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Proton and oxide ion conductivity in grain boundaries and grain interior of Ca-doped Er2Ti2O7 with Si-impurities
Solid State Ionics 179 (2008) 1849–1853
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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 e v i e r. c o m / l o c a t e / s s i
Proton and oxide ion conductivity in grain boundaries and grain interior of Ca-doped Er2Ti2O7 with Si-impurities Harald Fjeld a,⁎, Reidar Haugsrud a, Anette Eleonora Gunnæs b, Truls Norby a a b
Department of Chemistry, University of Oslo, FERMiO, Gaustadalléen 21, NO-0349 Oslo, Norway Department of Physics, University of Oslo, FERMiO, Gaustadalléen 21, NO-0349 Oslo, Norway
A R T I C L E
I N F O
Article history: Received 8 December 2007 Received in revised form 9 May 2008 Accepted 3 June 2008 Keywords: Er2Ti2O7 Grain boundaries Ionic conductivity Impedance spectroscopy
A B S T R A C T Dense samples of nominal composition Er1.96Ca0.04Ti2O6.98 were prepared by conventional ceramic synthesis to evaluate the ionic conductivity. The electrical properties were studied by impedance spectroscopy under oxidizing conditions as a function of pO2, pH2O and pD2O in the temperature range 360–600°C. The microstructure was investigated by analytical transmission electron microscopy before and after the impedance study. In both cases the grain boundaries contained an amorphous phase, enriched in Si and Ca. Three different contributions to the sample impedance, with different time constants, were identified, one assigned to the grain interior and two to grain boundaries. For the grain interior, oxygen vacancies were found to be the major charge carrier. In contrast, protons were found to play an important role as charge carriers in the grain boundary related processes. The ionic transport over grain boundaries in Er1.96Ca0.04Ti2O6.98 is discussed in terms of effects of additional phases in the grain boundaries and space charge layers. © 2008 Elsevier B.V. All rights reserved.
1. Introduction Acceptor-doped oxides are known to exhibit transport of ionic charge carriers such as oxide ions (in e.g. ZrO2 and CeO2 [1]) and protons (in e.g. SrCeO3 [2] and LaNbO4 [3]) at elevated temperatures, and find possible applications in electrochemical devices for future energy technology. In a polycrystalline material, both the grain interior and the grain boundaries contribute to the overall electrical properties. This implies that even though the grain interior of the material is an excellent ionic conductor, the practical use of the material may be limited by a high grain boundary resistance. In-depth knowledge of the transport of charge carriers in both the grain interior and grain boundaries is therefore of great importance. The detrimental role of the grain boundaries is known for e.g. acceptor-doped ZrO2, where the overall ionic conductivity is decreased due to the grain boundary resistance [4]. The grain boundary resistance can be divided into (at least) two contributions: 1) intergranular siliceous phases partially covering the grain boundary area [5], leading to constriction of current lines [6], and 2) depletion of oxygen vacancies in the vicinity of the grain boundaries; the space charge layer effect [7,8]. The grain boundary core-space charge layer model [9] has also been successfully applied to explain why electrons, in addition to oxygen vacancies, contribute to the grain boundary conductivity in acceptordoped CeO2 [10] and why electron holes dominate the grain boundary conductivity in Fe-doped SrTiO3 [11,12] under conditions where the grain interior conductivity is essentially ionic. ⁎ Corresponding author. Tel.: +47 22840657; fax: +47 22840651. E-mail address: [email protected] (H. Fjeld). 0167-2738/$ – see front matter © 2008 Elsevier B.V. All rights reserved. doi:10.1016/j.ssi.2008.06.003
An interesting question is whether the grain boundaries of other ionic conductors, such as proton conductors, are affected in a similar manner. Examples discussed in the literature so far include Ba3Ca1.18Nb1.82O9−δ where it is claimed that the grain boundary conductivity is electronic (p-type) under conditions where the grain interior conductivity is purely protonic [13]. For acceptor-doped BaZrO3 it has become clear that it exhibits high grain interior proton conductivity [14], while the grain boundaries are very resistive. Moreover, it appears that electrons contribute significantly to the grain boundary conductivity in undoped BaZrO3 under conditions where the grain interior conductivity is essentially protonic [15]. Recently, we have reported on new classes of proton conducting acceptor doped ternary oxides, including rare earth niobates [3,16,17], tantalates [3,18], tungstates [19,20] and titanates [21]. In general, these materials conduct protons at elevated temperature, and, additionally, oxide ions and electronic charge carriers contribute to the conductivity to different extents. Upon studying the conductive behaviour along the series of the pyrochlore structured rare earth titanates (Ln2Ti2O7), an interesting observation was made for Ca-doped Er2Ti2O7: at low temperatures, essentially no difference in the ac conductivity (10 kHz) was observed when comparing wet and dry atmospheres, despite that measurements of the open circuit voltage (OCV) in a water vapour concentration cell indicated that protons were the major charge carrier. Also, impedance spectroscopy indicated that the “high” and “medium” frequency contributions behaved differently upon changes in measurements conditions, notably pH2O. Consequently, these initial data suggest that a more complex model, including different transport characteristics in grain interior and grain boundaries, should be used to explain the conductivity data recorded for acceptor-doped Er2Ti2O7.
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H. Fjeld et al. / Solid State Ionics 179 (2008) 1849–1853
vapour (both H2O and D2O) and oxygen partial pressures in the temperature range 360–600°C. The results from the electrical measurements will be discussed in terms of a point-defect chemical model and sample microstructure. 2. Experimental
Fig. 1. Impedance spectra recorded at 400 °C (A) and 550 °C (B) in O2 as a function of pH2O. Numbers indicate log frequency of the low pH2O spectra. The inset in (A) shows the high frequency part of the spectra. The equivalent circuit used to fit the data is also shown schematically.
