- Email: [email protected]
Influence of microstructure on electrical properties in BaZr0.5In0.5O3 − δ proton conductor
Available online at www.sciencedirect.com
Solid State Ionics 179 (2008) 1155 – 1160 www.elsevier.com/locate/ssi
Influence of microstructure on electrical properties in BaZr0.5In0.5O3 − δ proton conductor Istaq Ahmed a,⁎, Sten-G. Eriksson a , Elisabet Ahlberg b , Christopher S. Knee b a
Department of Chemical and Biological Engineering, Chalmers University of Technology, SE-412 96, Göteborg, Sweden b Department of Chemistry, Göteborg University, SE-412 96 Göteborg, Sweden Received 9 July 2007; received in revised form 13 February 2008; accepted 20 February 2008
Abstract Two different preparative routes, one a wet chemical route (WCR) and the other a traditional solid state sintering route (SSR), have been used to synthesise the oxygen deficient perovskite BaZr0.5In0.5O3 − δ. Analysis of X-ray powder diffraction data showed that both phases adopt cubic crystal structures. Environmental scanning electron microscope (ESEM) images showed that a smaller grain size was obtained for the WCR sample. Dynamic thermo-gravimetric analysis indicated significant mass losses for hydrated samples compared to their respective dried samples. Proton conductivity was investigated for hydrated and as-prepared phases under dry and wet atmospheres, respectively. The hydrated sample prepared via the solid state sintering route displayed a higher proton conductivity under dry conditions, reflecting the presence of fewer grain boundaries in this sample. © 2008 Elsevier B.V. All rights reserved. Keywords: X-ray diffraction; ESEM; Proton conductor; Perovskite; Grain boundaries
1. Introduction High proton conductivity and chemical stability are highly desirable characteristics for a material to be used as an electrolyte for intermediate temperature range Solid Oxide Fuel Cells (SOFC). In this respect, zirconates are chemically more stable compared to cerates [1]. Protons accumulated at grain boundaries facilitate the transport process, for example, in Y2O3 [2]. In contrast, proton transports in the grain boundaries of perovskite structured oxides are relatively slower. This adverse effect is believed to be a greater problem for zirconates compared to cerates [3]. In order to improve the grain boundary conductivity, it is necessary to understand how their properties (for example, grain size, orientation, segregation, etc.) influence the proton conductivity. Iguchi et al. [4] have investigated the influence of grain structure on electrical conductivity on BaZr0.95Y0.05O3 and it was reported that proton conductivity was not affected by the grain structure. However, Haile et al. [5] previously reported that the grain boundary of BaCe0.85Gd0.15O3 was strongly influenced
⁎ Corresponding author. E-mail address: [email protected] (I. Ahmed). 0167-2738/$ - see front matter © 2008 Elsevier B.V. All rights reserved. doi:10.1016/j.ssi.2008.02.031
by the grain size. It was shown that the number of grain boundaries was dependent on the average grain size, i.e., samples sintered at higher temperature showed larger grains, fewer grain boundaries and a lower total grain boundary resistance. Indeed, the influence of grain size on electrical properties may vary for different systems. In this work we have chosen to investigate the heavily acceptor doped perovskite oxide, BaZr0.5In0.5O3 − δ. Structural proton positions in this material have already been located at low temperature by Ahmed et al. [6]. The grain structure and electrical conductivity were investigated in terms of two different preparative routes, namely, by a traditional solid state sintering route (SSR) and by a wet chemical route (WCR). The combined techniques of X-ray powder diffraction (XRPD), dynamic thermogravimetric analysis (TGA), environmental scanning electron microscopy (ESEM) and electrochemical measurements (impedance analysis) have been used in this study. 2. Experimental The material BaZr0.5In0.5O3 − δ was prepared by both a wet chemical route (WCR) and a solid state sintering route (SSR). In the case of the WCR, appropriate amounts of Ba(C2H3O2)2,
1156
I. Ahmed et al. / Solid State Ionics 179 (2008) 1155–1160
Zr(C5H8O2)4 and In(C2H3O2)3, were dissolved in hot acetic acid with a stirrer. Then the solution was heated slowly to 150– 250 °C in order to evaporate the acid. After evaporating the acid the solution formed a black solid. This black solid was ground in an agate mortar and heated to 400 °C for 2 h followed by another heating at 800 °C for 8 h. A yellow powder formed as a product which finally was ball milled at 500 rpm for 15 min and then pelletized and sintered at 1100 °C in air for 24 h. Regarding the solid state sintered material it was prepared in the same way as reported in Ref. [6] with a final sintering temperature of 1500 °C. After the final sintering the density of the samples were determined by dimension measurement. Calculated densities were ≈ 60% for WCR samples (both for hydrated and as-prepared), and ≈ 84% for hydrated and ≈ 90% for asprepared SSR samples. Measurement of the density by an immersion technique is accurate only when the samples have densities N 95%, as lower-density samples may allow water to enter