Trans. Mat. Res. Soc. Japan

41[1] 101-108 (2016)

Substitution effect of tetravalent and pentavalent elementson thermoelectric properties in In2O3-SnO2 system

Rintaro Abe, Hiroyuki Fujishiro* and Tomoyuki Naito

Faculty of Engineering, Iwate University, Morioka 020-8551, Japan* Corresponding author: Fax: 81-19-621-6363, and/or e-mail: fujishiro@iwate-u.ac.jp

Thermoelectric properties of In2-xMxO3 system doped with tetravalent metal (M) ions (Si4+, Ti4+,
Ge4+, Sn4+, Te4+, and Ce4+) were investigated at temperatures up to 1273 K. For theGe-substituted system, the power factor P takes a maximum of 9.5 × 10-4 W/K2m around 1150 K
and dimensionless thermoelectric figure of merit ZT shows a maximum of 0.18 at 1000 K for the
x=0.025 sample, which is the solubility limit of Ge. The maximum ZT of the In1.9M0.1O3 series
was about 0.13 around 1000 K for the Si4+, Ge4+ and Sn4+ doped samples. We also investigated
the thermoelectric properties of the In2O3-SnO2 system doped with pentavalent Sb element toreplace the expensive In2O3 to the plentiful SnO2. The Sb atom was preferentially substituted forthe In-site and the thermoelectric properties do not exceed those of pure In2O3. The potential of
the thermoelectric properties in the doped In2O3-SnO2 system was discussed.
Key words: thermoelectric material, In2O3, SnO2, spark plasma sintering, dimensionless figure of merit,
power factor
1. INTRODUCTION
The thermoelectric technology using waste heat for
power generation devices has been revived because of
the recent energy crisis together with other renewableenergy sources. Since the discovery of NaCo2O4 in 1997[1], oxide thermoelectric materials have been intensively
studied, as a substitute for conventional intermetallic
thermoelectric materials using heavy metal series suchas Bi2Te3 [2-4]. Despite their lower performance ascompared to the intermetallic materials, oxides are very
attractive for high temperature applications due to theirchemical stability in air.
Recently, doped indium oxide (In2O3) compounds
are reported to exhibit relatively high thermoelectricperformances. The substitution of Ge for In was reported
to decrease the electrical resistivity and, at the same time,
the precipitation of fine particles of the In2Ge2O7
impurity phase efficiently reduced the thermal
conductivity, resulting in the higher dimensionlessfigure of merit of ZT=0.45 at 1273 K in air (=S2T/ρκ, S:
Seebeck coefficient, T: absolute temperature, :electrical resistivity, and : thermal conductivity) in theIn1.8Ge0.2O3 system [5]. The single element substitutionof Co [6], Sn [7], and co-substitution of Zn+Nd [8],Zn+Ge [9], Zn+Ce [10], M+Sn (M=Zn, Cu, Ni) [11] forthe indium site has been recently investigated and ZT
values as high as 0.3 can be obtained at hightemperatures. The approaches of fine-grained and highpacking were also performed to decrease the thermal
conductivity [12, 13]. However, systematic study of thesingle element substitution for the indium site has notinvestigated yet, and the reason why the Ge-substituted
In2O3 shows the higher ZT value, has not clarified
compared to other element substitution.