In order to reach a better understanding of the electrical properties of grain interior and grain boundaries of Er1.96Ca0.04Ti2O6.98, we have systematically recorded impedance spectra as a function of water
Powders of Er1.96Ca0.04Ti2O6.98 and Er2Ti2O7 were synthesized by a standard ceramic method using Er2O3 (Molycorp, 98%, Ho main impurity), CaCO3 (Sigma-Aldrich, 99.95%) and TiO2 (Alfa Aesar, 99.5%) as reagents. To promote good mixing of Ca on Er-site for the acceptor doped sample, Er2O3 and CaCO3 were dissolved in hot dilute nitric acid, with citric acid as complexing agent. After decomposition of the citric acid, the viscous solution was heated to 700°C after which the powder was pink. TiO2 was then added in stoichiometric amounts, and the mixture was ballmilled and then calcined at 800°C for 15 h. Ballmilling and calcination (900 and 1000°C) were repeated twice. Dense samples (90–95% of theoretical density) were obtained by uniaxially pressing green disks of 25 mm diameter followed by sintering at 1500°C for 5 h. The crystal structures of Er1.96Ca0.04Ti2O6.98 and Er2Ti2O7 were characterized by means of X-ray powder diffraction using a Siemens D5000 powder diffractometer with CuKα1 radiation. The lattice parameters were refined from the obtained patterns in space group — Fd 3m (number 227) with Si as an internal standard using the FullProf Software [22] operating in a profile matching mode. The impedance measurements (oscillation voltage = 0.5 V RMS) were performed using Solartron 1260 FRA on samples with Ptelectrodes (Metalor A3788A paint + Pt mesh) in the frequency range 2 MHz–0.05 Hz. The gas atmosphere in the measurement cell (ProboStat, NorECs, Norway) was controlled using an in-house-built gas mixer [23]. Only oxidising conditions were applied (i.e. 10− 5 atm b pO2 b 1 atm). Controlled H2O or D2O contents were obtained by passing the gas mixture through saturated KBr (aq). Upon changing the gas composition, the conductivity was recorded at two frequencies, “high” and “low”, to assure that equilibrium was obtained throughout the microstructure of the sample. The impedance spectra were fitted to the equivalent circuit L(R1Q1) (R2Q2)(R3Q3)(R4Q4), where (RiQi) denotes a resistor and a constant phase element in parallel. Above 500°C (R1Q1) was replaced with R1, and below 400°C (R4Q4) was excluded. An inductive element (L) was used throughout the fitting of the data, mostly reflecting parasitic
Fig. 2. Rgi (A), R2 (B) and R3 (C) vs inverse temperature in O2/H2O (pH2O = 0.025 atm), O2/D2O (pD2O = 0.025 atm) and dry O2 (pH2O ~ 50 ppm H2O).
H. Fjeld et al. / Solid State Ionics 179 (2008) 1849–1853
apparatus contributions at high frequencies. To relate the different contributions of the impedance spectra to physical processes, the capacitances, Ci, of the sub-circuits were calculated using [24] 1 n
1 n
−1
Ci ¼ Yi i Ri i ;
ð1Þ
where Yi and ni define the admittance offfi the constant phase element pffiffiffiffiffi according to YCPE = Y(jω)n where j ¼ −1 and ω = angular frequency. For more details concerning equivalent circuits in solid state ionics, including constant phase elements, consult e.g. [25]. The microstructure and chemical compositions of the dense samples were investigated before and after the impedance study by means of analytical transmission electron microscopy (TEM). The microscope used was a JEOL2010F with a field emission gun and a Noran Vantage DI+ energy dispersive X-ray spectroscopy (EDS) system. 3. Results and discussion Fig. 1 (A and B) show the impedance spectra recorded for Er1.96Ca0.04Ti2O6.98 as a function of pH2O in O2 at 400°C (A) and 550°C (B); the inset in Fig. 1A displays the high-frequency part in more
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detail. The total impedance increases with decreasing water vapour pressure, whereas the high frequency part remains essentially constant. By fitting the impedance spectra into sub-circuits involving R1, R2, R3 and R4, as illustrated in Fig. 1 (A and B), geometrical capacitances were calculated with Eq. (1) to C1 ~ 5.5 ± 1 pF cm− 1, C2 ~ 2.5 ± 1 nF cm− 1, C3 ~ 25 ± 10 nF cm− 1 and C4 N 20 µF cm− 1 for the experimental conditions in this study. C1 reflects a value typical of grain interiors, and R1 is from now on denoted Rgi (grain interior). The values of C2 and C3 are within the range expected to represent grain boundary capacitances in polycrystalline samples of typical grain size [25]. It should be noted that the spectra obtained generally displayed two distinguishable contributions related to C2 and C3, except under the driest conditions where the two contributions were overlapping into one deformed semi-circle. The change in the spectra with temperature and atmosphere were reversible. Finally, we assigned C4 to electrode related processes for which it was possible to extract reasonable values for temperatures ≥ 400°C. Fig. 2 present Rgi (A), R2 (B) and R3 (C) as a function of inverse temperature in H2O/O2, D2O/O2 and dry O2. No change in grain interior resistivity was observed upon changing the atmosphere, indicating that protons do not significantly contribute as charge carriers. In addition, the grain interior resistivity was independent of pO2 (not shown here)
Fig. 3. R2 (A) and R3 (B) as a function of water vapour pressure in O2. R2 (C) and R3 (D) as a function of oxygen partial pressure in wet atmosphere (pH2O = 0.025 atm).
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Table 1 Composition (atomic percent of cations) of grain interior, grain boundary and triple grain junction in Fig. 4, measured by EDS
Grain interior, big grain (A) Grain interior, small grain (B) Grain boundary (C) Triple grain junction (D)
Er
Ti
Ca
Si
53 54 46 16
46 44 40 3
1 1 2 9
b1 1 12 72
within the present experimental window. On basis of this, one may conclude that oxygen vacancies are the major charge carriers in the grain interiors of Er1.96Ca0.04Ti2O6.98, charge compensating the acceptor Ca/Er. On the other hand, R2 and R3 (cf. Fig. 2B and C) were dependent on pH2O and hydrogen isotope, indicating that protons are important as charge carriers in the grain boundaries. The water vapour and oxygen partial pressure dependencies of R2 and R3 are displayed in Fig. 3 (A–D). R2 decreases with increasing pH2O, which is similar to what is typically observed for high temperature proton conductors. One should also note in this respect that the relative dependence of R2 on changes in pH2O decreases with increasing temperature. This type of behaviour is in agreement with a model where oxygen vacancies are hydrated exothermically,
Fig. 5. Profile matching of Er1.96Ca0.04Ti2O6.98 (and Si as internal standard, labelled with ⁎ in the figure) with recorded pattern in black, fitted pattern in grey and the difference shown below.
