Compatibility of GdxTi2O7 pyrochlores (1.72≤x≤2.0) as electrolytes in high-temperature solid oxide fuel cells

Solid State Ionics 158 (2003) 79 – 90 www.elsevier.com/locate/ssi

Compatibility of GdxTi2O7 pyrochlores (1.72VxV2.0) as electrolytes in high-temperature solid oxide fuel cells Masashi Mori a,*, Geoff M. Tompsett b, Nigel M. Sammes c, Eisaku Suda d, Yasuo Takeda d a

Central Research Institute of Electric Power Industry, 2-6-1 Nagasaka, Yokosuka, Kanagawa 240-01, Japan b Department of Chemical Engineering, The University of Massachusetts, Amherst, MA 01003, USA c Department of Mechanical Engineering, University of Connecticut, Storrs, CT 06269-3139, USA d Mie University, 1515 Kamihama, Tsu 514-8507, Japan Received 27 March 2002; received in revised form 24 August 2002; accepted 7 September 2002

Abstract The crystal structure, sintering characteristics, electrical conductivity and thermal expansion of GdxTi2O7
d pyrochlores (1.72 V x V 2.10) have been studied as potential electrolytes in high-temperature solid oxide fuel cells (SOFC). It was found that the A-site deficiency of GdxTi2O7
d with cubic symmetry showed a wide region of 1.72 < x < 2.03. A-site vacancies promoted the sintering characteristics of these materials. Electrical conductivity of the pyrochlores with A-site vacancies was lower than that of GdxTi2O7 in air although the reverse was observed in reducing atmospheres, due to the appearance of electronic conduction. Average linear thermal expansion coefficients of GdxTi2O7
d pyrochlores decreased with increasing A-site vacancies. The difference between the TEC in air and in a H2 atmosphere had a tendency to increase with increasing A-site vacancies. D 2002 Elsevier Science B.V. All rights reserved. Keywords: Gadolinium titanate; Crystal structure; Electrical conductivity; Non-stoichiometry; Sintering characteristics; Thermal expansion

1. Introduction High-temperature solid oxide fuel cells (SOFC) utilize Y2O3 stabilized ZrO2 (YSZ) as an electrolyte, because of the pure oxide-ion (O2
) conductivity, as well as excellent chemical and physical stability under reducing and oxidizing atmospheres at an operating temperature of approximately 1000 jC. However, the

* Corresponding author. Tel.: +81-468-56-2121; fax: +81-46856-3346. E-mail address: [email protected] (M. Mori).

oxide-ion conductivity of YSZ is approximately 10
1 S/cm at 1000 jC, which is not high enough. This conductivity value is 2 to 4 orders of magnitude lower than that of other cell component materials. An electrolyte with high oxide-ion conductivity is required so that voltage losses due to the internal resistance in the SOFC are not apparent. YSZ has a fluorite-type structure with cubic symmetry (space group Fm3m). In the fluorite structure, the anions are in simple cubic packing with half the interstices filled by cations. This gives rise to a unit cell in which the cations are characterized by facecentered cubic packing with a space at the center of

0167-2738/02/$ - see front matter D 2002 Elsevier Science B.V. All rights reserved. PII: S 0 1 6 7 - 2 7 3 8 ( 0 2 ) 0 0 7 6 1 - 0

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M. Mori et al. / Solid State Ionics 158 (2003) 79–90