In this study, we investigated the In2O3 polycrystals
doped with various tetravalent elements (Si, Ti, Ge, Sn,
Te and Ce) and measured the electrical and thermal
properties. The better effect of the Ge-doping on thethermoelectric properties in In2O3 system was discussed.
In order to reduce the amount of In as a rare metalelement, we also investigated thermoelectric propertiesof the In2O3-SnO2 system doped with pentavalent Sbelement. Because SnO2 shows similar physicalproperties to In2O3, which is inexpensive, plentiful andnontoxic, as a substitute for the expensive In2O3. The
Sb5+ doping is expected to perform as a donor by boththe In3+- and Sn4+-site substitutions simultaneously. The
potential of the In2O3-SnO2 system was discussed as athermoelectric material.
2. EXPERIMENTAL PROCEDURE
In2-xGexO3 (0≤x≤0.3), In1.9M0.1O3 (M=Si, Ti, Ge, Sn,
Te, Ce) and Sb-doped In2O3-SnO2 systems wereprepared by the spark plasma sintering (SPS) method.
In2O3 (99.9%), SiO2 (99.9%), TiO2 (99.99%), GeO2
(99.99%), SnO2 (99.9%), TeO2 (99.9%), CeO2 (99.9%)
and Sb2O5 (99.99%) fine powders were used as raw
materials. The powders were mixed with stoichiometoricratio for 1 h by automatic mortar and were sintered by
the SPS apparatus (NJS-LABOX-110C) at about 1273 K
for 5-10 min in vacuum. The sintered bulks were cut
into rectangles and annealed at 1173 K in air for 2 h to
reduce oxygen defects. Powder X-ray diffraction (XRD)
measurements were performed (Rigaku Multi Flex) at
room temperature using Cu K radiation in the range of
20° ≦ 2 ≦ 90° with 0.02° steps, and the lattice
parameters were determined using a Rietveld refinement
program, RIETAN 2000 [14]. Scanning electron

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Substitution effect of tetravalent and pentavalent elements on thermoelectric properties in In2O3-SnO2 system

microscope (SEM) images were taken and quantitative
element analyses were performed by electron probe
micro-analyzer (EPMA). The electrical resistivity ρ(T)(or electrical conductivity σ(T)=1/ρ(T)) and Seebeckcoefficient S(T) were simultaneously measured in thetemperature range from 300 to 1273 K using an
automated measuring system (Ozawa Science RZ2001i)and the thermoelectric power factor P=S2/ρ was
calculated. The thermal conductivity κ(T) was measured
by a laser flash method (Ulvac-Riko TC-7000) from 300
K to 1023 K. The  value below 300 K was alsomeasured by steady-state heat flow method using
home-made measuring system [15]. A dimensionlessfigure of merit ZT=S2T/ρκ was estimated using thesevalues. Hall effect experiments were carried out by
five-terminal method at room temperature to determine
the carrier concentration n and Hall mobility , which
was calculated by μ=σ|RH| (RH: Hall coefficient).
3. RESULTS AND DISCUSSION
3.1 In2-xGexO3 (0≤x≤0.3)
Figure 1 shows the XRD patterns of the In2-xGexO3
system (0≤x≤0.3). The impurity peaks of In2Ge2O7
appeared for x≥0.05. Berardan et al. reported that thesolubility limit of Ge in In2O3 was about x=0.01~0.02 byconventional solid-state reaction method [5], which wasslightly lower than that of our samples. The increase of
the solubility limit in our samples may result from thedifference in the fabricating method.

scattering by In2Ge2O7 precipitates.
Figure 4(a) shows the temperature dependence of thethermal conductivity  of the In2-xGexO3 system. For thepure In2O3, (T) is inversely proportional to T, which
results from phonon-phonon scattering. The κ value at
300 K decreased with increasing Ge contents x and (T)of the Ge-substituted samples is less temperaturedependent, which may come from the phonon scatteringby In2Ge2O7 impurity particles. Figure 4(b) depicts thetemperature dependence of dimensionless figure of merit
ZT of the In2-xGexO3 system. The ZT value increases
with increasing T and that of the Ge-substituted samples
is larger than that of pure In2O3 at high temperatures.
The maximum ZT value was 0.18 at 1023 K for the
x=0.025 sample, which was smaller than the reported
values for the In2-xGexO3 system; ZT=0.3 for the x=0.2
sample was interpolated at 1000 K from the Berardan’s
data [5].
Figure 5 shows the ZT and P values at 1000 K, as afunction of x. Both ZT and P values take a maximum atthe same x of 0.025, which is the solubility limit of Ge in
In2O3. The x=0.025 sample shows the best
thermoelectric performance both electrically andthermally in the In2-xGexO3 system. We compare ourresults of the x=0.025 sample at 1000 K to those of the
x=0.2 sample at 1000 K by Berardan et al. [5]. The P
value of our sample (=0.95 mW/K2m) is larger than that