where Acc/ denotes an effective negatively charged acceptor. This results in a gradual decrease in the proton concentration with increasing temperature. That protons play the major role as charge carriers in the transport process reflected by R2 is, furthermore, supported by Fig. 2B showing a clear H+/D+ isotope effect. In the temperature range covered here, the ratio of the resistivities measured in D2O and H2O varied from 1.3 to 1.6, increasing with decreasing temperature. R3 shows a more complex behaviour than R2 as a function of the water vapour pressure, as can be seen by comparing Fig. 3A and B. R3 decreases with pH2O at low temperatures and increases with pH2O at high temperatures, the slope gradually changing from being negative to positive with increasing temperature. R3 exhibits, like R2, (cf. Fig. 2B and C) a clear isotope effect at low temperature, which disappears at higher temperatures indicating that protons contribute less as charge carriers at high temperatures. Such a complex behaviour can again be explained in light of Eqs. (2) and (3). At low temperatures protons dominate as charge carriers so that an increase in pH2O reduces the resistivity. In contrast, at high temperatures the oxygen vacancies are the major charge carriers so
that increasing pH2O increases the resistivity by reducing the concentration of oxygen vacancies. One should, however, note that the shift in type of charge carrier that determines the resistivity can only be rationalised in this way if the enthalpy of mobility of oxygen vacancies is higher than that of protons. One may argue that electron holes, instead of oxygen vacancies, could be the major charge carrier determining the resistivity reflected by R3 at high temperatures. However, taking the essentially pO2-independent behaviour of R3 into account (cf. Fig. 3D), this seems unlikely as one would expect that R3 should be proportional to pO−2 1/4. This is also in agreement with OCV measurements (not shown here), which indicate an overall ionic transport number higher than 0.95 at 600°C. As discussed above and shown in Fig. 1 (A and B), the impedance spectra of the samples display two semi-circles with capacitances in the “grain boundary range” and the corresponding resistivities were pH2O dependant. In order to understand this complex behaviour, the chemistry of the grain boundaries was investigated by EDS in a TEM. Table 1 lists the chemical composition of grain interior, a grain boundary and a triple grain junction, from the corresponding areas marked in Fig. 4. The chemical analyses clearly show that the triple grain junction is richer in both Si and Ca than the grain interior and, additionally, that also the grain boundary contains significantly higher concentrations of Si and Ca. More areas of the grain boundaries were analysed and it was in general found that the concentration of Si and Ca increased towards the triple grain junctions. Moreover, no significant difference in the chemistry could be detected in the samples prior to and after the
Fig. 4. TEM micrograph of typical grain structure of Er1.96Ca0.04Ti2O6.98.
Fig. 6. SAD from a triple grain junction showing diffuse rings from the amorphous phase and reflections from the neighbouring Er1.96Ca0.04Ti2O6.98-phase.
H2 OðgÞ þ VO þOXO ¼ 2OHO
ð2Þ
and where the electroneutrality condition is given by i h OHO þ 2 VO ¼ Acc=
ð3Þ
H. Fjeld et al. / Solid State Ionics 179 (2008) 1849–1853
impedance study. The presence of Si in the grain boundaries may be a result of the powder preparation where Er2O3 and CaCO3 were dissolved in hot dilute nitric acid in a borosilicate glass beaker or of the milling of the powders in an agate ballmill. The XRD pattern and the result of the profile matching for Er1.96Ca0.04Ti2O6.98 (and Si as internal standard) are shown in Fig. 5. The lattice parameters were refined to 10.072(2) Å and 10.071(2) Å for the Ca-doped and nominally undoped sample, respectively. This is in agreement with the reported values of a = 10.076Å [26,27]. Even though no conclusion can be made from the profile matching on whether the lattice parameter changed as a result of the Ca addition, the EDS results (Table 1) show that Ca is present in the grain interior. XRD showed no evidence of secondary phases, however, selected area diffraction (SAD) from triple grain junctions showed diffuse rings indicating the presence of an amorphous phase (Fig. 6). High resolution electron microscopy supports this observation and it is therefore concluded that the Si- and Ca-enriched areas found in the triple grain junctions and along the grain boundaries are amorphous. Since the additional phase is amorphous and has a much lower volume fraction than the crystalline Er2Ti2O7 phase it is difficult to detect this phase by XRD. Based on the TEM-characterization, it is appropriate to ask how the additional phase detected in the grain boundaries influences the electrical properties of the grain boundaries. An insulating phase partially covering the grain boundaries exhibits in principle no response to changes in e.g. water vapour pressure. This is in contrast to what is encountered here, and we therefore suggest that the role of the siliceous phases is different in acceptor doped Er2Ti2O7 than what is observed for e.g. ZrO2 and CeO2, the siliceous phase in our case perhaps being a proton conductor. Interestingly, crystalline Al-doped Ca3(SiO4)O, which contains elements as were observed in the grain boundaries of our material (Table 1), has recently been reported as a high temperature proton conductor in wet atmospheres [28]. Crystalline Sr-doped La2Si2O7 is another silicate with reported high temperature proton conductivity [29]. As a final example from the literature it is worth mentioning LaPO4, in which amorphous phosphorous-rich films present in the grain boundaries exhibit proton conductivity at high temperatures [30]. According to the grain boundary core-space charge layer model, the depletion