BFO has
been considered a p-type semiconducting material resulting from
the Bi3+ loss during the sintering process, which causes vacancies
to act as p-type centers [42]. First-principles calculations suggested
that Bi vacancies (VBi) have lower formation energy than oxy-
gen vacancies under oxygen-rich condition and thus VBi become
electron-acceptor defects [43].
The aim of this work is to investigate the electronic config-
urations and domain structures and their correlations in the PV
responses of Sm3+ doped BFO polycrystalline ceramics with ITO
and Au electrodes. Fe K-edge, Fe L2,3-edge and oxygen K-edge
synchrotron X-ray absorptions were measured to study electron
structures and the hybridization of the O 2p and the unoccupied Fe
3d states. High-resolution TEM and diffraction were used to exam-
ine the nanoscale domain structures and symmetry. A theoretical
p-n-junction-like model is used to quantitatively describe the PV
open-circuit voltage and short-circuit current density as functions
of illumination intensity.
2. Experimental procedure
BiFeO3 (BFO), (Bi0.95Sm0.05)FeO3 (BFO5Sm), and
(Bi0.90Sm0.10)FeO3 (BFO10Sm) ceramics were prepared by the
solid state reaction, in which Bi2O3, Sm2O3, and Fe2O3 powders
(purity ≥99.0 %) were weighed in the ratios of 1.1:0:1.0 for BFO,
0.95:0.05:1.0 for BFO5Sm, and 0.90:0.10:1.0 for BFO10Sm. The
powdersweremixed in an agatemortarwith alcohol formore than
24h and then were calcined at 800 ◦C for 3h. The sintering temper-
atures were 830 ◦C (10h) for BFO and 870 ◦C (3h) for BFO5Sm and
BFO10Sm ceramics. Grain morphologies, oxygen atomic ratios, and
lattice structureswere characterized using aHitachi S-3400N scan-
ning electron microscope (SEM) equipped with energy-dispersive
X-ray spectroscopy (EDS) and a Rigaku Multiplex Diffractometer.
For SEM measurement, the samples were polished and thermally
etched at 800 ◦C for 30min. Standard deviations were calculated
to estimate error ranges of oxygen atomic ratios. The average grain
sizes were estimated by counting the number of grains intercepted
by several straight lines sufficiently long to include most of grains
on the SEM photomicrograph.
Micro-Raman spectra were measured using a Nanobase Model
XperRam 200 Raman spectrometer equipped with a green laser
of =532nm and a TE-cooled CCD detector. A high-resolution
TEM (JEOL JEM-2100 LaB6) was used to study domain structures
and symmetries of the unit cells. To determine electronic states,
the Fe K-edge synchrotron X-ray absorption near edge structure
(XANES) and extended X-ray absorption fine structure (EXAFS)
were obtained in transmission mode at the 01C1 beam line of
the National Synchrotron Radiation Research Center (NSRRC) in
Taiwan. The soft X-ray absorption spectra (XAS) of the Fe L2,3-edges
and oxygen K-edge were studied in total electron yield via current
mode at the 20A1 beam line of the NSRRC. For PV measurements
(I–V characteristic curve, open-circuit voltage, short-circuit cur-
rent, and power conversion efficiency), ITO (top electrode) and Au
(bottom electrode) thin films were deposited on the ceramic sur-
faces by dc sputtering. A diode laser (=405nm) was used as the
excitation source and the illuminated area (also ITO area) is ∼0.15
cm[2]. The ceramic thicknesses are 0.2mm for BFO and 0.15mm
for BFO5Sm and BFO10Sm.