Fig. 1. XRD patterns of In2-xGexO3 (0≤x≤0.3). Arrows
indicate the In2Ge2O7 impurity peaks.
Figure 2 presents the temperature dependence of theelectrical resistivity ρ, the Seebeck coefficient S and thepower factor P of In2-xGexO3 (0≤x≤0.3). ρ and |S| values
decreased with increasing x up to x=0.05 and thenslightly increased with increasing x. These results
suggest that electron was doped by the Ge substitution
up to the solubility limit. The decrease of the |S| value
accompanied by the decrease of the  value is a generaltrend in condensed matter. The power factor of the
x=0.025 sample showed the maximum value of about
9.5 × 10-4 W/K2m around 1150 K.
Figure 3 presents the carrier concentration n and theHall mobility μ of In2-xGexO3 at room temperature, as afunction of x. The n value sharply increases with theincrease in x, takes a maximum of 3.2 × 1020 cm-3 at
x=0.025, and then slightly decreases with increasing x.
The Hall mobility μ decreased sharply with increasing x
up to x=0.025 due to the Ge substitution for the In site
and then gradually decreased due to the increase of theFig. 2. Temperature dependence of the (a) electricalresistivity , (b) Seebeck coefficient S and (c) power
factor P of In2-xGexO3 (0≤x≤0.3).

Rintaro Abe et al. Trans. Mat. Res. Soc. Japan(=0.45 mW/K2m) of the Berardan’s sample, whichsuggests that the electrical property of our samples isbetter. However, the thermal conductivity  of oursample (=52 mW/cmK) is about three times larger thanthat (=16 mW/cmK) of the Berardan’s sample. The
reduction of the  value of our samples is a possible
solution to enhance the ZT value.

Fig. 3. Carrier concentration n and Hall mobility μ of
In2-xGexO3 at room temperature, as a function of thecontents x of Ge.

Fig. 4 Temperature dependence of the (a) thermal
conductivity  and (b) of dimensionless figure of merit
ZT of the In2-xGexO3 system.

Fig. 5 the ZT and P values at 1000 K, as a function ofcontents of Ge, x in In2-xGexO3.

41[1] 101-108 (2016)

3.2 In1.9M0.1O3 (M=Si, Ti, Ge, Sn, Te, Ce)
In the previous subsection, the P and ZT values tooka maximum at x=0.025 for the In2-xGexO3 system. On theother hand, these values by Berardan et al. slightly
increased with increasing x at each temperature [5].
Since the solubility limit and the particle size of the
impurity phase seem to change depending on the doped
element, the thermoelectric properties of In1.9M0.1O3
doped with various tetravalent M ions were investigated
for the x=0.1 impurity concentration, which is higher
than the solubility limit of Ge in In2O3 in this study.
Figure 6 shows the XRD patterns of In1.9M0.1O3(M=Si, Ti, Ge, Sn, Te, Ce). The impurity peak of the
In4Si5O12, In2TiO5, In2Ge2O7, SnO2 and CeO2 weredetected for the M=Si, Ti, Ge, Sn, Ce samples,
respectively, which were also confirmed by thequantitative analysis by EPMA shown in the later. In theIn1.9Te0.1O3 sample, there was no impurity peaks.
Figures 7(a) and 7(b) present the SEM images andthe characteristic X-ray images of each doped element in
the In1.9M0.1O3 samples, respectively. The distribution of
the impurity phases depends on the substituted M ions;in the In1.9Ge0.1O3 sample, the In2Ge2O7 particles aslarge as about 50 μm were precipitated, and the grain
size of the impurity phase in the M=Sn, Ce samples was
about 10~30 μm. On the other hand, the impurity phasein the M=Si, Ti, Te samples was finely and uniformly
distributed.
Figure 8 presents the temperature dependence of theρ, S and P values of the In1.9M0.1O3 samples. Comparedto the pure In2O3, the ρ and S values of the M=Ge andSn samples drastically decreased due to the carrier
doping. The ρ and S values of the M=Si and Ti samples
decreased moderately. The relationship between the ρ
and S values and the particle size of the impurity phaseshown in Fig. 7 does not necessarily seem to correlatewith each other, which is remained for further study. In
Fig. 8(c), the power factor P of the Ge-doped sample
shows the maximum value of 7.3 × 10-4 W/K2m at 1148
K. in the In1.9M0.1O3 system.
Table I summarizes the carrier concentration n,electrical conductivity  and Hall mobility  of theIn1.9M0.1O3 samples at room temperature.  was
calculated using the relation of =en. The Ge- andSn-doped samples showed higher carrier concentration nand, at the same time, relatively higher Hall mobility ,which generally decreases with increasing n value due tothe increase of the scattering by carriers. The size andthe distribution of impurity particles also might
influence on the electrical conductivity, in which thecarriers are scattered by grain boundaries and/or thelattice strain due to the impurity particles.