the unit cell, corresponding to the unfilled interstice in the simple cubic anion lattice. This structure shows a loose packing of oxide-ions, and thus conduction of the oxide-ion at high temperature. Since the oxide-ion is transported via hopping through the vacancy site, the high oxygen vacancy concentration gives rise to a high oxide-ion mobility. Because the vacancies on the O-site are formed by substituting a lower-valence ion on the Zr4 + ion site, the concentration of the vacancy is determined by the concentration of the dopant. However, isothermal plots of the conductivity, as a function of the dopant concentration, exhibit a maximum [1]. In the case of Y2O3 dopant, the conductivity increases with increasing Y2O3 content up to 9 mol% and then decreases. The decrease in conductivity, at higher dopant concentration, is believed to be due to vacancy clustering or electrostatic interaction [2]. The pyrochlore structure can be considered as an ordered vacancy fluorite structure and has the ideal composition of A2B2O7 with space group Fd3m. Compared to the fluorite structure, two kinds of cations are ordered into rows in (110) directions, resulting in the axes of the pyrochlore unit cell being twice that of the fluorite unit cell. Two kinds of polyhedra, AO8 and BO6 exist in the pyrochlore structure and there are two unoccupied O-sites (8(b) site) in the BO6. Thus, there is a possibility that the oxide-ion can be transported via a hopping mechanism through the vacancy sites, namely the 8(b) sites. Additionally, some of the pyrochlore oxides have been reported to be stable in a wide composition range, such as 0.33 V x V 0.60 for the (1
x)ZrO2 –xGdO1.5 system [3]. This means that the pyrochlores show a non-stoichiometry of the A-site, B-site or O-site, and more oxygen vacancies are formed by charge-compensation [4 – 7]. It is quite interesting to discuss the relationship between non-stoichiometry and electrical property of the pyrochlores. Since in Ln2Ti2O7 there are no d-electrons (Ti4 +: 0 d system), and the f-electrons, due to the Ln-ions, are localized in the inner 4f levels, no electronic conductivity is expected at room temperature; Ln2Ti2 O7 are found to be insulating with V 10
12 S/cm at approximately room temperature. On the other hand, p-type semiconduction has been noted by Brixner [8] in these compounds at temperatures V 1000 jC, and Uematsu et al. [9] have also predicted that oxide-ion

conduction may be exhibited with increasing temperature in some titanates. Recently, Tuller et al. [10 – 12] reported oxide-ionic conductivity in Gd2Ti2O7 pyrochlore, showing approximately 6.3  10
4 S/cm at 1000 jC in air, and that A- and B-site doping improves the electrical property of stoichiometric Gd2Ti2O7. In addition, it was found that the doping increases the degree of non-stoichiometry of the pyrochlores [13 – 15] although their details were not shown. The electrolyte in the SOFC should meet the following requirements in addition to low electronic and high oxide-ion conductivity; (1) excellent chemical stability under reducing and oxidizing atmospheres, (2) good sintering characteristics, (3) good compatibility of thermal expansion with the other cell components, and (4) low reactivity with electrodes during SOFC fabrication and operation Although there are many fundamental data for crystallographic properties of Ln2Ti2O7 pyrochlores with a stoichiometric composition [16,17], few papers for non-stoichiometric Gd2Ti2O7
d have been published on the relationship among crystal structures, sintering characteristics, electrical conductivity and thermal expansions of the pyrochlores. In this paper, we measured the required properties, and discussed the possibility for the compatibility of the Gd2Ti2O7-based pyrochlores as electrolytes in the SOFC.

2. Experimental The gadolinium titanate powders were synthesized by a solid-state reaction technique. Starting material powders, Gd2O3 (pre-heated at 1500 jC for 1 h) (99.9%, High Purity Chem., Japan), and TiO2 (preheated at 1400 jC for 1 h) (99.9%, rutile phase, High Purity Chem.) were used without further purification. They were weighed and mixed in a rotary-type Y2O3 partially stabilized ZrO2 ball mill for 24 h with isopropyl alcohol. After drying, the mixtures were heated to 1200 jC in air and held at this temperature for 6 h. The powders were uniaxially pressed into a tablet under pressure of 98 MPa, and the tablets were isostatically pressed at a pressure o f 200 MPa. The tablets were then sintered at 1540 – 1560 jC for 12 h in air, at a heating/cooling rate of 6 jC/min. The milling and heating procedures were repeated twice.