C.-S. Tu et al. / Journal of the European Ceramic Society 36 (2016) 1149–1157 1151
3. Results and discussion
Fig. 1(a–c) shows grain morphologies of thermally etched
ceramics with average grain sizes (D¯) of 3.4, 1.8, and 1.3m for
BFO, BFO5Sm, and BFO10Sm. As labeled in Fig. 1(a–c), the aver-
age oxygen atomic ratios (A¯O) for BFO, BFO5Sm, and BFO10Sm
are approximately 52.5±1.6%, 53.3±0.5%, and 55.6±0.6%, respec-
tively. The average oxygen atomic ratios were determined from 15
to 18 spots in the region where ITO thin film was later deposited
for the PV experiment. The average oxygen atomic ratio increases
with increasing Sm concentration, suggesting that oxygen vacan-
cies were reduced by the A-site Sm3+ substitution. This result is
consistentwith previous reports on rare-earth doped BiFeO3 mate-
rials [5]. XRD spectra of as-sintered samples are shown in Fig. 1(d),
which show similar splittings in XRD peaks, suggesting a rhom-
bohedral phase in all compounds. Compared with BFO, BFO5Sm
and BFO10Sm exhibit stronger (100) and (200) peaks, suggest-
ing a preferred (100) crystallographic orientation in the ceramic
grains. As illustrated by the dashed line, XRD peaks of BFO5Sm and
BFO10Smshift towardhigher 2, due to the smaller radius (1.098Å)
of the Sm3+ ion compared with 1.17Å for the Bi3+ ion [44].
Fig. 2(a) shows Raman vibration modes in the range of 50-600
cm−1, including Fe2O3, Bi2O3, and Sm2O3 powders. The Raman
Fig. 2. (a) Raman spectra in the region of 50–600 cm−1 and (b) fitting modes in the
region of 50–250 cm−1. The labeled numbers are frequencies and full-width at half
maxima (in parentheses). The red line is the sum of the fitting curves.
spectra suggest no second phase from incomplete reaction in the
as-sintered Sm-doped compounds. The Raman active modes in the
R3c rhombohedral BFO can be represented by =4A1 + 9E [45–47].
BFO, BFO5Sm, and BFO10Smexhibit similar Raman vibrations, con-
firming the same R3c space group. To identify the Sm3+ doping
effects on atomic vibrations, the four dominant Ramanmodes, E(1),
E(2),A1(1), andA1(2),werefittedbyusing a Lorentzianprofile in the
region of 50–250 cm−1 as shown in Fig. 2(b). The numbers labeled
onpeaksareRaman frequencies and full-widthathalfmaximum(in
parentheses). The low-frequency Raman modes (below 170cm−1)
and mid-frequency modes (150-270 cm−1) are respectively associ-
ated with Bi and Fe atoms [45]. The higher Raman modes (above
260 cm−1)mainly correlate to oxygenmotions. The vibrationmode
near 220 cm−1 is the A1 tilt mode of the FeO6 octahedra. Major
Raman modes, as illustrated by the dashed lines, shift significantly
to higher frequencies with increasing Sm concentration, indicating
that the A-site Bi3+ ions have been mostly substituted by lighter
Sm3+ ions. The E(1), E(2), A1(1), and A1(2) modes of BFO5Sm and
BFO10Smexhibit broader spectra than BFO, possibly resulting from
lattice anharmonicity and disorder caused by randomness of the
A-site Sm3+ substitution.
Fig. 3 shows the high-resolution bright-field TEM images (top)
and the selected-area diffraction patterns (SADPs). The inset in
Fig. 3(a) shows the domain configuration of BFO measured from
a two-beam condition via tilting technique. The arrows in the
bright-field TEM images indicate various domain boundaries. BFO
and BFO10Sm exhibit complex pinning domain structures at the
boundary intersections, which can cause high strain and reduce
electrical conductivity. BFO5Sm exhibits relatively homogeneous
domain structures with linear domain walls. The domain morphol-
ogy transforms from small nanodomains (in BFO) to large lamellar
domains in the Sm-doped compounds, suggesting a transformation
from the short-range-ordered to the long-range-ordered states.