the pores [7]. Dried samples were prepared by annealing the as-prepared samples at 950 °C under ≈ 10− 6 mbar pressure overnight. The hydration reaction was carried out at ≈300 °C under humid atmosphere (N2 saturated with water vapour at 76.2 °C corresponding to p(H2O) ≈ 0.40 atm) for several days. The X-ray diffraction measurements were carried out at ambient temperature using a Bruker AXS D8 ADVANCE VARIO powder diffractometer (CuKα1 = 1.54058 Å). Microstructure analysis has been performed in Electro Scan 2020 Environmental Scanning Electron Microscope (ESEM). The microscope was operated at 20 kV for secondary electron imaging. Dynamic TGA was carried out on dried and hydrated samples during heating run. The heating run was performed from ambient temperature to 850 °C with a NETZSCH STA 409 PC
at a rate of 1.5 °C/min in a stream of nominally dry O2 and Ar gas at a flow rate of 20 ml/min. The impedance was measured from 1 MHz to 1 Hz using a Solatron 1260 frequency response analyser in the stand alone mode. The sine wave amplitude was 1 V rms. The conductivity cell used was a ProboStat™ from Norwegian Electro Ceramics AS (NorECs) [8]. Approximately 0.7–0.8 cm2 of the oxide electrode surface was covered with conducting platinum paste and platinum grid to assure good ohmic contacts. The covered surface area was measured by an optical microscope. For the asprepared sample the impedance measurement was performed in porous AL23 tube under nominally dry and wet (saturated at 22 °C with water or heavy water vapour or p(H2O / D2O) = 0.0261 atm) Ar atmosphere on cooling from 1000 to 150 °C in 50 °C steps. After reaching the desired temperature, 1 h elapsed before the impedance spectra was recorded. For hydrated samples impedance measurements were performed from 150 to 900 °C (heating cycle) and down from 900 to 150 °C (cooling cycle) under dried Ar atmosphere with 10 min elapsed before the impedance was recorded. In this case the Ar gas was dried by flowing through two beds of P2O5. In order to limit any water up take from the ambient atmospheres, two silica tubes were used as extra protection during the experiments. 3. Results and discussion Fig. 1 shows the room temperature XRPD patterns of dried and hydrated samples. Indexing of these patterns indicates that all samples possess cubic symmetry of space group Pm-3m. Peak broadening for the WCR sample clearly shows a smaller grain size compared to the SSR sample (see inset a) in Fig. 1). The Scherrer equation [9] was used to calculate the average grain size of the samples. Nano-sized (≈ 46 nm) grains were
Fig. 1. X-ray powder diffraction patterns for BaZr0.5In0.5O3 − δ samples. Insets show the dependence of the width of the 110 Bragg peak on the preparative route a) and hydration level b) and c). Inset d) shows the 111 Bragg peak on hydrated and dried WCR samples, respectively.
I. Ahmed et al. / Solid State Ionics 179 (2008) 1155–1160
Fig. 2. ESEM images of BaZr0.5In0.5O3 −
obtained for the WCR sample, which was approximately five times smaller than the average crystallite size of the SSR sample. Similar findings have been reported previously [4,5,10,11]. The lower sintering temperature for the WCR sample caused the smaller grain size. In order to grow bigger grains a higher surface energy is required. This higher surface energy is provided externally by sintering the samples at higher temperatures. An expansion of the unit cell after hydration was observed, Fig. 1, insets b) and c) for SSR and WCR samples, respectively. The indexed unit cell parameters (a) have been determined to be 4.1922(12) Å, 4.23(0.05) Å, 4.1976(1) Å and 4.2376(4) Å for WCR dried and hydrated and SSR dried and hydrated samples, respectively. This expansion, which is generally common for UU hydrated samples, is a result of filling oxygen vacancies (VO ) by U hydroxyl groups (OHO) in the presence of water vapour according to the following reaction, H2 OðgÞ þ OOx þ V UUO ↔2OHOU
ð1Þ
After hydration both WCR and SSR samples show two sets of diffraction lines. One set occurs close to the 2θ angles obtained for the dried sample which is non-hydrated phase whilst the other set of diffraction peaks occurs at lower 2θ angles due to the hydrated phase. This is commonly observed for heavily doped hydrated samples and indicates that it is not a result of a lowering of symmetry after hydration (see splitting of the 111 reflection on hydrated WCR sample, inset d) in Fig. 1). We have previously observed similar results from both X-ray and neutron powder diffraction analysis of In3+ doped BaZrO3 samples [6,12]. ESEM images (Fig. 2) clearly showed the smaller grain size for the sample prepared by the WCR, which is consistent with XRD results. No secondary phase or specific grain boundary orientation was observed from the images. Significant mass losses occurred for all hydrated samples (Fig. 3) from the TGA experiments. During heating the samples mass loss occurred beyond T ≈ 275 °C. The maximum mass losses for hydrated WCR and SSR samples were ≈ 1.53% and 1.32%, respectively. These results suggest that approximately 97% and 83% of the theoretically possible protonic defects [OHU] were filled during the hydration process. Similar results
δ
1157
samples prepared by a) SSR and b) WCR.