M. Mori et al. / Solid State Ionics 158 (2003) 79–90

XRD patterns of powdered samples were obtained on a Philips X’Pert MPD diffractometer (12 kW) using monochromated CuKa radiation and a scintillation detector, at a scanning rate of 0.02 j/s. Raman spectroscopy was performed on samples of powders using a Jobin Yvon U1000 double beam pass spectrometer equipped with a microscope stage for analyzing small samples utilizing 180 j incident geometry. A Spectra Physics argon-ion laser was employed to excite laser Raman Spectra using a 514 nm laser line at an incident power of ca. 4 – 5 mW, and a water-cooled photomultiplier tube. Spectra were obtained using 500 Am slit width and scanning rate used to collect the spectra was kept at 0.5 cm s
1. Sintering characteristics of the pyrochlores were evaluated as follows. The heating rate was fixed at 200 jC/h on every run, and the holding time, at selected temperatures, was changed from no holding to 10 h. The density of the sintered tablets was determined from observed values of size and weight of the tablets. Relative density was derived using the theoretical value determined from the experimental lattice parameters and unit formula. Electrical conductivity for the sintered samples was measured as a function of temperature by the fourterminal method. After the Pt-terminals were attached using Pt-past, they were fired at 1000 jC for 1 h. Measurements were performed stepwise in air between 600 and 1000 jC. O2, air, N2 and H2 gas mixtures of the desired ratio, prepared in the gas mixing system, were introduced into the equipment after it had been bubbled through pure water at 30 jC. TEC was measured using a Mac Science TD5000S dilatometer equipment, at a temperature range from 50 to 1000 jC in air and in a H2 atmosphere. References used for the present measurements were a rod of fused SiO2 for air and sapphire for H2. The measurements were performed using a heating/cooling rate of 2 jC/min. A flow rate of 50 ml/min of H2 gas was used during these measurements after it had been bubbled through pure water at 10 jC. Thermogravimetry and differential thermal analysis (TG-DTA) measurements were carried out using a Mac Science TG-DTA 5000S equipment in the temperature range from room temperature to 1000 jC under the H2 atmosphere at a heating/cooling rate of 5 jC/min.

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3. Results and discussion 3.1. Crystal structure 3.1.1. XRD analysis Fig. 1(a) and (b) shows the XRD patterns of the GdxTi2O7
d samples (2h = 10 –40 j) after heating at 1540 jC for 12 h in air. In Fig. 1(b), the open circle represents a peak of Gd2TiO5 and the closed one represents a peak of TiO2. Gd2Ti2O7 pyrochlore crys˚ , V = 1058.2(1) tallized in a cubic unit a = 10.1905(4) A ˚ 3 and Z = 8 (Fd3m) which is in good agreement with A that previously reported [18]. The XRD patterns showed that no compound, other than the pyrochlore phase, was found in the samples with a Gd content of 1.72 < x < 2.03. On the other hand, the samples contained a very small amount of TiO2 for Gd1.72Ti2O7
d, and Gd2TiO5 for the Gd2.03Ti2O7
d in addition to the pyrochlore phase. Fig. 2 shows the lattice parameters of the GdxTi2 O7
d samples after heating at 1540 jC for 12 h in air as a function of Gd content. The closed circles represent the samples with a second phase. The lattice parameters of GdxTi2O7
d increased with increasing Gd content and reached a constant value in the region of x z 2.0. The non-stoichiometric range of A-site vacancies in the pyrochlores is considered to be 1.72 < x V 2.0. However, it was observed that the standard deviation of lattice parameters has a tendency to decrease with increasing A-site vacancies. This leads to an expectation that the formation of A-site and oxygen vacancies might cause some distortion in the cubic pyrochlore structure. In the pyrochlore structure, the Gd3 + ion is located on the 16(c) site, the Ti4 + ion on the 16(d) site, and the Gd3 + ion is 8-coordinated by oxide-ions, while the Ti4 + ion is 6-coordinated. This results in three kinds of tetrahedral interstices for anions; the 48(f) site has two nearly neighbor A-site and B-site ions. There are two similar sites for the 8(a) and 8(b); the 8(a) site has four B-site ions, and the 8(b) site has four A-site ions. In the Gd2Ti2O7 pyrochlore structure, the 8(b) sites are vacant. The non-stoichiometry of the A-site ion in the pyrochlore structure results in charge-compensation by the release of oxygen, which should be 8-coordinated. This can be explained by the fact that the content of the oxygen vacancy in Gd1.72Ti2O6.61 is almost comparable to that of the vacancy on the 8(a) site.

