Proton conductivity in Sm2Sn2O7 pyrochlores

Solid State Ionics 181 (2010) 1577–1585

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Solid State Ionics j o u r n a l h o m e p a g e : w w w. e l s ev i e r. c o m / l o c a t e / s s i

Proton conductivity in Sm2Sn2O7 pyrochlores K.E.J. Eurenius, E. Ahlberg, C.S. Knee ⁎ Department of Chemistry, University of Gothenburg, SE-412 96 Göteborg, Sweden

a r t i c l e

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Article history: Received 17 October 2009 Received in revised form 20 August 2010 Accepted 3 September 2010 Keywords: Sm2Sn2O7 Proton conductor Pyrochlore Oxide ion conductivity p-type conductivity Infra-red spectroscopy Thermogravimetric analysis

a b s t r a c t The electrical conductivity of the pyrochlore systems, Sm2Sn2O7, Sm1.92Ca0.08Sn2O7 − δ and Sm2Sn1.92Y0.08O7 − δ was studied using impedance spectroscopy under wet and dry gas (O2 and Ar) in the temperature range 150– 1000 °C. Enhancements of the bulk conductivity of all samples at temperatures up to ~550 °C were observed for wet conditions consistent with significant levels of proton conduction. The presence of dissolved protons in the acceptor-doped materials, Sm1.92Ca0.08Sn2O7 − δ and Sm2Sn1.92Y0.08O7 − δ, is supported by infrared spectroscopy and thermogravimetric analysis. Proton conduction was confirmed by isotope effects under heavy water (O2/D2O and Ar/D2O). The A-site substituted sample Sm1.92Ca0.08Sn2O7 − δ yielded the highest levels of proton conduction and displayed mixed ionic and electronic conduction under dry oxidising conditions. Electron hole conduction dominates in dry oxygen for Sm2Sn1.92Y0.08O7 − δ and Sm2Sn2O7. For the A-site doped sample bulk and grain boundary conduction could be separated. The specific grain boundary conduction was calculated using the brick layer model and was found to be two orders of magnitude lower compared to the bulk conductivity. The unexpected increase in conductivity seen for the undoped sample under wet gas is discussed in the context of structural disorder and possible filling of the un-occupied anion site in the pyrochlore structure by OH-groups. © 2010 Elsevier B.V. All rights reserved.

1. Introduction Perovskite systems have been the centre of attention regarding the development of oxide based proton conducting electrolytes for many years [1–4]. Acceptor-doped variants of BaCeO3 [5,6] and BaZrO3 [7,8] are examples of simple phases which offer reasonably high proton conductivities, ~ 10− 3 S cm− 1, in the attractive intermediate temperature range. In spite of a large number of studies relating to high temperature oxide ion migration in A2B2O7 pyrochlore systems, see for example [9–13], there are relatively few studies relating to proton conduction at lower temperatures. Further investigation of proton mobility in pyrochlore oxides that adopt a more open frameworkrelated structure, but nonetheless contain a network of apex-linked BO6 octahedra as found for perovskites, is therefore of interest. Investigations of proton conductivity in pyrochlores have so far been focused on acceptor-doped derivatives of Ln2B2O7, where Ln = lanthanide and Y, and B = Zr and Ti. Original work claiming significant proton conduction in La2Zr2 − xYxO7 − δ was performed by Shimura et al. [14]. Further work, including recordings of infrared spectra [15,16], extensive electrochemical measurements [17] and quantum mechanical simulations [18] have been used to give additional information in relation to proton mobility in La2Zr2O7. Shimura et al. [14] also investigated Y2Ti1.8M0.2O7 − δ (M = In and Mg) and concluded that these materials did not exhibit proton conductiv-

⁎ Corresponding author. Tel.: + 46 317869036. E-mail address: [email protected] (C.S. Knee). 0167-2738/$ – see front matter © 2010 Elsevier B.V. All rights reserved. doi:10.1016/j.ssi.2010.09.008

ity. Haugsrud and Norby [19] studied a monoclinic A2B2O7 variant, La1.98Ca0.02Ti2O7 − δ, which was shown to exhibit mixed protonicelectronic conductor behaviour in wet hydrogen. In 2008 Fjeld et al. [20] reported on the conductivity of Er1.96Ca0.04Ti2O7 − δ containing Si impurities in wet oxidising conditions and found no contribution to the bulk conduction from protonic charge carriers. In contrast, the present authors found bulk proton conductivity in both Sm1.92Ca0.08Ti2O7 − δ and Sm2Ti1.92Y0.08O7 − δ [21]. Eurenius et al. [22] have recently expanded the available data on proton conductivity in pyrochlores with a report on proton conduction in tin based pyrochlores, Ln1.96Ca0.04Sn2O7 − δ (Ln = La, Sm and Yb). This study revealed that proton conduction is strongly dependent on the lanthanide size. In the present contribution the synthesis and characterisation of two acceptor-doped tin based pyrochlore systems, (Sm1.92Ca0.08)Sn2O7 − δ and Sm2(Sn1.92Y0.08)O7 − δ as well as the undoped system Sm2Sn2O7 is reported. 2. Experimental Sm1.92Ca0.08Sn2O7 − δ, Sm2Sn1.92Y0.08O7 − δ and Sm2Sn2O7 were prepared via conventional solid state reactions of high purity reactants (Sm2O3 (99.9%), CaCO3 (99.9%), SnO2 (99.9%) and Y2O3 (99.9%)). All oxides were dried (800 °C) to remove moisture or carbonates and increase accuracy upon weighing. The reactants were ground manually with ethanol in an agate mortar, sintered in air (1400 °C, 24 h; 2 × 1550 °C, 50 h) with regrinding and pelletizing (13 mm diameter die; P ~ 672.5 MPa = 10 tons) carried out between the heating steps. To run IR and thermogravimetric (TGA) analyses, small batches of the as-