Enhanced electrical conductivity has been reported along uniform
domain walls in (110)-oriented La3+ doped (10%) BFO thin films
[48]. It was proposed that the enhanced conduction at the domain
walls was driven by a local reduction in the electronic band gap
[48]. Therefore, higher electrical conductivity can be expected in
BFO5Sm because of uniform domain walls in the grain matrix.
The SADPs were measured from representative grains along
the [11¯0] zone axis. The [11¯0]-viewed SADPs reveal 1⁄2(111)
(or 1⁄2{ooo}) superlattice diffraction spots without appearance of
1⁄2{ooe} diffraction spots. Here, “o” and “e” represent odd and even
Miller indices. Thisphenomenoncanbeattributed to the tiltedFeO6
octahedra, indicating a predominantly rhombohedral R3c symme-
try [49,50]. The high-resolution TEM confirms the rhombohedal
R3c space group as suggested by the XRD (Fig. 1d) and Raman
spectra (Fig. 2). The intensity of the 1⁄2(111) superlattice diffrac-
tion is enhanced as Sm content increases, suggesting that the FeO6
octahedral tilting is enhanced by the Sm substitution.
Fig. 4 displays the Fe K-edge XANES spectra of BFO, BFO5Sm,
BFO10Sm, and the reference powders FeO and Fe2O3. The FeK-edge
absorption arises primarily from the 1s→4p electronic transition
[51]. The Fe K-edge absorption is sensitive to the octahedral envi-
ronment and local electrostatic interaction. As indicated by lines
and labels in the inset, BFO shows a K-edge absorption peak at
approximately 7133eV, which is close to the K-edge energy of ref-
erence powder Fe2O3. This result suggests an Fe3+ oxidation state
in BFO. The Fe K-edge absorption peaks in BFO5Sm and BFO10Sm
appear respectively at approximately 7134eV and 7136eV, indi-
cating that the oxidation state of Fe3+ does not transfer into Fe2+
as Sm concentration increases. It is also consistent with the EDS
result (Fig. 1), in which BFO5Sm and BFO10Sm exhibit fewer oxy-
gen vacancies. First-principles density functional theory suggests
that the incorporation of oxygen vacancies can lead to the forma-
tion of the Fe2+ oxidation state [52]. As indicated by the arrow near
Fig. 3. Bright-field TEM images and selected-area diffraction patterns (SADPs).
Fig. 4. Fe K-edge and Sm L2-edge absorption spectra vs. photon energy. The inset is
an enlargement with energy labels for the absorption peaks.
7110eV, weak pre-edge peaks appear at ∼5eV below the major
absorptions and can be attributed to the dipole-forbidden 1s→3d
transitions [51,53]. The 1s→3d pre-edge transition is associated
with electric quadrupole and dipole states, and involves 4p mixing
with the 3d orbital configuration around the Fe site [53]. Fig. 4 also
shows the Sm L2-edge absorption peaks near 7320eV in BFO5Sm
and BFO10Sm, due to the 2p→5d transition [54,55].
Fig. 5(a) shows the Fe L2,3-edge XAS spectra, which correspond
to the transition from the Fe 2p core level to the unoccupied Fe 3d
states [56]. The spectra include the L3 (2p3/2) band at ∼710eV and
the L2 (2p1/2) band at ∼723eV. The splittings of the L2,3 absorptions
are due to the t2g and eg orbital configurations [57]. The absorption
configurations and energy peaks of the Fe L2,3 bands are similar
to the L2,3-edge band structures of -Fe2O3 [58,59] and LaFeO3 as
found from first-principles calculations [57]. This result suggests
that the dominant oxidation states of the Fe ion are Fe3+ in all
three compounds and are consistent with the Fe K-edge spectra as
Fig. 5. (a) Fe L2,3-edges and (b) O K-edge absorption spectra vs. photon energy.