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Substitution effect of tetravalent and pentavalent elements on thermoelectric properties in In2O3-SnO2 system

Fig. 6. XRD patterns of In1.9M0.1O3 (M=Si, Ti, Ge, Sn,
Te, Ce).

Fig. 8. Temperature dependence of ρ, S and P values of
the In1.9M0.1O3 samples.

Fig. 7. (a) SEM images and (b) characteristic X-ray
images of each doped element for the In1.9M0.1O3
samples.
We investigated the influence of the substituted M
ion in the In1.9M0.1O3 system on the thermoelectric andcrystallographic properties. Figure 9(a) shows therelationship between the ionic radius of the impurity M
ion with coordination number (CN) of VI [16] and thelattice parameter of the In1.9M0.1O3 samples. Indium
oxide has a Bixbyite-type cubic lattice structure with
Ia3-(206) space group [17]. The lattice parameter seems
to be correlated with the ionic radius of M ion, whichsuggests the substitution of the M ion for the In-site,
except for the M=Te and Ce samples. For the M=Te andCe samples, the Rietveld refinement might not be
successful, since the solubility may be too small.

Table I. Carrier concentration n, electrical resistivity and Hall mobility  of the In1.9M0.1O3 samples.
Carrier
ElectricalHallconcentration conductivity mobilityn (cm-3)
 (S/cm)
 (cm2/Vs)
18
pure
1.3×1029
157
M=Si
6.7×1018
22
24
13
0.8
M=Ti
1.2×1020
M=Ge
3.1×1020
1780
43
116
4.9
M=Sn
1.8×1020
M=Te
1.1×1019
3
1.7
0.4
0.5
M=Ce
6.0×1018
Figure 9(b) presents the relationship between the
carrier concentration n and the Hall mobility  of theIn1.9M0.1O3 samples. The M=Ge and Sn samples show
the relatively higher n values. It should be noted that the
Ge-doped sample presents the higher Hall mobility,
though the carrier doping introduces the crystal defects
in general. Figure 9(c) depicts the relationship between
the Seebeck coefficient S and the power factor P at 1200
K of the In1.9M0.1O3 samples. Due to the introduction of
the carriers shown in Fig. 9(b), the M ion dependence of
the S value at 1200 K is resemble with that of the nvalue. Seebeck coefficient S of a conductor is given by

Rintaro Abe et al. Trans. Mat. Res. Soc. Japan 2 k k T  (E)S   B  B
3  e   E

(1)where E is the electron energy. Roughly speaking, |S| isinversely proportional to the =en value. In theIn1.9M0.1O3 system, the Ge-substituted sample shows thehighest power factor P, which results from the relatively
higher  value due to both higher n and  values, andthe relatively not small |S| value. These properties areclear contrast with other substituted ions for the In-site.

41[1] 101-108 (2016)

Figure 10(b) depicts the temperature dependence ofdimensionless figure of merit ZT. The ZT value
increases with increasing T. The In1.9M0.1O3 samples
doped by Ge, Sn and Si shows the higher ZT values thanthat of the pure In2O3 sample. The maximum ZT value
was about 0.13 around 1000 K for the Si-, Ge- andSn-doped samples. The optimization of x for eachIn2-xMxO3 system may enhance the P and ZT values
further as indicated for the Ge diped sample shown in
Fig. 5.

Fig. 10. Temperature dependence of the (a) thermal
conductivity κ and (b) dimensionless figure of merit ZTof the In1.9M0.1O3 (M=Si, Ti, Ge, Sn).