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prepared powder samples were hydrated in a humidifier (300 °C, 120 h) under a flow of nitrogen saturated with H2O or D2O at 74.4 °C. Dried samples were prepared by annealing portions of the as-prepared samples at 1000 °C under vacuum (5× 10− 6 mbar, 12 h). To assess phase purity, powder X-ray diffraction (PXRD) was carried out at room temperature using a Siemens D5000 diffractometer (Cu-Kα = 1.5418 Å, 25–65° 2θ). High quality long scan data was then collected using a Bruker D8 Advance diffractometer with a primary monochromator (Cu-Kα1 = 1.54056 Å, 12–105° 2θ, 0.0197° step size). Unit cell constants were obtained from these patterns using the Rietveld profile fitting program GSAS [23]. The Fourier Transform Infrared (FTIR) measurements were performed in diffuse reflectance mode in the range (560–6000 cm− 1). The data was recorded with a Bruker IFS 66v/S vacuum FTIR interferometer with a KBr beam splitter and a Mercury Cadmium Tellurium detector. The system was flushed with dry CO2-free air. A reference spectrum was measured on ground KBr before collecting each sample (400 scans/run). The spectra were then derived by taking the logarithm of the ratio between the reference spectrum and the sample spectrum. TGA was performed on hydrated and vacuum dried powders during heating (25–1000 °C, 15 °C/min) with a NETZSCH STA 409 PC, under nominally dry gas (N2, 10 ml/min). Electrochemical Impedance Spectroscopy (EIS) measurements in the frequency range 4.5 MHz–1 Hz were carried out with a Solartron 1260 frequency response analyser in the stand alone mode connected to a conductivity cell (ProboStat™; Norwegian Electro Ceramics AS, (NorECs) [24]. The sine wave amplitude was 1 V rms. To enable good ohmic contact, around 0.8 cm2 (A- and B-site doped) and 0.4 cm2 (undoped) of the pellets' surfaces were painted with conducting Ptpaste. The area of Pt coverage was obtained accurately using an Optic Zeiss 120 HD Microscope. The data sets were collected upon cooling (1000–150 °C) in 50 °C intervals (30 min. equilibration time) under a flow of gas on pellets with relative densities of 86%, 83% and 79% for A-site, B-site and undoped materials respectively. The data were collected on the same pellet first under dry (Ar/O2 flowed over P2O5) and then wet gas (Ar/O2 bubbled through H2O or D2O at RT). Two silica tubes were used to protect the cell from humidity during the dry runs while the wet runs were carried out with an alumina outer tube (AL23). Given the reported instability of related tin based pyrochlores under reducing atmospheres [12], conductivity measurements under H2 were not attempted. The impedance spectra were fitted to equivalent circuits containing one to three time constants. For all time constants a constant phase element (CPE) was used in parallel with a resistor. The effective capacitance was calculated using Eq. (1) [25]. C=Q

1 = α ð1 = α−1Þ

R

ð1Þ

intensity of the pyrochlore supercell reflections is very weak, and we attribute this to the lack of scattering contrast (i.e. similar atomic numbers) between the Sm and Sn ions. The Rietveld scans also revealed the presence of weak reflections not accounted for by the pyrochlore structure for the Sm2Sn1.92Y0.08O7 − δ sample. The position of these peaks indicate that they probably arise from a cubic (Y,Sm)2O3 phase suggesting that doping with Y2O3 may be incomplete, Fig. 1. The IR spectra for the vacuum dried, as-prepared, deuterated, and protonated samples are shown in Fig. 2. For both A and B-site doped materials several differences are clear between the as-prepared and vacuum dried samples, and the samples after treatment in H2O and D2O rich atmospheres. The protonated A-site doped sample (Fig. 2a) has four intense peaks located at 3454, 3423, 3388 and 3314 cm− 1 that are assigned to distinct O–H stretch vibrations indicating the presence of dissolved protons in the structure. Very weak signs of these peaks are present in the as-prepared sample but they are completely absent from the dried Sm1.92Ca0.08Sn2O7 − δ sample and also un-doped Sm2Sn2O7. The hydrated B-site doped system (Fig. 2b) shows a single peak at 3600 cm− 1 that is absent in both the asprepared and dried samples. Confirmation that these features are linked to dissolved protons comes from the data collected on the deuterated samples. Here, isotopic shifts of the most intense peaks at 3454 cm− 1 and 3600 cm− 1 for the A and B-site substituted samples respectively are observed, with ratios between the υOH and υOD ~ 1.35. This compares with an expected ratio of ~1.37 reflecting the change in mass between O–H and O–D groups. An O–D peak is also seen at 2459 cm− 1 (not shown) which is related by the same ratio to the 3314 cm− 1 band of the A-site protonated sample. It is noteworthy that both the deuterated Sm1.92Ca0.08Sn2O7 − δ and Sm2Sn1.92Y0.08O7 − δ samples display the same O–H vibrations as the protonated phases, indicating that protonation, as well as deuteration, occurred. The IR spectra of the hydrated Sm2Sn2O7 sample did not show any sign of peaks attributable to dissolved protons. All samples show peaks located at approximately 3523 cm− 1 and 2540 cm− 1 (labelled with * in the figures). As these bands show no dependence on the hydration history, and are also present in the as-prepared parent undoped material (Fig. 2), we assign them as overtones of the structural phonons. The TGA data shown in Fig. 3 reveal small but distinct mass losses beginning at ~400 and 275 °C for the hydrated A-site and B-site substituted samples respectively, which signifies the loss of protons as H2O (g). The traces reveal a 0.08% mass loss for the A-site substituted sample and a 0.02% decrease for Sm2Sn1.92Y0.08O7 − δ. The theoretical mass change expected for complete filling of all oxygen vacancies by [OH.] defects and protonation of a nearby lattice oxygen is, in both, cases 0.11%. This calculation assumes complete acceptor substitution at the desired crystallographic site and charge compensation through oxygen vacancy formation only, i.e. δ = 0.04. The results therefore