revealed in Fig. 4. Fig. 5(b) shows the normalized O K-edge XAS due
to the transition from the O 1s core level to O 2p states hybridized
with the unoccupied Fe 3d orbital configurations [56,60,61]. The
well-resolved profiles in BFO5Sm and BFO10Sm suggest crystalline
structures as revealed in the (001) XRD peaks in Fig. 1(a) [62]. The
A and B peaks can be identified as the t2g and eg orbital bands sep-
arated by the ligand-field (LF) splitting [63,64]. The LF splitting is
associated with the local electrostatic interaction of oxygen with
the Fe 3d orbitals [61]. The LF splittings between the t2g and eg
bands are ∼1.3 eV in BFO and ∼2eV in BFO5Sm and BFO10Sm. The
intensity ratios between the t2g and eg bands are respectively about
1.8:1.7, 2.4:2.1, and 2.2:2.0 for BFO, BFO5Sm, and BFO10Sm.
In-Fe2O3, the LF splitting is about 1.3 eV and the intensity ratio
of the t2g and eg bands is approximately 1:1 [60,61]. The intensities
of the t2g and eg bands correlate to the number of Fe 3d holes mod-
ulated by the hybridization effects [61]. There are three t2g and two
eg holes in iron and thus an intensity ratio of 3:2 is expected with-
out the hybridization effect in -Fe2O3 [61]. As the eg orbitals shift
toward the oxygen ligands, the O 2p–eg hybridization is stronger
than the O 2p–t2g hybridization, thus a ratio of 1:1 could be experi-
mentallyobserved [61].However, aspointedoutbyGroot et al. [60],
the splitting of the d orbitals is complex and other mechanisms can
affect the relative intensities of the t2g and eg bands, such as elec-
tronic exchange interactions and the non-stoichiometry effect. The
intensities of the A andBpeaks in Fig. 5(b) show strong dependence
on the Sm concentration, although their absorption energies do not
shift remarkably. This result suggests that the hybridization of the
O 2p and Fe 3d states depends on the Sm atomic ratio. In BFO mate-
rials, the conduction band is mainly composed of the Fe 3d state
hybridized with the O 2p state. Thus, the stronger hybridization of
Fe3dandO2p inBFO5Smsuggest anenhancedelectrical conductiv-
ity. The energy positions and shapes of the O K-edge bands are very
similar to the O K-edge structures of -Fe2O3 [61,63], confirming
the Fe3+ oxidation state in BFO, BFO5Sm, and BFO10Sm.
The broad C band in the region of 535–545eV in Fig. 5(b) can
be attributed to the mixture of Bi 6p, Sm 6s, Fe 4s/4p, and O 2p
configurations[61,62]. BFO shows a broader C band than BFO5Sm
and BFO10Sm, indicating a highly disordered configuration of elec-
tronic states. The broadness of the XAS lines has been considered as
an intrinsic property of oxygen vacancies[65]. This result suggests
that Sm doping can reduce oxygen vacancies. This phenomenon is
consistent with the average oxygen atomic ratios obtained by SEM
EDS as given in Fig. 1(c). As indicated by the arrow in Fig. 5(b), the
pre-edge peaks occur at ∼2eV below the A peak and are correlated
to the dipole transition from O 1s to O 2p up-spin states hybridized
with the unoccupied Fe 3d up-spin states [62]. This implies that the
Fe 3d5 orbitals are in the high spin polarization with e2gt
3
2g electron
configuration, as based on ligand-field theory [64].
Fig. 6 shows plots of optical transmission and (h) [2] vs.