Fig. 9. The relations (a) between the lattice parameterand the ionic radius of the substituted M ion and (b)between the carrier concentration n and the Hall
mobility  and (c) between the Seebeck coefficient S
and power factor P at 1200 K for each doped M ion in
In0.9M0.1O3 system.
Figure 10(a) shows the temperature dependence ofthe thermal conductivity κ of the In1.9M0.1O3 (M=Si, Ti,
Ge, Sn) samples. κ at 300 K was decreased by thesubstitution of the M ion, compared to the pure In2O3.
(300 K) of the Si-doped sample was most strongly
depressed in the samples. (T) of the In1.9M0.1O3
samples is less temperature dependence, which may
mainly come from the phonon scattering by impurity
phases and particles identified by XRD and EMPA
shown in Figs. 6 and 7.

3.3 In2O3-SnO2 system
In order to reduce the amount of the In metal, weinvestigated thermoelectric properties of the In2O3-SnO2
system. SnO2 shows similar physical properties to In2O3,
which is inexpensive, plentiful and nontoxic, as a
substitute for the expensive In2O3. Figure 11 shows the
XRD patterns of the In2O3-SnO2 samples with various
mixing ratios. It was found that all the samples were themixture of SnO2 and In2O3 phases without otherimpurity phase.
Figure 12(a) shows the electrical resistivity  andSeebeck coefficient S at 1173 K for the In2O3-SnO2
system, as a function of the fraction of SnO2. As shownin the subsection of 3.2, the  value was decreased by
two orders of magnitude by the 5% SnO2 mixing in theIn1.9Sn0.1O3 sample, which results from the slight Sn4+
substitution for the In3+ site. As a result, electrons were
doped in the sample and, at the same time, the |S| value
was also decreased. And then, the  and |S| values
increased with increasing the contents of the SnO2 phaseand approached those of the pristine SnO2. The power
factor P at 1173 K for the system, as a function of thefraction of SnO2, is shown in Fig. 12(b). The P value
took a maximum by the 5% SnO2 mixture and thendecreased with the increase in the fraction of SnO2.

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Substitution effect of tetravalent and pentavalent elements on thermoelectric properties in In2O3-SnO2 system

Fig. 11. XRD patterns of the In2O3-SnO2.samples withvarious mixing ratios fabricated by SPS method.

Fig. 13. Temperature dependence of the resistivity ρ andSeebeck coefficient S of In2-xSbxO3 ((a), (b)) and of
Sn0.985-xSbxZn0.015O2 ((c), (d)).

Fig. 12. (a) The resistivity  and Seebeck coefficient S
and (b) power factor P of the In2O3-SnO2 samples at1173 K, as a function of the fraction of the SnO2.
3.4 Sb doping to the In2O3-SnO2 system
In order to introduce the carriers to the In2O3-SnO2
system, pentavalent antimony (Sb5+) was doped, whichis expected to substitute both for In3+- and Sn4+-sites.
First, we investigated the Sb doping effect to the In2O3.
Figures 13(a) and 13(b) show the electrical resistivity 
and Seebeck coefficient S of the In2-xSbxO3 (x=0.1),
respectively. The  and |S| values decreased with the Sb
substitution, which suggests that Sb acts as a donor.
Next, we investigated the Sb doping effect to the SnO2.
Figures 13(c) and 13(d) show the  and S values of theSn0.985-xSbxZn0.015O2, respectively. Since many cracks
were created and the mass density was relatively low for
the Sn1-xSbxO2, a small amount of ZnO (y=0.015) was
added as an auxiliary agent, which is a conventional
technique to fabricate ceramics [18]. As a result, theSnO2 samples as high density as 90% of the ideal one
without crack can be fabricated by the Zn addition. The
 and |S| values of the Sn0.985-xSbxZn0.015O2 systemdecreased with increasing Sb substitution, whichsuggests that electrons were doped by the Sb substitution
in the SnO2 system.
The Sb doping was investigated for the In2O3-SnO2
system. Figure 14(a) shows the XRD patterns of(Sn0.5In0.5)1-x-ySbxZnyO2 as a function of the Sb contents
x, in which the ratio of the In2O3:SnO2=1:1 was fixed asa typical example. The contents of ZnO were fixed to be
y=0.02, which was slightly larger than y=0.015 shown in
Figs. 13(c) and 13(d). The difference in the contents ofauxiliary agent y was confirmed not to influence on theelectrical properties. In Fig. 14(a), only the peaks of
SnO2 and In2O3 phases can be indexed without impurity,even though the Sb was doped up to x=0.06. Figures
14(b) and 14(c) show the magnification of the (110)
XRD peak of SnO2 phase and the (222) peak of In2O3
phase for each sample, respectively. It should be notedthat the (222) peak of In2O3 shifts to higher angle withincreasing the contents of Sb, but the (110) peak of SnO2
hardly shifts. The ionic radius of In3+, Sn4+ and Sb5+ is
0.80Å, 0.69Å and 0.60Å, respectively [16]. These
results suggest that the Sb atom was preferentiallysubstituted for the In-site rather than the Sn-site.
Because the peak shift of the (222) peak of In2O3 tohigher angle results from the decrease of lattice
parameter. On the other hand, unchanged (110) peak of
the SnO2 phase suggested that the most of the doped Sb
substituted for the In-site.
Figure 15 shows the temperature dependence of ρ, S
and P values of the (Sn0.5In0.5)1-x-ySbxZnyO2 system. The
lowest  value was realized for the Sn0.5In0.5O3 sample
because of the Sn substitution for the In-site. As a result,the |S| value also decreased and the P value was not sohigh. Since the Sb atoms were expected to substitutemainly for the In-site as discussed in Fig. 14, the  value