where Q and α are the parameters related to the constant phase element (1/ZCPE = (Q(jω)α) and R is the resistance. Characterisation of the microstructure was carried out using a Leo Ultra 55 SEG SEM, operated with an acceleration potential of 2.5 kV and secondary electron (SE) detector. 3. Results The products obtained were off-white powders and the PXRD patterns collected after each heating step revealed that a second annealing at 1550 °C was necessary to ensure a complete reaction. Rietveld analysis of the long scan data sets allowed cell parameters of 10.5140(2), 10.5098(2) and 10.5158(2) Å for Sm1.92Ca0.08Sn2O7 − δ, Sm2Sn1.92Y0.08O7 − δ and Sm2Sn2O7 to be determined respectively. The presence of weak reflections at 2θ ≈ 37 and 45° which may be fitted using Fd-3m symmetry confirm that the long range structure of the materials is pyrochlore type, rather than oxygen deficient fluorite, reflecting ordering of the Sm and Sn ions on the A and B site. The relative

Fig. 1. PXRD patterns for Sm2Sn1.92Y0.08O7 − δ, Sm1.92Ca0.08Sn2O7 − δ and Sm2Sn2O7 from top to bottom respectively. * Marks the position of reflections that confirm the pyrochlore structure for Sm2Sn2O7 and arrows mark the presence of an apparently cubic additional phase in the Sm2Sn1.92Y0.08O7 − δ sample.

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2.4 ± 0.1 10− 11 Fcm− 2 for A-site doped, B-site doped and undoped samples, respectively. These capacitance values give dielectric constants of 35± 4, 28 ± 3 and 45 ± 3. For the A-site doped sample a time constant for the grain boundary conduction was observed in the intermediate frequency region with a capacitance of 2.8 ± 0.4 10− 9 Fcm− 2 after correcting for the frequency dispersion using Eq. (1). At the lowest frequencies a time constant for the electrode processes is also observed with a high capacitance value, see Fig. 4a. For the B-site doped sample it is not possible to separate the bulk and the grain boundary processes and the impedance was fitted to

Fig. 2. Infrared absorbance spectra of Sm1.92Ca0.08Sn2O7 − δ (a) and Sm2Sn1.92Y0.08O7 − δ (b). The spectra from top to bottom are of hydrated, deuterated, as-prepared and vacuum dried samples. The scan of as-prepared Sm2Sn2O7 is also included at the bottom of both figures. The spectra have been vertically shifted to aid clarity. * indicate features not linked to O–H/D vibrations.

indicate that ~ 72% and ~18% of vacancies are filled for the A- and Bsite substituted samples respectively. The undoped sample is not included since it showed no detectable mass loss for the vacuum dried or the hydrated sample. The impedance spectra recorded for the three samples differ. For the A-site doped sample in general three time constants were observed, the undoped sample showed overall two time constants whilst for the B-site only one is seen as shown in Fig. 4. At high frequencies the time constant can be attributed to the bulk conduction. The capacitances in the high frequency region were calculated using Eq. (1) and data taken at 150 °C to give values of 2.4 ± 0.3· 10− 11 Fcm− 2, 1.0 ± 0.1 · 10− 11 Fcm− 2 and

Fig. 3. TGA plots for acceptor-doped Sm2Sn2O7. The lines shows scans obtained for vacuum dried Sm1.92Ca0.08Sn2O7 − δ (provided as a reference), hydrated Sm2Sn1.92Y0.08O7 − δ and hydrated Sm1.92Ca0.08Sn2O7 − δ from top to bottom.

Fig. 4. Complex plane plots for Sm1.92Ca0.08Sn2O7 − δ (a), Sm2Sn1.92Y0.08O7 − δ (b) and Sm2Sn2O7 (c) under dry (spectra 1) and wet (H2O (spectra 2), D2O (spectra 3)) oxygen. In all cases the temperature was 400 °C. A fit to the dry run impedance data is shown in Fig. 4a and the equivalent circuit used is also shown in the figure.