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M. Mori et al. / Solid State Ionics 158 (2003) 79–90

Vanderborre and Husson [19] described the following irreducible representation for the Raman and infrared modes predicted. Gopt ¼ A1g ðRÞ þ Eg ðRÞ þ 2 F1g ðiÞ þ 4 F2g ðRÞ þ 3 A2u ðiÞ þ 3 Eu ðiÞ þ 7 F1u ðirÞ þ F2u ðiÞ Gac ¼ F1u where R = Raman active, ir = infrared active, I = inactive. Gopt = optical modes, Gac = acoustic modes. There are a total of seven Raman active modes and seven infrared active modes. Experimentally, Vanderborre and Husson [19] observed (and assigned) Raman bands 110, 215, 227 (F2g), 317 (F2g), 515 (A1g, F2g), 580 (F2g) and 705 cm
1, while a band at 347 (Eg) cm
1 was calculated but not observed. Oueslati et al. [20] investigated the Raman spectroscopy and structural disorder of Gd2(ZrxTi1
x)2O7. These workers observed a band broadening with increase in x, which in turn was related to the increase in structural disorder of the material. Oueslati et al. [20] disputed the mode assignments of Vandenborre et al. [21] of Raman bands observed for Gd2Ti2O7. The band at 310 cm
1 previously attributed to the O –Ti– O bending mode, did not shift in wavenumber posi-

Fig. 1. (a) XRD patterns of the GdxTi2O7
d pyrochlore samples (2h = 10 – 40j) after heating at 1540 jC for 12 h in air. (b) XRD patterns of the Gd2.03Ti2O7
d and GdxTi2O7
d samples (2h = 10 – 40j) after heating at 1540jC for 12 h in air. The open circle represents a peak of Gd2TiO5, and the closed one represents a peak of TiO2.

3.1.2. Raman spectroscopy analysis Previously, Vanderborre and Husson [19] and Oueslati et al. [20] have studied the Gd2Ti2O7 pyrochlore using Raman spectroscopy. However, the non-stoichiometry of this material has not been studied. For the cubic pyrochlore structure with Fd3m space group,

Fig. 2. Lattice parameters of the GdxTi2O7
d pyrochlores after heating at 1540 jC for 12 h as a function of Gd content in the Gdx Ti2O7
d system. The closed circles represent the samples with second phase.

M. Mori et al. / Solid State Ionics 158 (2003) 79–90

tion when Ti was substituted by Zr, and was therefore re-assigned to the O – Gd – O bending mode. Similarly, the band observed at 520 cm
1 was assigned to a Gd – O stretching vibration. The band at 450 cm
1 was assigned to the Ti –O stretching vibration. Finally, broad weak bands observed at 765 – 775 cm
1 were assigned by Oueslati et al. [20] to forbidden modes activated by the solid-solutions Gd2(ZrxTi1
x)O7. Fig. 3 shows the comparison of Raman spectra for GdxTi2O7
d (x = 2.03, 2.0, 1.9, 1.8 and 1.72) and the observed band positions are listed in Table 1. In general, there is good agreement between the observed spectrum for Gd2Ti2O7 with that observed in the literature by Vanderborre and Husson [19] and Oueslati et al. [20]. However, an extra band is observed at ca. 125 cm
1 and no band at 379 cm
1 is observed (as reported by Vanderborre and Husson [19] assigned to the Eg mode). The band at 125 cm
1 is attributed to

Fig. 3. Comparison of Raman spectra for the Gd2
xTi2O7
d pyrochlores, (a) x = 2.03, (b) x = 2.0, (c) x = 1.9, (d) x = 1.8 and (e) x = 1.72.