photon energy (hv) measured by using a Cary 5E UV-Vis-NIR spec-
trometer. h and  are Planck’s constant and photon frequency. The
optical absorption coefficient  was determined by ˛ = − ln(T)/d,
where T and d are the measured optical transmission and sam-
ple thickness. The optical band gap can be estimated using the
Tauc relation [66], (˛hv)2 = A(hv − Eg), where Eg is the optical
band gap between valence and conduction bands, and A is a
material-dependent constant. Eg can be evaluated by extrapolat-
ing the straight line portions of the curves of (h)[2] vs. hv as
given in Fig. 6(b). BFO5Sm and BFO10Sm have smaller Eg values
of 2.18 eV and 2.15 eV than Eg∼2.24 eV in BFO. Rhombohedral BFO
thin films have been reported with direct band gaps in the range of
∼2.55–2.75 eV [67–70]. Band gaps of 2.18 and 2.1 eV were reported
respectively for BFO nanoparticles [71] and nanocubes (with sizes
of 0.05–0.2m) [72]. Band gaps of BFO nanoparticles show strong
dependence on particle size and were correlated to micro-strain
and oxygen defects [73]. From first-principles density functional
theory, the oxygen vacancies (OVs) in BFO can shift the optical
absorption to lower energy below the conduction band [74,75]. The
slightly different band gaps in BFO, BFO5Sm, and BFO10Sm may be
correlated to OVs and grain sizes.
Fig. 7 shows PV open-circuit voltages (Voc) and short-circuit
current densities (Jsc) of ITO/(Bi1-xSmx)FeO3/Au heterostructures
Fig. 6. (a) Optical transmission and (b) (h)[2] vs. photon energy.
Fig. 7. (a) Open-circuit voltage (Voc) and (b) short-circuit current density (Jsc) as
light was switched on and off with increasing intensity. The labeled numbers are
light intensities in mW/cm2.
Table 1
Comparison of PV studies in perovskite ferroelectric and BiFeO3 materials. PZT=Pb(Zr,Ti)O3; PLZT= (Pb,La)(Zr,Ti)O3; BSFCO= (Bi0.9Sm0.1)(Fe0.95Co0.05)O3;
LSCO= (La0.67Sr0.33)CoO3; LSMO= (La,Sr)MnO3; STO=SrTiO3; SRO=SrRuO3; AZO=Al2O3–ZnO. ITO is indium tin oxide. FTO is a commercial glass substrate.
Photovoltaic heterostructures Voc (V) Jsc (mA/cm2) Light wavelength 
(nm)
Light intensity
(mW/cm2)
Efficiency (%)
Pt/PZT(52/48)/Ni [17] 0.8 6×10−5 300–390 1.0
Pt/PZT(20/80)/Pt [18] 8×10−3 350–450 10
ITO/poly-PZT(53/47)/ITO [19] 0.25 5×10−6 632 0.45 0.22
Nb:STO/PLZT(3/52/48)/ LSMO [20] 0.7 0.0008 0.06 0.28
ITO/PZT/Cu2O/Pt [83] 0.4 4.5 Sunlight 100 0.57
ITO/BFO/SRO [21] 0.3 4×10−4 435 0.75
Pt/BFO/Pt [22] 20 0.05 375 100
Pt/BFO crystal/Pt [23] 13 0.001 405 3×10−5 (EQE)
ZnO:Al/BFO/LSCO [24,25] 0.22 0.004 White light 1
Pt/BFO/SRO [26] 0.08 0.063 White light 100
Pt/BSFCO/Pt [27] 0.9 5×10−5 White light 100
Pt/0.9BFO-0.1YCrO3/Pt [28] 0.51 0.00148 Sunlight 100
Pt/BFO/STO [29] 3.9 2.6×10−4 White light 100
Nb:STO/BFO/Au [32] 0.04 6 White light 285 0.03
Ag/Pr-BFO NTs/Ag [33] 0.2 Sunlight 10 0.5
AZO/poly-BFO/FTO [34] 0.63 0.13 Sunlight 100 0.0208 7 (EQE)
at =340 nm
Fe/BFO/LSMO/STO [35] 0.21 6×10−4 White light 20
Au/BFO/Au [36] 0.0075 532 <20
Graphene/poly-BFO/Pt [37] 0.44 0.025 Sunlight 100
ITO/poly-BFO/Pt [38] 0.1 0.0025 0.45 0.125
Pt/BFO/Pt [40] 16 0.12 White light 285
ITO/BFO/SRO [41] 0.8− 1.5 Sunlight 285 10 (EQE) at
=325 nm
ITO/BFO ceramic/Au (this work) 0.58 0.005 405 10 0.005 0.16
(EQE)
TO/BFO5Sm ceramic/Au (this work) 0.7 0.094 405 10 0.25 3.0 (EQE)
as the laser (=405nm) was switched on and off with increas-
ing illumination intensity (I) by steps. The illumination-intensity
dependent Voc and Jsc are plotted in Fig. 8(a) and (b). The
ITO/(Bi1− xSmx)FeO3/Au heterostructures exhibit rapid increase in
PV response below I∼50 mW/cm2. We apply a previously devel-
oped p-n-junction-like model to describe Voc and Jsc as functions of
illumination intensity[76]. The photodiode current id under a bias
voltage V is expressed as [77].