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one time constant, Fig. 4b. The capacitance value indicates that bulk conduction dominates. For the undoped sample only the high frequency time constant was evaluated and related to the bulk conduction. The time constant at lower frequencies is not well defined and the capacitance value indicates that the impedance is due to electrode processes, Fig. 4c. A comparison of the impedance spectra collected at 400 °C in dry O2, and O2 saturated with H2O and D2O for all samples is shown in Fig. 4. From the figure it is immediately clear that there is a significant decrease in the resistance for all materials in the presences of protons at this temperature. From further examination of the impedance spectra it is readily seen that all samples exhibit significantly higher conductivity at temperatures up to ~550 °C in wet atmospheres. A clear isotope effect is also apparent from the experiments performed in O2/D2O atmosphere. The general trends seen here are the same with Ar as the carrying gas. Fig. 5 shows Arrhenius plots of the bulk conductivity for the asprepared Sm1.92Ca0.08Sn2O7 − δ, Sm2Sn1.92Y0.08O7 − δ and Sm2Sn2O7 samples under flows of dry gas (O2 and Ar). At 1000 °C the conductivity in O2 atmosphere is relatively high (7 · 10− 3 S cm− 1) for the A-site doped compound. For the B-site substituted and undoped compounds the maximum in conductivity is a factor of ten lower than for the A-site substituted sample. The conductivity is significantly higher in oxygen compared to argon for the B-site and undoped samples while for the A-site substituted sample the conductivity is only slightly higher in oxygen atmosphere. For the undoped and B-site doped samples the conductivity in Ar at low temperatures tends to cross over towards the curve obtained in O2, indicating two different conduction processes. In Figs. 6–8 the Arrhenius plots for the A-site, B-site and undoped materials in dry gas (O2 and Ar) and wet gas (O2/Ar + H2O or O2/Ar + D2O) are shown. At temperatures below ~ 550 °C the bulk conductivity of the samples under wet gas was higher compared to the equivalent dry gas runs. The enhancement is in general larger in Ar compared to O2 and differs between the samples. At 400 °C the enhancement is about 20 (Ar) and 10 (O2), 20 (Ar) and 3 (O2) and 10 (Ar) and 3 (O2) for A-site, B-site and undoped samples, respectively. The conductivity of the deuterated runs lie below the runs performed in H2O at low temperatures, which agrees with the expectations for a proton conductor. The isotope effect in oxygen is 2.5 ± 0.2 for A-site, 1.7 ± 0.1 for B-site and 1.4 ± 0.1 for undoped sample, measured in the temperature range 200–350 °C. For the A-site (Ar and O2) and B-site (Ar) doped materials, the conductivity change is larger than expected from the isotope effect, while for the B-site (O2) and undoped (Ar) samples the change in conductivity is consistent with the isotope effect. The conductivity values at 1000 °C and 300 °C are given in Table 1

Fig. 5. Bulk conductivity of Sm1.92Ca0.08Sn2O7 − δ (▲), Sm2Sn1.92Y0.08O7 − δ (■) and Sm2Sn2O7 ( ) under dry oxygen (open symbols) and dry Argon (filled symbols).

•

Fig. 6. Arrhenius plots of the bulk conductivity of Sm1.92Ca0.08Sn2O7 − δ under Ar (a) and O2 (b). Dry gas ( ), H2O (▲) and D2O (Δ).

•

Fig. 7. Arrhenius plots of the bulk conductivity of Sm2Sn1.92Y0.08O7 − δ under Ar (a) and O2 (b). Dry gas ( ), H2O (▲) and D2O (Δ).

•

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Table 1 Activation energies (Ea) and bulk conductivities (σ) for Sm2Sn2O7, Sm1.92Ca0.08Sn2O7 − δ and Sm2Sn1.92Y0.08O7 − δ. Ea/eVb

Dry Ar Ar/H2O Ar/D2O Dry O2 O2/H2O O2/D2O

0.98 0.76 0.80 1.03 0.81 0.83

0.74 0.74 0.80 0.97 0.91 0.93

Sm2Sn1.92Y0.08O7 − δ

Ea/eVa

Ea/eVb

σ/nScm− 1c

σ/mScm− 1d

Dry Ar Ar/H2O Ar/D2O Dry O2 O2/H2O O2/D2O

0.73 0.85 1.05 1.17 0.80 0.94

1.41 1.20 1.34 1.17 1.10 1.09

1.53 34.6 9.40 1.30 18.2 10.5

0.19 0.17 0.19 0.59 0.58 0.66

Sm2Sn2O7

Ea/eVa

Ea/eVb

σ/nScm− 1c

σ/mScm− 1d

Dry Ar Ar/H2O Ar/D2O Dry O2 O2/H2O O2/D2O

0.88 0.94 0.88 1.05 0.94 0.96

1.12 1.09 1.06 1.01 0.94 0.96

3.11 28.6 18.1 4.29 19.6 14.0

0.21 0.22 0.22 0.78 0.80 0.81

b c d

•

together with the activation energies estimated in the low (150– 400 °C) and high (750–1000 °C) temperature regions. At intermediate temperatures the wet gas conductivity curves display the characteristic hydration/dehydration behaviour of proton conductors, and above ~600 °C the conductivity of the wet and dry runs were almost identical in each gas. Under wet conditions the conductivity was again significantly higher for the A-site doped system than the undoped and B-site doped samples. To obtain further insight into the hydration process, the hydration enthalpy, along with the activation energies and pre-exponential factors, were estimated by modelling the conductivity data for the A- and B-site doped samples utilising well known expressions for the equilibrium constant of the hydration reaction and the conductivity of species migrating via activated diffusion [26,27]. The values obtained are summarised in Table 2. SEM was carried out on the conductivity pellets, see Fig. 9. The variation in grain size is large with a typical size of ~ 5 μm for Sm1.92Ca0.08Sn2O7 − δ. For the Sm2Sn1.92Y0.08O7 − δ and Sm2Sn2O7 samples the grain sizes are larger, typically N10 μm. 4. Discussion 4.1. Characterisation by PXRD, FTIR and TGA In an effort to increase the conductivity of Sm2Sn2O7, acceptor doping on either the A or B-site was carried out to achieve the same nominal level oxygen vacancies, via partial substitution of Sm3+ for Ca2+ at the A-site and Y3+ for Sn4+ on the B-site. The ionic radii of the dopant ions are larger than the Sm and Sn ions that they are expected to replace. The refined cell parameters, however, do not show an expansion for the acceptor-doped samples in comparison to Sm2Sn2O7. For the B-site substituted sample the minor impurity present (Fig. 1) suggests incomplete doping, and potentially some Y