of doubly charged defects (e.g. oxygen vacancies) is quadratic compared to singly charged defects (e.g. protons) [9]. In a material containing both protons and oxygen vacancies, the space charge layer effect should therefore result in the ratio of the concentrations of protons to oxygen vacancies becoming higher in the space charge layers than in grain interiors. This may, alternatively or additionally to the siliceous phase, explain why protons play a more important role in the conductive behaviour of the grain boundaries than in the grain interiors of Er1.96Ca0.04Ti2O6.98. Another interesting feature in the impedance spectra (Fig. 1A and B) is the fact that two frequency-resolved semi-circles exhibit capacitances corresponding to the “grain boundary range”. For mixed ionic and electronic conductors such as SrTiO3 [31] and CeO2 [32,33] it has been experimentally confirmed that diffusion-related features appear in impedance spectra in addition to the “traditional” ones resulting from grain interior, grain boundaries and electrodes. This feature can be explained in terms of the chemical capacitance associated to the coupling between the different current flow rails, i.e. the ionic and electronic rails. The effect is similar to what is found when employing selectively blocking electrodes (Hebb–Wagner polarization) and may result in a tear-drop shaped “semi-circle” in a Nyquist-plot [34]. On basis of this, one may speculate whether a similar feature can appear in the case of mixed protonic and oxide ion conduction and that this may explain the R3-contribution. Finally, one may reflect on whether similar effects from the grain boundaries on the overall ionic transport play a role in other materials. The samples investigated in this study were prepared by traditional
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ceramic methods, resulting in this case in an additional siliceous phase unintentionally present (and not observed in XRD) at the grain boundaries. Notably, when measuring the resistivity at a fixed frequency, it is easy to overlook that the grain interior and grain boundaries respond differently to changes in e.g. partial pressures, which indeed emphasises the use of impedance spectroscopy. 4. Conclusions The study of conventionally prepared samples of Er1.96Ca0.04Ti2O6.98 can be summarized as follows: 1. TEM showed that an amorphous phase, enriched in Si and Ca, was present in the triple grain junctions and along the grain boundaries. 2. The impedance spectra contained two contributions with capacitances in the range typically assigned to grain boundary processes. 3. While oxide ions were identified as the major charge carriers in the grain interiors, protons were found to play an important role as charge carriers in the grain boundary related processes. 4. The behaviour of the two resistivities associated with the grain boundary processes could be explained by applying a defect chemical model in which protons and oxygen vacancies are the major charge carriers. Acknowledgements RH acknowledges funding by the Research Council of Norway, Grant No. 15851/431 (Functional Oxides for Energy Technology, NANOMAT). References [1] T. Ishihara, N.M. Sammes, O. Yamamoto, High Temperature Solid Oxide Fuel Cells, Elsevier, Oxford, 2003, p. 83. [2] H. Iwahara, T. Esaka, H. Uchida, N. Maeda, Solid State Ionics 3–4 (1981) 359. [3] R. Haugsrud, T. Norby, Nature Materials 5 (2006) 193. [4] J.E. Bauerle, J. Phys. Chem. Solids 30 (1969) 2657. [5] S.P.S. Badwal, J. Drennan, Solid State Ionics 40–41 (1990) 869. [6] J. Fleig, J. Maier, J. Am. Ceram. Soc. 82 (1999) 3485. [7] X. Guo, J. Maier, J. Electrochem. Soc. 148 (2001) E121. [8] X. Guo, W. Sigle, J. Fleig, J. Maier, Solid State Ionics 154–155 (2002) 555. [9] X. Guo, R. Waser, Prog. Mater. Sci. 51 (2006) 151. [10] X. Guo, W. Sigle, J. Maier, J. Am. Ceram. Soc. 86 (2003) 77. [11] X. Guo, J. Fleig, J. Maier, J. Electrochem. Soc. 148 (2001) J50. [12] X. Guo, J. Fleig, J. Maier, Solid State Ionics 154–155 (2002) 563. [13] H.G. Bohn, T. Schober, T. Mono, W. Schilling, Solid State Ionics 117 (1999) 219. [14] H.G. Bohn, T. Schober, J. Am. Ceram. Soc. 83 (2000) 768. [15] C. Kjølseth, H. Fjeld, P.I. Dahl, C. Estournès, R. Haugsrud, T. Norby, submitted to Solid State Ionics. [16] R. Haugsrud, T. Norby, Solid State Ionics 177 (2006) 1129. [17] R. Haugsrud, B. Ballesteros, M. Lira-Cantu, T. Norby, J. Electrochem. Soc. 153 (2006) J87. [18] R. Haugsrud, T. Norby, J. Am. Ceram. Soc. 90 (2007) 1116. [19] R. Haugsrud, H. Fjeld, K.R. Haug, T. Norby, J. Electrochem. Soc. 154 (2007) B77. [20] R. Haugsrud, Solid State Ionics 178 (2007) 555. [21] R. Haugsrud, T. Norby, Solid State Electrochemistry, Proceedings of the Risoe International Symposium on Materials Science, 26th, Roskilde, Denmark, Sept. 4–8, 2005, 2005, p. 209. [22] J. Rodríguez-Carvajal, FullProf Ver. Feb-2007, 2007. [23] T. Norby, Solid State Ionics 28–30 (1988) 1586. [24] T. Jacobsen, B. Zachau-Christiansen, L. Bay, S. Skaarup, Proc. 17th Risø International Symposium on “High Temperature Electrochemistry: Ceramics and Metals”, Roskilde, Denmark, 1996, p. 29. [25] S.M. Haile, D.L. West, J. Campbell, J. Mater. Res. 13 (1998) 1576. [26] O. Knop, F. Brisse, L. Castelliz, Sutarno, Canadian Journal of Chemistry 43 (1965) 2812. [27] G.C. Lau, B.D. Muegge, T.M. McQueen, E.L. Duncan, R.J. Cava, J. Solid State Chem. 179 (2006) 3126. [28] J.M. Porras-Vazquez, A.G. De la Torre, E.R. Losilla, M.A.G. Aranda, Solid State Ionics 178 (2007) 1073. [29] K. Amezawa, J.-i. Yamada, N. Kitamura, Y. Tomii, T. Handa, N. Yamamoto, Solid State Ionics 176 (2005) 341. [30] G. Harley, R. Yu, L.C. De Jonghe, Solid State Ionics 178 (2007) 769. [31] J. Jamnik, X. Guo, J. Maier, Appl. Phys. Lett. 82 (2003) 2820. [32] P. Jasinski, V. Petrovsky, T. Suzuki, H.U. Anderson, J. Electrochem. Soc. 152 (2005) J27. [33] W. Lai, S.M. Haile, J. Am. Ceram. Soc. 88 (2005) 2979. [34] J. Jamnik, Solid State Ionics 157 (2003) 19.