id = io
{
exp
[
q (V − idRs)
kBT
]
− 1
}
. (1)
Here, io, Rs, and  are dark current, source resistance, and
diode-quality factor, respectively. To estimate io,Rs, and , the char-
acteristic curves were measured in the dark as shown in Fig. 9. By
using Eq. (1) with q=1.6×10−19 C and T=300K, we obtained io,
Rs, and  as given in Fig. 9. From Eq. (1), the illuminated i can be
expressed as [76].
i = ip − id = ip − io(exp{[V − (id − ip)Rs]
q
kT
} − 1). (2)
where ip, id, and V are the photovoltaic current, diode current,
and measured voltage. Through a considerable derivation [76], the
PV open-circuit voltage (Voc) and short-circuit current (isc) can be
expressed as [76].
Voc = Uo − Bˇ2i2o
[
exp(Vocq/kT) − 1
]2
(qSI/hc)2
(3)
isc = Uo
Rs
− Bˇ2
isc + io[exp(iscRsq/kT) − 1] 2
Rs(qSI/hc)
2
(4)
B = qnp
2εoεp
)(
1 + npεp
εnnn
(5)
where S,, h, c, and I are illumination area, lightwavelength, Planck
constant, light speed, and illumination intensity. Uo is the voltage
barrier across thedepletion region in thedark. is theoptical atten-
uation length determined directly from the optical transmission in
Fig. 6(a), i.e. ˇ = −d/ ln(T). The measured  values at =405nm
are respectively about 2.0, 1.0 and 1.0m for BFO, BFO5Sm, and
BFO10Sm. np and nn are carrier densities of p-type (Bi1-xSmx)FeO3
ceramics and n-type ITO film. εp and εn are p-type and n-type
dielectric permittivities. εp values of BFO, BFO5Sm, and BFO10Sm
ceramics are respectively about 100, 145, and 108 for measuring
frequency f=1MHz at room temperature. The reported carrier den-
sity [78] and dielectric permittivity [79] of n-type ITO thin films are
nn ∼1021 cm−3 and εn ∼4 at photon energy of 3 eV. FromEq. (5), the
calculated np values are about 3×1020, 3×1019, and 2×1020 cm−3
for BFO, BFO5Sm, and BFO10Sm ceramics, respectively. The junc-
tion widths in the dark, do∼(Uo/B)1/2 [76], are about 2, 17, and 3nm
for BFO, BFO5Sm, and BFO10Sm, respectively.
The solid lines in Figs. 8(a) and (b) are fits of Voc and Jsc (= isc/S)
by using Eqs. (3) and (4) with fitting parameters in Fig. 8(a). The
theoretical fits of Voc agree well quantitatively with experimental
data. The fits of Jsc in Fig. 8(b) show discrepancies with the exper-
imental points in the low-illumination region for ITO/BFO5Sm/Au,
likely due to the leakage currents. Note that this p-n-junction-
like model only considers the photo-excited electron-hole creation
[76]. Other factors, such as oxygen and Bi vacancies, domain struc-
ture, electronic states, and charge recombinationmay influence the
conductivity [73].