σ/mScm− 1d

Ea/eVa

a

Fig. 8. Arrhenius plots of the bulk conductivity of Sm2Sn2O7 under Ar (a) and O2 (b). Dry gas ( ), H2O (▲) and D2O (Δ).

σ/μScm− 1c

Sm1.92Ca0.08Sn2O7 − δ

0.06 2.63 0.72 0.06 1.06 0.47

3.9 4.0 4.6 7.5 7.5 7.6

For temperatures b400 °C. For temperatures N700 °C. At 300 °C. At 1000 °C.

has replaced Sm at the A-site as this would explain the observed decrease in cell parameter. For the A-site substituted sample other factors linked to the strength of the Sm–O(1) 48f interaction [28], covalency of the Sn–O(1) bond [29], and possible structural relaxation around the doped oxygen vacancies may influence the cell parameter. After exposure to wet gas, acceptor-doped oxides are often found to hydrate via the filling of oxygen vacancies by OH-groups. To examine the hydration levels in the materials qualitative IR spectroscopy measurements were performed and complemented by quantitative TGA. IR also provides information on the local structure of the proton sites in the host material. The spectra shown in Fig. 2 gave clear indications of an increase in proton concentration for the hydrated Sm1.92Ca0.08Sn2O7 − δ and Sm2Sn1.92Y0.08O7 − δ samples. For Sm1.92Ca0.08Sn2O7 − δ four relatively sharp peaks were seen for the sample in the region 3300–3460 cm− 1 suggesting that there are four specific proton positions, with a small spread of energies. For Sm2Sn1.92Y0.08O7 − δ only a single sharp υOH peak is observed at 3600 cm− 1 suggesting one, specific, proton site. Closer inspection of Fig. 2b also reveals a weak and broad feature in the region 3100– 3400 cm− 1 is present for the protonated and deuterated samples. This

Table 2 Hydration enthalpies (ΔH), pre-exponential factors, and activation energies for proton and deuteron conductivity extracted from fitting the conductivity data as described in the text. During the fitting process the entropy term (ΔS) was fixed =−120 J mol− 1 K− 1 [26,27]. Sm1.92Ca0.08Sn2O7 − δ

ΔH/kJ mol− 1

U0,H+/cm2KV− 1 s− 1

Ea,

Ar/H2O Ar/D2O O2/H2O O2/D2O

− 100 ± 3 − 90 ± 1 − 102 ± 2 − 100 ± 3

37 ± 10 56 ± 6 34 ± 8 35 ± 12

0.84 0.96 0.87 0.92

Sm2Sn1.92Y0.08O7 − δ

ΔH/kJ mol− 1

U0,H+/cm2KV− 1 s− 1

Ea, H+/eV

Ar/H2O Ar/D2O O2/H2O O2/D2O

− 85 ± 1 − 80 ± 2 − 112 ± 2 − 97 ± 6

17 ± 3 74 ± 24 83 ± 3 84 ± 8

0.96 1.19 0.86 0.87

H+/eV

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O(2), 8b, oxygen and the O(1), 48f site, that bonds solely to the octahedrally coordinated B-site ion, respectively. The greater number of bands seen in the IR results suggest the presence of protonic defects at both the O(2) and O(1) sites for the A-site doped sample, and only at the O(1) position for Sm2Sn1.92Y0.08O7 − δ. This behaviour is similar to that observed for hydrated samples of Sm1.92Ca0.08Ti2O7 − δ and Sm2Ti1.92Y0.08O7 − δ [21]. The TGA results support the findings from the IR analysis and indicate that the level of protonic [OH.] defects is much greater in the A-site doped sample compared to the B-site substituted material. At present we are not certain why the hydration levels are so different for the two different doping strategies, but note that Ca2+ will create oxygen vacancies regardless of the cationic site it occupies, whilst the introduction of Y3+ should be effective only at the B-site. The presence of a minor impurity phase in Sm2Sn1.92Y0.08O7 − δ also suggests that incorporation of the acceptor dopant Y is not complete, and this is consistent with a lower proton concentration. It should be remembered that the conductivity runs were performed under much lower H2O vapour pressures compared to the hydration experiments, so a direct correlation between the proton concentrations determined from the TGA data and the level of proton charge carriers present is not possible. We may, though, conclude that Sm1.92Ca0.08Sn2O7 − δ has a greater tendency to absorb protons. For the hydrated undoped material no bands related to protons were observed in the IR spectra, which is in agreement with the TGA data where no mass loss was seen. However, these observations contrast with the impedance results where a clear enhancement in conductivity was observed at lower temperatures in wet gas as discussed further in Section 4.2.3 below. 4.2. Electrochemical impedance spectroscopy 4.2.1. Dielectric properties The dielectric constants obtained from the impedance measurements vary with the doping and follow the order undoped (45), A-site doped (35) and B-site doped (28). These values are much larger than those reported for other pyrochlore stannates [31]. The dielectric constant gives information about the polarisability of the material and consists of two terms, the induced and orientation polarisabilities. At high frequencies only the induced polarisability is measured, i.e. the electron cloud movement. At lower frequencies both processes takes place. The relationship between oxide-ion conductivity and dielectric relaxation has been studied for the related Ln2Zr2O7 system [32] and a clear frequency dependence of the dielectric properties was obtained. The low values of the dielectric constant reported previously [31] were measured at high frequencies (1 GHz) and is probably related to the induced polarisation while the values obtained in the present paper were calculated from the capacitance values measured at frequencies lower than 1 MHz and include also the orientation polarisation. Fig. 9. SEM images for (a) Sm1.92Ca0.08Sn2O7 − δ, (b) Sm2Sn1.92Y0.08O7 − δ and (c) Sm2Sn2O7.