Contents lists available at ScienceDirect
Solid State Ionics j o u r n a l h o m e p a g e : w w w. e l s e v i e r. c o m / l o c a t e / s s i
Proton and oxide ion conductivity in grain boundaries and grain interior of Ca-doped Er2Ti2O7 with Si-impurities Harald Fjeld a,⁎, Reidar Haugsrud a, Anette Eleonora Gunnæs b, Truls Norby a a b
Department of Chemistry, University of Oslo, FERMiO, Gaustadalléen 21, NO-0349 Oslo, Norway Department of Physics, University of Oslo, FERMiO, Gaustadalléen 21, NO-0349 Oslo, Norway
A R T I C L E
I N F O
Article history: Received 8 December 2007 Received in revised form 9 May 2008 Accepted 3 June 2008 Keywords: Er2Ti2O7 Grain boundaries Ionic conductivity Impedance spectroscopy
A B S T R A C T Dense samples of nominal composition Er1.96Ca0.04Ti2O6.98 were prepared by conventional ceramic synthesis to evaluate the ionic conductivity. The electrical properties were studied by impedance spectroscopy under oxidizing conditions as a function of pO2, pH2O and pD2O in the temperature range 360–600°C. The microstructure was investigated by analytical transmission electron microscopy before and after the impedance study. In both cases the grain boundaries contained an amorphous phase, enriched in Si and Ca. Three different contributions to the sample impedance, with different time constants, were identified, one assigned to the grain interior and two to grain boundaries. For the grain interior, oxygen vacancies were found to be the major charge carrier. In contrast, protons were found to play an important role as charge carriers in the grain boundary related processes. The ionic transport over grain boundaries in Er1.96Ca0.04Ti2O6.98 is discussed in terms of effects of additional phases in the grain boundaries and space charge layers. © 2008 Elsevier B.V. All rights reserved.
1. Introduction Acceptor-doped oxides are known to exhibit transport of ionic charge carriers such as oxide ions (in e.g. ZrO2 and CeO2 [1]) and protons (in e.g. SrCeO3 [2] and LaNbO4 [3]) at elevated temperatures, and find possible applications in electrochemical devices for future energy technology. In a polycrystalline material, both the grain interior and the grain boundaries contribute to the overall electrical properties. This implies that even though the grain interior of the material is an excellent ionic conductor, the practical use of the material may be limited by a high grain boundary resistance. In-depth knowledge of the transport of charge carriers in both the grain interior and grain boundaries is therefore of great importance. The detrimental role of the grain boundaries is known for e.g. acceptor-doped ZrO2, where the overall ionic conductivity is decreased due to the grain boundary resistance [4]. The grain boundary resistance can be divided into (at least) two contributions: 1) intergranular siliceous phases partially covering the grain boundary area [5], leading to constriction of current lines [6], and 2) depletion of oxygen vacancies in the vicinity of the grain boundaries; the space charge layer effect [7,8]. The grain boundary core-space charge layer model [9] has also been successfully applied to explain why electrons, in addition to oxygen vacancies, contribute to the grain boundary conductivity in acceptordoped CeO2 [10] and why electron holes dominate the grain boundary conductivity in Fe-doped SrTiO3 [11,12] under conditions where the grain interior conductivity is essentially ionic. ⁎ Corresponding author. Tel.: +47 22840657; fax: +47 22840651. E-mail address: [email protected] (H. Fjeld). 0167-2738/$ – see front matter © 2008 Elsevier B.V. All rights reserved. doi:10.1016/j.ssi.2008.06.003
An interesting question is whether the grain boundaries of other ionic conductors, such as proton conductors, are affected in a similar manner. Examples discussed in the literature so far include Ba3Ca1.18Nb1.82O9−δ where it is claimed that the grain boundary conductivity is electronic (p-type) under conditions where the grain interior conductivity is purely protonic [13]. For acceptor-doped BaZrO3 it has become clear that it exhibits high grain interior proton conductivity [14], while the grain boundaries are very resistive. Moreover, it appears that electrons contribute significantly to the grain boundary conductivity in undoped BaZrO3 under conditions where the grain interior conductivity is essentially protonic [15]. Recently, we have reported on new classes of proton conducting acceptor doped ternary oxides, including rare earth niobates [3,16,17], tantalates [3,18], tungstates [19,20] and titanates [21]. In general, these materials conduct protons at elevated temperature, and, additionally, oxide ions and electronic charge carriers contribute to the conductivity to different extents. Upon studying the conductive behaviour along the series of the pyrochlore structured rare earth titanates (Ln2Ti2O7), an interesting observation was made for Ca-doped Er2Ti2O7: at low temperatures, essentially no difference in the ac conductivity (10 kHz) was observed when comparing wet and dry atmospheres, despite that measurements of the open circuit voltage (OCV) in a water vapour concentration cell indicated that protons were the major charge carrier. Also, impedance spectroscopy indicated that the “high” and “medium” frequency contributions behaved differently upon changes in measurements conditions, notably pH2O. Consequently, these initial data suggest that a more complex model, including different transport characteristics in grain interior and grain boundaries, should be used to explain the conductivity data recorded for acceptor-doped Er2Ti2O7.
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H. Fjeld et al. / Solid State Ionics 179 (2008) 1849–1853
vapour (both H2O and D2O) and oxygen partial pressures in the temperature range 360–600°C. The results from the electrical measurements will be discussed in terms of a point-defect chemical model and sample microstructure. 2. Experimental
Fig. 1. Impedance spectra recorded at 400 °C (A) and 550 °C (B) in O2 as a function of pH2O. Numbers indicate log frequency of the low pH2O spectra. The inset in (A) shows the high frequency part of the spectra. The equivalent circuit used to fit the data is also shown schematically.