83

the presence of trace Gd2O3. The literature Raman band positions and assignments are summarized in Table 1 and the assignment of the spectra in this work is made according to that by Oueslati et al. [20]. It can be seen from the spectra in Fig. 3 that there is, in general, little change in the Raman profile with shift in A-site vacancy in the GdxTi2O7 pyrochlore system. However, there are several significant features, namely, the increase in bands at ca. 103, 126 and 582 cm
1. Fig. 4 shows the comparison of the Raman spectra of (a) Gd2Ti2O7, (b) Gd1.72Ti2O7
d and the subtraction spectrum of the two. The subtraction spectrum shows significant features at 71, 102, 120, 334, 440, 470, 518, 584, 701 and 739 cm
1. Several of these bands indicate the presence of either of the starting materials, namely TiO2 [22] as rutile (440 cm
1) or Gd2O3 [23] (bands at 71, 102 and 120 cm
1) since they match the band positions listed in Table 1 [24]. TiO2 was also observed by the XRD analysis in trace amounts. The feature at 332 cm
1 shows that the band at 311 cm
1 is considerably broadened towards higher frequency, with increase in A-site vacancy from 0.0 to 0.2. This is consistent with increase in disorder in the system as observed by Oueslati et al. [20] for the same band with increasing x in Gd2(ZrxTi1
x)O7. The feature at 332 is most likely due to the 311 cm
1 band broadening since neither TiO2 or Gd2O3 have bands at this position. The XRD results showed a shift in the cubic pattern to higher 2h values. With increasing the stoichiometry of Gd in Gdx Ti2O7
d there is no apparent change in the spectra as shown in Fig. 5(b) compared to (a). However, the subtraction of the Raman spectrum of Gd2Ti2O7 with Gd2.03Ti2O7
d shows several weak bands at 106, 128, 322, 521, 595, 681 and 799 cm
1. The bands at 322 and 521 cm
1 appear to be due to incomplete subtraction, and there are no bands assigned to the Gd2TiO5 secondary phase, which was observed in the XRD pattern in trace amounts. The other features observed are assigned to vibrations related to the distortion of the pyrochlore structure. The full width at half maximum (FWHM) of the intense band at 311 cm
1 assigned to O – Gd – O bending mode (Eg) was plotted against x in GdxTi2 O7
d as shown in Fig. 6. It can be seen that with decreasing x in GdxTi2O7
d, i.e. decreasing Gd content, there is an increase in the band width before x = 1.9. At dopant concentrations of greater than

84 Table 1 Raman band positions (cm
1) and assignments for stoichiometric and non-stoichiometric GdxTi2O7
y pyrochlores compared with literature values for base oxides and pyrochlores TiO2 (rutile) [22]

247 two phonon

Gd1.72Ti2O7 [this work]

Gd1.8Ti2O7 [this work]

Gd1.9Ti2O7 [this work]

71 s (Bg) 84 mw (Ag) 98 vs (Bg) 109 vs (Ag)

72 w

68 w

61 w 70 w

104 m,b

106 m

129 w

115 w (Bg) 123a (Ag) f 150 vw (Ag) 176 w (Ag) 218 w (Ag) 256 mw (Ag) 269 m (Ag) 299 m (Bg)

Gd2Ti2O7 [this work]

Gd2.03Ti2O7 [this work]

Gd2Ti2O7 [21]

Gd2Ti2O7 [20]

La2Zr2O7 [24]

Mode vibration type [20]

102 m

105 w

103 w

110 ( – )

105

n.o. (calc.238/246)

F2g (Eg/A1g)

129 w

127 w

125 sh

126 w

214 m,b

219 m,b

208 m 221 w

222 sh

217 m

215 ( – ) 227 (F2g)

205

298

F2g

314 vs

313 vs

312 vs

311 vs

311 vs

317 (F2g)

310

395

379 sh

347b (Eg)

450

490

Eg O – Gd – O bend A1g Ti – O stretch

515/515 (F2g/A1g) 580 (F2g)

520

523

580 640

590

387 m (Ag) 447 Eg

612 A1g

826 B2g

416 s (Bg) 442 m (Bg + Ag) 483 m (Bg) f 580 sh (Ag) 590 w (Ag)

432 vw 440 w 465 vw

448 w

450 w

452 w

451 w

518 s

519 s

517 s

518 s

517 s

584 w,b 698 w,b 739 sh,b

582 w 695 w,b 740 sh

578 sh 682 w 783 vw

585 sh

582 sh 681 m 798 w

846 vw,b 924 vw

846 vw

825 vw

705 ( – )

853 vw

vw = Very weak, w = weak, mw = medium weak, m = medium, s = strong, vs = very strong and b = broad relative intensity. a Observed at 34 K. b Calculated not observed.