resembles the IR signal commonly observed for hydrated perovskites [30], and it may be linked to protons that are strongly hydrogen bonded in the pyrochlore structure. From the IR data we can therefore conclude that there are more distinct proton sites in the hydrated A-site doped sample, suggesting, but not confirming, a greater concentration of dissolved protons. Furthermore, the lower frequency of the O–H stretch seen for the Ca containing sample suggests longer O–H separations which may reflect differences in the local environment around the dopant ions and/or possibly stronger hydrogen bonding in the material. Omata and Otsuka-Yao-Matsuo [16] have previously assigned three bands for protonated La1.96Ca0.04Zr2O7 − δ with two intense bands at 3517 and 3401 cm− 1. These bands were attributed to protons sites bound to the

4.2.2. Dry gas behaviour The significantly higher conductivity for the A-site substituted sample (dry gas behaviour, Fig. 5) is consistent with Ca doping producing an increased level of oxygen vacancies as observed for related pyrochlores by, for example, Tuller and co-workers [10,12]. In contrast, the conductivity of the B-site substituted sample in fact lies slightly lower than the undoped sample. This suggests that either the increased level of oxygen vacancies introduced via substitution at the Bsite is offset by an accompanying reduction in the average mobility of the vacancies, and/or that the ionic (oxide ion) contribution to the conductivity is in fact low. Support for the latter scenario is seen as the conductivity is, in all cases, higher in oxygen compared to argon, indicating that the conduction is not purely ionic, and more specifically that an enhanced level of p-type electron hole conduction occurs under oxidising conditions as found previously for Gd2 − xCaxSn2O7 systems [10,12]. Note that Sm2Sn2O7 shows the largest contribution of hole

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conduction followed by the B-site doped sample. For the A-site doped sample only a small difference in conductivity is observed in Ar and O2, indicating that the oxide ion conduction is significant in this case. The activation energies in oxygen atmosphere are roughly the same for the A-site, B-site and undoped samples in the entire temperature range studied, with values close to 1 eV as seen in Table 1. Thus the conduction mechanism seems to be the same with a mixture of ionic and electronic conduction. The activation energies in Ar differ more between the samples and in the different temperature intervals. For the A-site compound a lower activation energy, 0.74 eV, is found at high temperatures reflecting a possible change from mixed to purely ionic conduction. The same activation energy is found for the B-site doped material in the low temperature range. Since for this material a clear difference in conductivity is observed in Ar and O2, the conduction process should be attributed to the transport of holes. In general the activation energies for the B-site material are higher than the corresponding values for the A-site material. For the undoped material the activation energies follow the same trend. 4.2.3. Wet gas behaviour As can be seen in Figs. 6 and 7, both acceptor-doped samples show higher bulk conductivity at T ≤ 500 °C under wet gas which is typical for proton conductivity. At higher temperatures, the data from both wet and dry runs coincide, being characteristic for loss of protons and a shift to another central conduction process. Evidence pointing towards proton diffusion being the main conduction process comes from the IR and TGA results, which respectively show significant amounts of dissolved protons in the doped materials. The conductivity data collected in the presence of D2O again support this conclusion, due to the expected decrease in conductivity for the heavy water runs. The present findings follow the trend of higher levels of proton conduction found for A-site substituted pyrochlores in comparison to B-site doped systems [17,21]. As for these previous studies, the experimental evidence from IR and TGA leads us to conclude that this effect is mainly linked to a greater proton concentration in the A-site substituted material. For the undoped material the same general trends are observed under wet gas conditions, Fig. 8. In particular, the presence of an apparent isotope effect for the D2O runs supports proton conduction as the dominant process for the material even in the absence of acceptor-doped oxygen vacancies. In spite of the high disordering energies associated with stannate pyrochlores [29] some intrinsic anti-site defects must exist in the Sm2Sn2O7 sample. There is thus the likelihood of oxygen vacancies at either the 48f (x, 1/8, 1/8) and/or 8b (1/8, 1/8, 1/8) site(s) associated with Sm3+ residing on the Sn4+ site, which may allow the material to absorb and conduct protons. Alternatively, partial occupation of the vacant 8a (3/8, 3/8, 3/8) position in the pyrochlore structure by [OH] defects may occur. In either case the level of protonic charge carriers would be anticipated to be extremely low and this possibly explains the fact that IR and TGA experiments fail to show the presence of protons in the hydrated sample, see Section 4.1. The low temperatures regions where the proton conduction can be seen (Table 1), show activation energies of 0.76 eV for A-site, 0.82 eV for B-site and 0.94 eV for undoped materials in Ar, respectively. Similar values are obtained with O2 as the carrying gas. These values are higher than typical values found for proton conducting perovskites (~ 0.4–0.7 eV) [32,33] and are also somewhat higher than those observed for acceptor-doped La2Zr2O7 [17], as well as Sm1.92Ca0.08Ti2O7 − δ and Sm2Ti1.92Y0.08O7 − δ pyrochlores [21]. The highest activation energies found for proton conduction in perovskites have been attributed to strong “trapping” of the protons at the dopant cation site [33,34]. It was found that the activation energy increases as the size of the doping cation decreases. A similar mechanism can be envisioned for pyrochlores, where strong interaction between protons and defect sites or dopant ions will lead to higher activation