In order to reach a better understanding of the electrical properties of grain interior and grain boundaries of Er1.96Ca0.04Ti2O6.98, we have systematically recorded impedance spectra as a function of water
Powders of Er1.96Ca0.04Ti2O6.98 and Er2Ti2O7 were synthesized by a standard ceramic method using Er2O3 (Molycorp, 98%, Ho main impurity), CaCO3 (Sigma-Aldrich, 99.95%) and TiO2 (Alfa Aesar, 99.5%) as reagents. To promote good mixing of Ca on Er-site for the acceptor doped sample, Er2O3 and CaCO3 were dissolved in hot dilute nitric acid, with citric acid as complexing agent. After decomposition of the citric acid, the viscous solution was heated to 700°C after which the powder was pink. TiO2 was then added in stoichiometric amounts, and the mixture was ballmilled and then calcined at 800°C for 15 h. Ballmilling and calcination (900 and 1000°C) were repeated twice. Dense samples (90–95% of theoretical density) were obtained by uniaxially pressing green disks of 25 mm diameter followed by sintering at 1500°C for 5 h. The crystal structures of Er1.96Ca0.04Ti2O6.98 and Er2Ti2O7 were characterized by means of X-ray powder diffraction using a Siemens D5000 powder diffractometer with CuKα1 radiation. The lattice parameters were refined from the obtained patterns in space group — Fd 3m (number 227) with Si as an internal standard using the FullProf Software [22] operating in a profile matching mode. The impedance measurements (oscillation voltage = 0.5 V RMS) were performed using Solartron 1260 FRA on samples with Ptelectrodes (Metalor A3788A paint + Pt mesh) in the frequency range 2 MHz–0.05 Hz. The gas atmosphere in the measurement cell (ProboStat, NorECs, Norway) was controlled using an in-house-built gas mixer [23]. Only oxidising conditions were applied (i.e. 10− 5 atm b pO2 b 1 atm). Controlled H2O or D2O contents were obtained by passing the gas mixture through saturated KBr (aq). Upon changing the gas composition, the conductivity was recorded at two frequencies, “high” and “low”, to assure that equilibrium was obtained throughout the microstructure of the sample. The impedance spectra were fitted to the equivalent circuit L(R1Q1) (R2Q2)(R3Q3)(R4Q4), where (RiQi) denotes a resistor and a constant phase element in parallel. Above 500°C (R1Q1) was replaced with R1, and below 400°C (R4Q4) was excluded. An inductive element (L) was used throughout the fitting of the data, mostly reflecting parasitic
Fig. 2. Rgi (A), R2 (B) and R3 (C) vs inverse temperature in O2/H2O (pH2O = 0.025 atm), O2/D2O (pD2O = 0.025 atm) and dry O2 (pH2O ~ 50 ppm H2O).
H. Fjeld et al. / Solid State Ionics 179 (2008) 1849–1853
apparatus contributions at high frequencies. To relate the different contributions of the impedance spectra to physical processes, the capacitances, Ci, of the sub-circuits were calculated using [24] 1 n
1 n
−1
Ci ¼ Yi i Ri i ;
ð1Þ
where Yi and ni define the admittance offfi the constant phase element pffiffiffiffiffi according to YCPE = Y(jω)n where j ¼ −1 and ω = angular frequency. For more details concerning equivalent circuits in solid state ionics, including constant phase elements, consult e.g. [25]. The microstructure and chemical compositions of the dense samples were investigated before and after the impedance study by means of analytical transmission electron microscopy (TEM). The microscope used was a JEOL2010F with a field emission gun and a Noran Vantage DI+ energy dispersive X-ray spectroscopy (EDS) system. 3. Results and discussion Fig. 1 (A and B) show the impedance spectra recorded for Er1.96Ca0.04Ti2O6.98 as a function of pH2O in O2 at 400°C (A) and 550°C (B); the inset in Fig. 1A displays the high-frequency part in more
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detail. The total impedance increases with decreasing water vapour pressure, whereas the high frequency part remains essentially constant. By fitting the impedance spectra into sub-circuits involving R1, R2, R3 and R4, as illustrated in Fig. 1 (A and B), geometrical capacitances were calculated with Eq. (1) to C1 ~ 5.5 ± 1 pF cm− 1, C2 ~ 2.5 ± 1 nF cm− 1, C3 ~ 25 ± 10 nF cm− 1 and C4 N 20 µF cm− 1 for the experimental conditions in this study. C1 reflects a value typical of grain interiors, and R1 is from now on denoted Rgi (grain interior). The values of C2 and C3 are within the range expected to represent grain boundary capacitances in polycrystalline samples of typical grain size [25]. It should be noted that the spectra obtained generally displayed two distinguishable contributions related to C2 and C3, except under the driest conditions where the two contributions were overlapping into one deformed semi-circle. The change in the spectra with temperature and atmosphere were reversible. Finally, we assigned C4 to electrode related processes for which it was possible to extract reasonable values for temperatures ≥ 400°C. Fig. 2 present Rgi (A), R2 (B) and R3 (C) as a function of inverse temperature in H2O/O2, D2O/O2 and dry O2. No change in grain interior resistivity was observed upon changing the atmosphere, indicating that protons do not significantly contribute as charge carriers. In addition, the grain interior resistivity was independent of pO2 (not shown here)
Fig. 3. R2 (A) and R3 (B) as a function of water vapour pressure in O2. R2 (C) and R3 (D) as a function of oxygen partial pressure in wet atmosphere (pH2O = 0.025 atm).
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Table 1 Composition (atomic percent of cations) of grain interior, grain boundary and triple grain junction in Fig. 4, measured by EDS
Grain interior, big grain (A) Grain interior, small grain (B) Grain boundary (C) Triple grain junction (D)
Er
Ti
Ca
Si
53 54 46 16
46 44 40 3
1 1 2 9
b1 1 12 72
within the present experimental window. On basis of this, one may conclude that oxygen vacancies are the major charge carriers in the grain interiors of Er1.96Ca0.04Ti2O6.98, charge compensating the acceptor Ca/Er. On the other hand, R2 and R3 (cf. Fig. 2B and C) were dependent on pH2O and hydrogen isotope, indicating that protons are important as charge carriers in the grain boundaries. The water vapour and oxygen partial pressure dependencies of R2 and R3 are displayed in Fig. 3 (A–D). R2 decreases with increasing pH2O, which is similar to what is typically observed for high temperature proton conductors. One should also note in this respect that the relative dependence of R2 on changes in pH2O decreases with increasing temperature. This type of behaviour is in agreement with a model where oxygen vacancies are hydrated exothermically,
Fig. 5. Profile matching of Er1.96Ca0.04Ti2O6.98 (and Si as internal standard, labelled with ⁎ in the figure) with recorded pattern in black, fitted pattern in grey and the difference shown below.