743

F2g Gd – O stretch F2g LO, forbidden mode

M. Mori et al. / Solid State Ionics 158 (2003) 79–90

143 B1g

Gd2O3 [23]

M. Mori et al. / Solid State Ionics 158 (2003) 79–90

85

at the selected temperatures on the relative density of Gd2Ti2O7 and Gd1.9Ti2O6.85. An acceptable relative density of the electrolyte for SOFC has reported to be 94% [25], and the shaded region in Fig. 7(b) indicated this. The Gd2Ti2O7 showed poor sintering characteristics. The relative density of this material was about 60%, although a slight dependence could be seen on the sintering temperature and holding time. The relative densities of the GdxTi2O7
d pyrochlores had a tendency to increase with increasing A-site deficiency. A relative density of 94% was attained for the Gd1.9Ti2O6.85 after firing at 1500 jC for 10 h. For the pyrochlores with A-site vacancy, the promotion of the densification seems to relate to vacancy formation of Gd ions and oxide-ions. 3.3. Electrical conductivity Fig. 8 shows the temperature dependences of electrical conductivity for the GdxTi2O7
d pyro-

Fig. 4. Comparison of the Raman spectra of (a) Gd2Ti2O7, (b) Gd1.72 Ti2O7 and (c) the subtraction spectrum of the two.

x = 1.9, there is little change in the band width. This is indicative of structural disorder in the pyrochlore system as described by Oueslati et al. [20]. Although details remain uncertain, there is a possibility that from both the Raman and XRD measurements, their results suggest the indication of structural disorder for the pyrochlore with increasing A-site vacancies. 3.2. Sintering characteristic In the SOFC, O2
ions move through the electrolyte and electrochemically react with H + ions without the intermediate form of thermal energy. Since a densely sintered electrolyte is needed without mixing the oxidant and fuel gases, high sintering characteristics of the materials is important for the production of impervious bodies with high density at a low sintering temperature. Fig. 7(a) shows the relative density of GdxTi2O7
d pyrochlores as a function of firing temperature, where the holding time at the highest temperature is zero. Fig. 7(b) shows the effect of holding time

Fig. 5. Comparison of the Raman spectra of (a) Gd2Ti2O7, (b) Gd2.03 Ti2O7 and (c) the subtraction spectrum of the two.

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M. Mori et al. / Solid State Ionics 158 (2003) 79–90

Ti-doping causes the appearance of electronic conduction at high temperatures in low oxygen partial pressures, because of the reduction of Ti4 + into Ti3 + ion [26]. Therefore, the higher electrical conductivity of the pyrochlores is related to the appearance of electronic conductivity at high temperature.

Fig. 6. Full width at half maximum (FWHM) for the 311 cm
1 Raman band as a function of Gd content in the GdxTi2O7
d system.

chlores under various atmospheres. The electrical conductivity of the Gd2Ti2O7 at 1000 jC in air is comparable to that reported previously [10]. The electrical conductivity of the GdxTi2O7
d pyrochlores increased with increasing temperature. Since the electrical conductivity should be due to thermally activated hopping conduction of the oxide-ion for Gd2Ti2O7 in oxidizing atmospheres, it should be represented by the function rT~exp(
E/K B T), where r is the electrical conductivity, T is the absolute temperature, E is the activation energy for electrical conduction and KB is the Boltzmann constant. Since the linearity of the increased conductivity was observed for all the pyrochlore samples, their condition should be due to oxide-ions. For these pyrochlores, the electrical conductivity in the reducing atmosphere is higher than that in the oxidizing one. The color of the samples with A-site deficiency changed from white into black after the H2 measurements. The oxygen vacancies of the pryochlores at 1000 jC in the H2 atmosphere are summarized in Table 2. All the samples showed the weight decrease. This weight loss would correspond to the oxygen vacancies formation by a charge-compensation of Ti4 + to Ti3 + ion. For YSZ with the fluorite structure,