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energies [18]. The activation energies obtained in wet D2O atmosphere are somewhat higher than the corresponding energies in H2O gas, which is reasonable. The dry and wet conductivity runs merge at ~ 700 °C for all samples reflecting dehydration at higher temperatures. The dehydration during TGA is complete at ~ 550 °C for the A-site doped compound and at ~400 °C for the B-site doped sample, Fig. 3. The lower dehydration temperatures found in the TGA analyses can be explained by the fact that the TGA runs were performed under a flow of nominally dry gas in comparison to the conductivity runs which were performed in humid gas flows. The hydration enthalpies of both A and B-site substituted samples lie within the range −100±20 kJmo1− 1 (Table 2) which compares with the extremes for proton incorporation into perovskites of −163 kJmol− l for BaCe0.9Y0.1O3 − δ and −22 kJmol− l for SrTi0.98Sc0.02O3 − δ [35]. The calculated hydration enthalpies for the A- and B-site substituted samples can be compared to the significantly lower dehydration temperature observed for the pre-hydrated B-site sample from the TGA data. The experimental evidence thus suggests that the hydration process is less exothermic for Sm2Sn1.92Y0.08O7 − δ when compared to Sm1.92Ca0.08Sn2O7 − δ. This shows reasonable agreement with the somewhat lower ΔH values extracted for the B-site doped sample under Ar (−85 kJmol− 1, Table 2). The TGA de-hydrations were also performed under nominally non-oxidising (N2) conditions. It should, however, be stressed that the reliability of extracting the three parameters listed in Table 2 was affected by the relative lack of data points available. The curve fitting also required several assumptions, with for example the entropy term (ΔS) being fixed= −120 Jmol− 1 K− 1, i.e. equivalent to the change in entropy for the loss of one mole of gaseous water [26,27]. Such factors may contribute to the activation energies determined via the modelling process being, in general, ~0.1 eV higher than those determined from fitting of the Arrhenius plots (see Tables 1 and 2). Finally, we mention that the conductivity data reveals a significant decrease in conductivity for O2/H2O runs compared to Ar/H2O runs for all samples in the proton conducting temperature interval. Our understanding of this is that at higher oxygen partial pressure the formation of electron holes competes more strongly with the formation of protonic [OH.] defects, thus suppressing the level of proton conductivity observed. 4.2.4. High temperature behaviour The disordering tendencies in pyrochlore oxides were recently investigated theoretically [29]. Besides the general view that the size differences between A3+ and B4+ cations is decisive for ordering in the pyrochlore structure it was shown that the covalency of the B–O (1)48f bond plays an important role for the ordering [28]. The covalency of the Sn–O(1)48f bond is large and accompanied by a large disordering energy. The disordering energy is also dependent on the nature of the A-site cation and increases with increasing radius for pyrochlore stannates. One might expect that a low disordering energy should lead to increased oxide ion conductivity, as greater cationic disorder in pyrochlore systems generally favours increased levels of vacancies and as a result higher conductivity [10]. However, comparing conductivity results for stannates from different sources [10,11], shows that for pyrochlores with larger A-site cations, which are expected to have a more perfectly ordered cation sub-lattice, the conductivity in the high temperature region is higher than for smaller A-site ions, Fig. 10. The figure shows a comparison of the conductivity in O2 for different undoped stannate pyrochlores, including our data for Sm2Sn2O7. The conductivity increases and the activation energy decreases with an increase in the A-site cation radius. The same trends are observed in Ar or low oxygen partial pressures with, in general, somewhat lower activation energies. This may be taken to support the view that cation disordering in Ln2Sn2O7 systems is very unfavourable for all lanthanide ions as predicted computationally [29] and found experimentally for Gd2Sn2O7 [12]. The enhanced conductivity and

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The brick layer model was used in the present work to extract the specific grain boundary conduction using Eq. (2). spec

σgb

=

  L δ 1 A D Rgb

ð2Þ

where L is the thickness of the sample, A the cross sectional area, δ is the grain boundary thickness and D the grain size. While D can be approximated from SEM analysis, δ is not easily accessible. Assuming that the dielectric constant for the grain interior and the grain boundaries are the same, the ratio between δ and D can be estimated from the impedance measurements, Eq. (3). δ C = bulk D Cgb Fig. 10. Arrhenius plots for Y2Sn2O7, La2Sn2O7 [both 11], Gd2Sn2O7 [10] and Sm2Sn2O7 [this work] under high temperature conditions.