where Acc/ denotes an effective negatively charged acceptor. This results in a gradual decrease in the proton concentration with increasing temperature. That protons play the major role as charge carriers in the transport process reflected by R2 is, furthermore, supported by Fig. 2B showing a clear H+/D+ isotope effect. In the temperature range covered here, the ratio of the resistivities measured in D2O and H2O varied from 1.3 to 1.6, increasing with decreasing temperature. R3 shows a more complex behaviour than R2 as a function of the water vapour pressure, as can be seen by comparing Fig. 3A and B. R3 decreases with pH2O at low temperatures and increases with pH2O at high temperatures, the slope gradually changing from being negative to positive with increasing temperature. R3 exhibits, like R2, (cf. Fig. 2B and C) a clear isotope effect at low temperature, which disappears at higher temperatures indicating that protons contribute less as charge carriers at high temperatures. Such a complex behaviour can again be explained in light of Eqs. (2) and (3). At low temperatures protons dominate as charge carriers so that an increase in pH2O reduces the resistivity. In contrast, at high temperatures the oxygen vacancies are the major charge carriers so
that increasing pH2O increases the resistivity by reducing the concentration of oxygen vacancies. One should, however, note that the shift in type of charge carrier that determines the resistivity can only be rationalised in this way if the enthalpy of mobility of oxygen vacancies is higher than that of protons. One may argue that electron holes, instead of oxygen vacancies, could be the major charge carrier determining the resistivity reflected by R3 at high temperatures. However, taking the essentially pO2-independent behaviour of R3 into account (cf. Fig. 3D), this seems unlikely as one would expect that R3 should be proportional to pO−2 1/4. This is also in agreement with OCV measurements (not shown here), which indicate an overall ionic transport number higher than 0.95 at 600°C. As discussed above and shown in Fig. 1 (A and B), the impedance spectra of the samples display two semi-circles with capacitances in the “grain boundary range” and the corresponding resistivities were pH2O dependant. In order to understand this complex behaviour, the chemistry of the grain boundaries was investigated by EDS in a TEM. Table 1 lists the chemical composition of grain interior, a grain boundary and a triple grain junction, from the corresponding areas marked in Fig. 4. The chemical analyses clearly show that the triple grain junction is richer in both Si and Ca than the grain interior and, additionally, that also the grain boundary contains significantly higher concentrations of Si and Ca. More areas of the grain boundaries were analysed and it was in general found that the concentration of Si and Ca increased towards the triple grain junctions. Moreover, no significant difference in the chemistry could be detected in the samples prior to and after the
Fig. 4. TEM micrograph of typical grain structure of Er1.96Ca0.04Ti2O6.98.
Fig. 6. SAD from a triple grain junction showing diffuse rings from the amorphous phase and reflections from the neighbouring Er1.96Ca0.04Ti2O6.98-phase.
H2 OðgÞ þ VO þOXO ¼ 2OHO
ð2Þ
and where the electroneutrality condition is given by i h OHO þ 2 VO ¼ Acc=
ð3Þ
H. Fjeld et al. / Solid State Ionics 179 (2008) 1849–1853
impedance study. The presence of Si in the grain boundaries may be a result of the powder preparation where Er2O3 and CaCO3 were dissolved in hot dilute nitric acid in a borosilicate glass beaker or of the milling of the powders in an agate ballmill. The XRD pattern and the result of the profile matching for Er1.96Ca0.04Ti2O6.98 (and Si as internal standard) are shown in Fig. 5. The lattice parameters were refined to 10.072(2) Å and 10.071(2) Å for the Ca-doped and nominally undoped sample, respectively. This is in agreement with the reported values of a = 10.076Å [26,27]. Even though no conclusion can be made from the profile matching on whether the lattice parameter changed as a result of the Ca addition, the EDS results (Table 1) show that Ca is present in the grain interior. XRD showed no evidence of secondary phases, however, selected area diffraction (SAD) from triple grain junctions showed diffuse rings indicating the presence of an amorphous phase (Fig. 6). High resolution electron microscopy supports this observation and it is therefore concluded that the Si- and Ca-enriched areas found in the triple grain junctions and along the grain boundaries are amorphous. Since the additional phase is amorphous and has a much lower volume fraction than the crystalline Er2Ti2O7 phase it is difficult to detect this phase by XRD. Based on the TEM-characterization, it is appropriate to ask how the additional phase detected in the grain boundaries influences the electrical properties of the grain boundaries. An insulating phase partially covering the grain boundaries exhibits in principle no response to changes in e.g. water vapour pressure. This is in contrast to what is encountered here, and we therefore suggest that the role of the siliceous phases is different in acceptor doped Er2Ti2O7 than what is observed for e.g. ZrO2 and CeO2, the siliceous phase in our case perhaps being a proton conductor. Interestingly, crystalline Al-doped Ca3(SiO4)O, which contains elements as were observed in the grain boundaries of our material (Table 1), has recently been reported as a high temperature proton conductor in wet atmospheres [28]. Crystalline Sr-doped La2Si2O7 is another silicate with reported high temperature proton conductivity [29]. As a final example from the literature it is worth mentioning LaPO4, in which amorphous phosphorous-rich films present in the grain boundaries exhibit proton conductivity at high temperatures [30]. According to the grain boundary core-space charge layer model, the depletion