Fig. 7. (a) Relative density of the Gd2Ti2O7
d pyrochlore samples as a function of firing temperature, where the holding time at the highest temperature was zero. (b) Effect of holding time at the selected temperatures (1300, 1400 and 1500 jC) on the relative density of Gd2Ti2O7 and Gd1.9Ti2O6.85.

M. Mori et al. / Solid State Ionics 158 (2003) 79–90

87

Fig. 8. Temperature dependences of electrical conductivity for the GdxTi2O7
d pyrochlores under various PO2. The open circles, open triangles, open squares and closed circles represent the conductivites for PO2 = 1 atm, PO2 = 2  10
1 atm, PO2 = 7  10
6 atm and PO2 = 1  10
18 atm at 1000 jC, respectively. (a) Gd2Ti2O7, (b) Gd1.9Ti2O6.85, (c) Gd1.8Ti2O6.7 and (d) Gd1.72Ti2O6.61.

In order to make clear the effect of A-site vacancy in the pyrochlore structure on electrical conductivity, the conductivity at 1000 jC is plotted as a function of

temperature in Fig. 9. Each data point in this figure corresponds to that of a temperature at 1000 jC, under various atmospheres. The conductivity of GdxTi2O7
d

Table 2 Oxygen vacancies at 1000 jC in a H2 atmosphere, TECs in the temperature range from 50 to 1000 jC in air and in a H2 atmosphere, and lattice parameters and cell volumes after firing at 1000 jC in a H2 atmosphere for the GdxTi2O7
d pyrochlores Sample

Oxygen vacancies d at 1000 jC

TEC(  10
6/jC)

Gd2Ti2O7
d

0.05

Gd1.9Ti2O6.85
d

0.06

Gd1.8Ti2O6.7
d

0.07

Gd1.72Ti2O6.61
d

0.07

10.8 10.8 10.6 10.6 10.6 10.5 10.5 10.4

(in (in (in (in (in (in (in (in

air), 10.8 (in H2 (2nd)) air), 10.5 (in H2 (2nd)) air), 10.4 (in H2 (2nd)) air), 10.3 (in H2 (2nd))

After reducing in H2 at 1000 jC H2 (1st)),

˚ , V = 1059.09 A ˚3 a = 10.1932(3) A

H2 (1st)),

˚ , V = 1056.53 A ˚3 a = 10.1850(3) A

H2 (1st)),

˚ , V = 1051.50 A ˚3 a = 10.1688(5) A

H2 (1st)),

˚ , V = 1039.94 A ˚3 a = 10.1314(5) A

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M. Mori et al. / Solid State Ionics 158 (2003) 79–90

3.4. Thermal expansion property Fig. 10(a) and (b) shows the linear thermal expansion of the GdxTi2O7
d pyrochlores in air and H2 atmosphere in comparison with that of 8 mol% Y2O3 stabilized ZrO2 (8YSZ) fluorite. Because these pyrochlores show a cubic symmetry, no anomaly in the thermal expansion curve for phase transformation in

Fig. 9. Electrical conductivities of the GdxTi2O7
d pyrochlores under various PO2 as a function of Gd content.

pyrochlores was higher than that of Gd2Ti2O7 at low oxygen partial pressures V 7  10
6 atm, although the reverse was observed in high oxygen partial pressures z 0.2 atm. The A-site deficient dependency of the electrical conductivity at 800 and 900 jC showed the same behavior. An important point is that in some cases, the conductivity could not be measured at temperatures V 700 jC, because the GdxTi2O7
d pyrochlores showed quite low electrical conductivity in high oxygen partial pressures z 0.2 atm. In the case of doped zirconia with a fluorite structure, the high oxide-ion conductivity originates from the existence of vacancies on the oxygen site. For the A-site deficient GdxTi2O7
d pyrochlores, the formation of oxygen vacancies occurs by a charge-compensation of A-site deficiency. Therefore, it is expected that the electrical conductivity of the GdxTi2O7
d pyrochlores might be improved by the vacancy formation. However, the reverse result was obtained as seen in Fig. 9.