lower activation energy then predominantly reflects more facile oxide ion migration through a relaxed (larger) crystal lattice rather than changes in the concentration of mobile species linked to disorder. Our results show that the activation energies in the high temperature region are dependent on the carrying gas. For the A-site doped material in Ar the activation energy is similar in the high temperature interval 700–1000 °C with Ea = 0.74–0.8 eV. In dry O2, the activation energy is the same in the low and high temperature region while in wet atmosphere the activation energies are somewhat higher in the high temperature region, see Table 1. The unexpectedly low value for the activation energy at high temperatures in dry Ar indicates a possible change to a purely ionic conduction process as discussed earlier. The conductivities at 300 °C and 1000 °C are also listed in Table 1. At 1000 °C the conductivity under dry and wet conditions are the same but differ with the carrier gas and is in general larger in O2. The ratio is less than 2 for A-site and about 3 and 4 for B-site and undoped material, respectively. As discussed in Section 4.2.2 this shows that the conduction is dominated by holes for the B-site and undoped samples under oxygen, while oxide ion conduction also contributes to the total conduction for the A-site doped material.

4.2.5. Grain boundary conduction Fig. 4a displays the complex plane plots for Sm1.92Ca0.08Sn2O7 − δ showing that both bulk and grain boundary resistance play an important role for the total conductivity. This system follows the trend generally observed for BaZrO3 perovskites where high grain boundary impedance greatly limits the total conductivity [8,34,36] but contrasts with the low grain boundary impedance observed for Sm1.92Ca0.08Ti2O7 − δ [21]. In order to compare the grain boundary conduction for different systems the specific grain boundary conduction must be calculated. The most commonly used model to account for grain boundary conduction is the brick layer model [37,38]. This model was first introduced for oxide ion conductors and later also used for proton conductors [39,40]. In the brick layer model the grains are approximated as cubes of the same size and the grain boundary conduction is divided in a parallel and a perpendicular part compared to the electric field. For materials where the grain boundaries are easy conduction parts the parallel component is dominating while for low conducting grain boundaries the perpendicular part will, together with the in grain conductivity, make up the total conductivity for the material. Although the brick layer model is very simple, it gives a good qualitative description of the grain boundary conduction and for example the use of more realistic models for the grains only marginally influences the calculated conduction values [41]. The effect of microstructure on the conductivity of polycrystalline oxides have been reviewed recently [42].

ð3Þ

where Cbulk and Cb are the bulk and grain boundary capacitances, respectively. For the A-site doped sample, the ratio was determined to 0.009 ± 0.002, which means that the specific grain boundary conduction is about two orders of magnitude lower than the total conductivity. It was found that the specific grain boundary conduction follows a similar temperature dependence as the bulk conductivity, indicating that proton conduction also occurs in the grain boundaries but with a much lower conductivity. It has previously been shown that for systems where the time constant for the specific grain boundary conduction is clearly separated from the time constant for the bulk conduction, then the specific grain boundary conduction will limit the total conductivity [41]. The SEM analysis shows that the materials are somewhat porous with large variation in the grain size For the A-site doped sample a typical grain size is 5 μm and this value has been used to estimate the thickness of the grain boundaries from the ratio between the bulk and grain boundary capacitances. A value of 45 ± 10 nm is obtained indicating that the grain boundaries are heavily distorted as this corresponds to a region ~45 unit cells thick provided that the Brick Layer Model is applicable. The large value for the grain boundary thickness may also indicate that electronic effects, such as the creation of space charge zones in the vicinity of the grain boundary, are important [43,44]. For B-site and undoped material, the time constant for the grain boundary conduction cannot be separately obtained. The grain size for these materials is larger than for the A-site doped material, typically larger than 10 μm, which might explain the lack of a separate time constant for grain boundary conduction. For a more detailed description of the grain boundary conduction better control of the grain size and the density of the pellets is necessary. 5. Conclusions In the present paper it was demonstrated, with evidence from IR spectroscopy, TGA and isotope measurements of the conductivity, that Sm1.92Ca0.08Sn2O7 − δ can conduct protons at intermediate temperatures, b550 °C. The bulk proton conductivity at 300 °C amounts to 2.6 and 1 μScm− 1 in H2O-saturated Ar and O2, respectively. In dry gas mixed oxide ion and hole conduction takes place with a maximum conductivity at 1000 °C of 4.2 (Ar) and 7.5 (O2) mScm− 1. The conductivity of Sm2Sn1.92Y0.08O7 − δ and undoped Sm2Sn2O7 generally lies about two orders of magnitude lower than for the A-site doped sample. In dry oxygen these materials show mainly hole conduction and both exhibit significant proton conduction at low temperatures in the presence of water vapour. Finally, we note that the relatively low level of bulk proton conduction, high grain boundary impedance, and the tendency for Snbased systems to decompose in reducing atmospheres are all factors likely to exclude Sm2Sn2O7 derivatives from any potential SOFC related applications. Nonetheless, the enhancement of conductivity seen for

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Sm2Sn2O7 under wet conditions is an intriguing result worthy of further investigation. Acknowledgements C.S. Knee acknowledges the financial support of the Swedish research council (Vetenskapsrådet) for this research. K.E.J. Eurenius acknowledges support for travel costs to attend conferences from Stiftelsen Wilhelm and Martina Lundgrens Vetenskapsfond. The authors are grateful to Istaq Ahmed, Rikard Elmén, Ezio Zanghellini, all Chalmers University of Technology, and Patrick Steegstra, University of Gothenburg, for helping to collect EIS, IR and SEM data respectively. References [1] [2] [3] [4] [5] [6] [7] [8] [9] [10] [11] [12] [13] [14] [15] [16] [17]

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