of doubly charged defects (e.g. oxygen vacancies) is quadratic compared to singly charged defects (e.g. protons) [9]. In a material containing both protons and oxygen vacancies, the space charge layer effect should therefore result in the ratio of the concentrations of protons to oxygen vacancies becoming higher in the space charge layers than in grain interiors. This may, alternatively or additionally to the siliceous phase, explain why protons play a more important role in the conductive behaviour of the grain boundaries than in the grain interiors of Er1.96Ca0.04Ti2O6.98. Another interesting feature in the impedance spectra (Fig. 1A and B) is the fact that two frequency-resolved semi-circles exhibit capacitances corresponding to the “grain boundary range”. For mixed ionic and electronic conductors such as SrTiO3 [31] and CeO2 [32,33] it has been experimentally confirmed that diffusion-related features appear in impedance spectra in addition to the “traditional” ones resulting from grain interior, grain boundaries and electrodes. This feature can be explained in terms of the chemical capacitance associated to the coupling between the different current flow rails, i.e. the ionic and electronic rails. The effect is similar to what is found when employing selectively blocking electrodes (Hebb–Wagner polarization) and may result in a tear-drop shaped “semi-circle” in a Nyquist-plot [34]. On basis of this, one may speculate whether a similar feature can appear in the case of mixed protonic and oxide ion conduction and that this may explain the R3-contribution. Finally, one may reflect on whether similar effects from the grain boundaries on the overall ionic transport play a role in other materials. The samples investigated in this study were prepared by traditional
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ceramic methods, resulting in this case in an additional siliceous phase unintentionally present (and not observed in XRD) at the grain boundaries. Notably, when measuring the resistivity at a fixed frequency, it is easy to overlook that the grain interior and grain boundaries respond differently to changes in e.g. partial pressures, which indeed emphasises the use of impedance spectroscopy. 4. Conclusions The study of conventionally prepared samples of Er1.96Ca0.04Ti2O6.98 can be summarized as follows: 1. TEM showed that an amorphous phase, enriched in Si and Ca, was present in the triple grain junctions and along the grain boundaries. 2. The impedance spectra contained two contributions with capacitances in the range typically assigned to grain boundary processes. 3. While oxide ions were identified as the major charge carriers in the grain interiors, protons were found to play an important role as charge carriers in the grain boundary related processes. 4. The behaviour of the two resistivities associated with the grain boundary processes could be explained by applying a defect chemical model in which protons and oxygen vacancies are the major charge carriers. Acknowledgements RH acknowledges funding by the Research Council of Norway, Grant No. 15851/431 (Functional Oxides for Energy Technology, NANOMAT). References [1] T. Ishihara, N.M. Sammes, O. Yamamoto, High Temperature Solid Oxide Fuel Cells, Elsevier, Oxford, 2003, p. 83. [2] H. Iwahara, T. Esaka, H. Uchida, N. Maeda, Solid State Ionics 3–4 (1981) 359. [3] R. Haugsrud, T. Norby, Nature Materials 5 (2006) 193. [4] J.E. Bauerle, J. Phys. Chem. Solids 30 (1969) 2657. [5] S.P.S. Badwal, J. Drennan, Solid State Ionics 40–41 (1990) 869. [6] J. Fleig, J. Maier, J. Am. Ceram. Soc. 82 (1999) 3485. [7] X. Guo, J. Maier, J. Electrochem. Soc. 148 (2001) E121. [8] X. Guo, W. Sigle, J. Fleig, J. Maier, Solid State Ionics 154–155 (2002) 555. [9] X. Guo, R. Waser, Prog. Mater. Sci. 51 (2006) 151. [10] X. Guo, W. Sigle, J. Maier, J. Am. Ceram. Soc. 86 (2003) 77. [11] X. Guo, J. Fleig, J. Maier, J. Electrochem. Soc. 148 (2001) J50. [12] X. Guo, J. Fleig, J. Maier, Solid State Ionics 154–155 (2002) 563. [13] H.G. Bohn, T. Schober, T. Mono, W. Schilling, Solid State Ionics 117 (1999) 219. [14] H.G. Bohn, T. Schober, J. Am. Ceram. Soc. 83 (2000) 768. [15] C. Kjølseth, H. Fjeld, P.I. Dahl, C. Estournès, R. Haugsrud, T. Norby, submitted to Solid State Ionics. [16] R. Haugsrud, T. Norby, Solid State Ionics 177 (2006) 1129. [17] R. Haugsrud, B. Ballesteros, M. Lira-Cantu, T. Norby, J. Electrochem. Soc. 153 (2006) J87. [18] R. Haugsrud, T. Norby, J. Am. Ceram. Soc. 90 (2007) 1116. [19] R. Haugsrud, H. Fjeld, K.R. Haug, T. Norby, J. Electrochem. Soc. 154 (2007) B77. [20] R. Haugsrud, Solid State Ionics 178 (2007) 555. [21] R. Haugsrud, T. Norby, Solid State Electrochemistry, Proceedings of the Risoe International Symposium on Materials Science, 26th, Roskilde, Denmark, Sept. 4–8, 2005, 2005, p. 209. [22] J. Rodríguez-Carvajal, FullProf Ver. Feb-2007, 2007. [23] T. Norby, Solid State Ionics 28–30 (1988) 1586. [24] T. Jacobsen, B. Zachau-Christiansen, L. Bay, S. Skaarup, Proc. 17th Risø International Symposium on “High Temperature Electrochemistry: Ceramics and Metals”, Roskilde, Denmark, 1996, p. 29. [25] S.M. Haile, D.L. West, J. Campbell, J. Mater. Res. 13 (1998) 1576. [26] O. Knop, F. Brisse, L. Castelliz, Sutarno, Canadian Journal of Chemistry 43 (1965) 2812. [27] G.C. Lau, B.D. Muegge, T.M. McQueen, E.L. Duncan, R.J. Cava, J. Solid State Chem. 179 (2006) 3126. [28] J.M. Porras-Vazquez, A.G. De la Torre, E.R. Losilla, M.A.G. Aranda, Solid State Ionics 178 (2007) 1073. [29] K. Amezawa, J.-i. Yamada, N. Kitamura, Y. Tomii, T. Handa, N. Yamamoto, Solid State Ionics 176 (2005) 341. [30] G. Harley, R. Yu, L.C. De Jonghe, Solid State Ionics 178 (2007) 769. [31] J. Jamnik, X. Guo, J. Maier, Appl. Phys. Lett. 82 (2003) 2820. [32] P. Jasinski, V. Petrovsky, T. Suzuki, H.U. Anderson, J. Electrochem. Soc. 152 (2005) J27. [33] W. Lai, S.M. Haile, J. Am. Ceram. Soc. 88 (2005) 2979. [34] J. Jamnik, Solid State Ionics 157 (2003) 19.