Rintaro Abe et al. Trans. Mat. Res. Soc. Japan

41[1] 101-108 (2016)

slightly decreases with increasing Sb content. However,the P value of the substituted samples does not exceedthat of In2O3.
Figures 16(a) and 16(b) depict the temperaturedependence of the κ and ZT values of the typical(Sn0.5In0.5)1-x-ySbxZnyO2 samples, respectively. The κ
value was fairly reduced at low temperatures by mixing
the SnO2, Sb and Zn into the parent In2O3 material. On
the other hand, at higher temperatures, thephonon-phonon scattering for the pure In2O3 sample isdominant, compared to the impurity scattering in theother (Sn0.5In0.5)1-x-ySbxZnyO2 samples. The ZT of(Sn0.5In0.5)1-x-ySbxZnyO2 could not exceed that of the pureIn2O3.

Fig. 15. Temperature dependence of (a) resistivity ρ, (b)Seebeck coefficient S and (c) power factor P of
(Sn0.5In0.5)1-x-ySbxZnyO2 system.

Fig. 14. (a) XRD patterns of the (Sn0.5In0.5)1-x-ySbxZnyO2.,
Enlarged views around the (b) (110) SnO2 peak (2~
26.6°) and (c) (222) In2O3 peak (2~ 30.6°).
4. SUMMARY
We have investigated the thermoelectric propertiesof the In2-xMxO3 system doped with tetravalent elements(M=Si, Ti, Ge, Sn, Te, Ce) and those of the In2O3-SnO2
system doped with pentavalent Sb element.
For the In2-xMxO3 system, the maximum power
factor P=S2/ and the maximum dimensionless figure of
merit ZT were realized for the M=Ge sample, whichsuggest that the Ge is the best dopant. These results
qualitatively reproduced those reported by Berardan etal.
For the In2O3-SnO2 system doped with Sb5+, which
is proposed to reduce the amount of the expensive In, Sb
atom was preferentially substituted for the In-site rather
than the Sn-site. The P value of the samples does notexceed that of the pure In2O3, contrary to our
expectation. Other substitution element and/or mixed
system must be explored to enhance the thermoelectricperformance of the In2O3 system.

Fig. 16. Temperature dependence of (a) thermal
conductivity κ, (b) dimensionless figure of merit ZT of(Sn0.5In0.5)1-x-ySbxZnyO2 system.

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Acknowledgments
The authors thank Professor Takashi Goto and Dr.
Hirokazu Katsui of Tohoku University, Japan, for theassistance in thermal conductivity measurement using
laser flash method and for valuable discussion. This
work was performed under the Inter-university
Cooperative Research Program of the Institute forMaterials Research, Tohoku University.
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(Received January 21, 2015; Accepted August 27, 2015)