Fig. 8(c) shows external quantum efficiencies (EQE) calculated
from the short-circuit current density (Jsc) in Fig. 8(b). The EQE is
the conversionefficiency from incidentphotons to conductionelec-
trons, i.e. EQE = hfJsc/qeI [76], where I and qe are light intensity and
electroncharge. ThemaximalEQEof ITO/BFO5Sm/Aureachesabout
4%, which is significantly higher than the 0.00003% reported in the
Pt/BFO crystal/Pt configuration under illumination of =405nm
[23], and is comparable to the 10% observed in ITO/BFO/SrRuO3 at
ultraviolet wavelength of =325nm [41].
Fig. 10 gives curves of power-conversion efficiency () vs.
load voltage (V) for several illumination intensities. The power-
conversion efficiencies were calculated using Pout/Pin. Pin and Pout
Fig. 8. (a) Voc, (b) Jsc, and (c) external quantum efficiencies (EQE) as a function of
illumination intensity. The solid lines in (a) and (b) are theoretical fits of Eq. (3) and
(4) with parameters in (a). The dashed lines in (c) are guides for the eye.
Fig. 9. Characteristic curves of current vs. bias voltage without illumination. The
solid lines are fits of Eq. (1) with parameters given in the Figure.
Fig. 10. Power-conversion efficiency () vs. load voltage for various illumination
intensities.
are respectively the incident illuminationpower andelectric power
output measured from the curves of current vs. load voltage
under illumination. Themaximal PV power-conversion efficiencies
(max) occur at low light intensities and reach respectively about
0.017%, 0.37%, and 0.13% for ITO/BFO/Au, ITO/BFO5Sm/Au, and
ITO/BFO10Sm/Au. Themicrostructure, domain structures, polariza-
tion, and electronic states are believed to play important roles in
the photovoltaic conductivity [80]. Compared with BFO as shown
in Fig. 1(d), BFO5Sm and BFO10Sm ceramics exhibit a relatively
stronger (100) orientation, which may enhance the photocurrent,
similar to the effect of ordered polarization as reported in (100)
oriented BFO films [81,82].
To compare the photovoltaic responses with previous reports,
a brief summary of photovoltaic studies in some major perovskite
ferroelectric and BFO-related materials is given in Table 1, includ-
ing the results from this work. At low illumination intensity,
the ITO/BFO5Sm/Au heterostructure has demonstrated compara-
ble photovoltaic open-circuit voltage (Voc), short-circuit current
density (Jsc), power-conversion efficiency (Pout/Pin), and external
quantum efficiency (EQE) with most other PV devices using per-
ovskite ferroelectric and BiFeO3 thin films and crystals as listed in
Table 1.
4. Conclusions
Enhanced photovoltaic (PV) responses have been observed in
ITO/(Bi1-xSmx)FeO3/Au heterostructures for x=0.05 and 0.10 under
near-ultraviolet irradiation of =405nm. The maximal PV power-
conversion and external quantum efficiencies reach 0.37% and 4.1%
in ITO/BFO5Sm/Au. The PV open-circuit voltage and short-circuit
current density can be described quantitatively by a p-n-junction-
like theoretical model as functions of illumination intensity. This
work reveals that oxygen vacancies were reduced by A-site Sm3+
substitution. A slight shift toward higher energies was observed in
the Fe K-edge absorption spectra due to the A-site Sm3+ substitu-
tion. Domain structures and hybridization between the O 2p and Fe
3d electronic configurations play important roles in the enhanced
PV effects. The oxygen K-edge absorption indicates that the A-site
Smsubstitution enhances hybridizationof theO2p andunoccupied
Fe 3d states. The optical band gaps of BFO, BFO5Sm, and BFO10Sm
are respectively about 2.24, 2.18, and 2.15 eV.
Acknowledgement
This work is supported by the Ministry of Science and Tech-
nology of Taiwan under Project Nos. MOST 104-2221-E-030-014,
104-2221-E-146-001, and 102-2221-E-131-006.
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