Fig. 10. Linear thermal expansion of the GdxTi2O7
d pyrochlores in air and H2 atmosphere in comparison with that of 8 mol% Y2O3 stabilized ZrO2 (8YSZ) fluorite. (a) In air, (b) in a H2 atmosphere.

M. Mori et al. / Solid State Ionics 158 (2003) 79–90

air was observed. Although the pyrochlore samples with A-site deficiency showed the oxygen vacancies by oxygen release in the H2 measurement, no isothermal expansion in the H2 atmosphere was observed for the pyrochlores. Fig. 11 shows the TECs of GdxTi2O7
d pyrochlores in the temperature range from 50 to 1000 jC as a function of Gd content. The TEC values, and the lattice parameters, of the samples in the H2 atmosphere are summarized in Table 2. The TEC of the Gd2Ti2O7 was 10.8  10
6/jC and was observed to be constant regardless of the oxygen partial pressure. The TECs of the GdxTi2O7
d pyrochlores decreased with decreasing Gd content. For the pyrochlores with A-site vacancies, a difference between the TECs in air and in the H2 atmosphere was observed and had a tendency to increase with increasing A-site vacancies. For the measurements made in the H2 atmosphere, TECs of the pyrochlores during the first heating cycle were larger than those observed during the second one. It is known that AE-doped lanthanum chromites are isothermally expanded by reduction of B-site ions and formation of oxygen vacancy under reducing atmospheres [27,28]. The reverse result was observed for these materials although they showed oxygen vacan-

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cies in the pyrochlore structure. In general, the TEC of the materials has a tendency to increase with the ionic bond strength. Although details remain uncertain, there is a possibility that the ionic bond strength in the pyrochlores decreases via the formation of oxygen vacancies, and a change from Ti4 + into Ti3 + ions. This can be explained by the fact that the lattice parameters of the pyrochlores are smaller that those after reducing at 1000 jC in the H2 atmosphere as shown in Table 2.

4. Conclusion The crystal structure, sintering characteristics, electrical conductivity and thermal expansion of GdxTi2O7
d pyrochlores (1.72 V x V 2.10) have been measured and discussed as potential electrolytes in the high-temperature solid oxide fuel cells. A non-stoichiometry in GdxTi2O7
d was 1.72 < x < 2.00. Sintering characteristics of the pyrochlores were promoted by A-site deficiency. The pyrochlores showed linear thermal expansion behavior in air and in an H2 atmosphere, and with increasing A-site deficiencies, their TEC values decreased and were close to that of the YSZ electrolyte. The electrical conductivity of these materials was V 1.8  10
4 S/cm2 at 1000 jC in air and its conductivity values were slightly lower than is acceptable for an electrolyte in a SOFC. In addition, the A-site deficiency in the pyrochlores caused an appearance of electronic conduction at high temperature under reducing atmospheres. Thus, this problem remains to be solved. From the viewpoint of these properties, when GdxTi2O7
d pyrochlores are considered as potential electrolytes in the high-temperature SOFC’s, an enhancement of their electrical conduction for oxide-ion and a prevention of their electronic conduction at high temperatures under reducing atmospheres are necessary.

Acknowledgements

Fig. 11. TECs of GdxTi2O7
d pyrochlores in the temperature range from 50 to 1000 jC as a function of Gd-content.

We would like to thank Dr. John Seakins and the Chemistry Department, University of Auckland, New Zealand, for the use of the Raman spectrometer. This work was supported by the New Zealand Foundation of Research, Science and Technology, Contract UOW 804.

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