for the SSR sample have been reported under humid atmospheres [12,13]. The TGA results show that approximately full hydration was observed for hydrated WCR sample which contradicts the diffraction results (see insets c) and d) in Fig. 1) where the amount of non-hydrated phase is larger compared to the hydrated phase. It is likely that after hydration of the WCR sample, most of the protons are incorporated in the grain-boundary region. Only a small fraction of the incorporated protons are present in the bulk region corresponding to the ones we detected from the diffraction experiment and most of the incorporated protons during the hydration process in WCR sample were not observed in the diffraction study. In addition, the WCR sample shows more gradual mass loss compared to the SSR sample. This may be a result of “trapped protons” at grain boundaries which require higher energy to leave from the sample. Both XRD and SEM results show that the WCR sample on average has a smaller grain size which indicates that WCR sample have a larger number of grain boundaries compared to the SSR sample. Higher mass loss for the WCR sample may be due to the larger surface area which probably facilitates the hydration reaction (Eq. (1)). The dried SSR sample showed very little mass loss during the de-hydration process. It is likely that some protons remain in the dry sample.
Fig. 3. Comparison of TGA results under different gas atmospheres for BaZr0.5In0.5O3 − δ samples prepared by both SSR and WCR.
1158
I. Ahmed et al. / Solid State Ionics 179 (2008) 1155–1160
Fig. 4. Complex plane plots at 300 °C for different samples together with the equivalent circuit used to extract the conductivity data. Fig. 6. Total conductivity of the BaZr0.5In0.5O3 − δ sample prepared via SSR.
Infra-red (IR) spectroscopy on dried BaZr0.25In0.75O3 − δ sample showed that it is difficult to remove all the protons from the sample [12]. In contrast, the dried WCR sample showed a much greater mass loss during the heating run. This behaviour is discussed together with conductivity results in the text below. Typical impedance spectra are shown in Fig. 4 together with the equivalent circuit used in the data analysis. Some data sets could not be evaluated at higher temperatures N 700 °C and also at lower temperatures (especially for hydrated samples during cooling) b 250 °C. At higher temperatures mainly electrode polarization (capacitance value ≤ 10− 6 F cm− 2) was observed, whilst at lower temperatures high electrode resistance was observed, which limited our ability to extract the conductivity data reliably. The total conductivity for as-prepared samples measured under wet Ar atmosphere (Figs. 5 and 6) showed a higher conductivity compared to the measurements under dry Ar atmosphere and the effect is stronger at lower temperatures. This is typical for proton conducting materials [3,10,13–17]. Conductivity depend both on the concentration (c) of charge carriers and on their mobility (µ). In a wet atmosphere, the protons are incorporated according to Eq. (1), hence, the concentration of
Fig. 5. Total conductivity of the BaZr0.5In0.5O3 − δ sample prepared via WCR.
charge carriers (protons) is higher compared to the concentration in dry atmosphere. It is surprising that the as-prepared samples under dried atmosphere show some proton conducting behaviour. This is probably due to 1) significant number of protons incorporated during sample synthesis time, since the sample was cooled down under ambient atmosphere, 2) water can diffuse from ambient atmosphere through porous Alumina tube during the electrical measurement. The latter case has also been observed by Norby et al. [18]. Total conductivity (Figs. 5 and 6) measured under heavy water (D2O) atmosphere was lower, especially in the intermediate temperature range (150–550 °C), compared to a normal water (H2O) atmosphere. This is due to the isotope effect, which in principle confirms that the material is a proton conductor. This effect is stronger for the WCR sample, particularly in the low temperature range. As stated in Experimental section that as-prepared WCR sample is more porous (ρtheoretical ≈ 60%) than that of SSR sample (ρtheoretical ≈ 90%). Therefore, it is relatively easier to establish equilibrium in WCR compared to SSR sample especially in the low temperature region [10,19]. This may explain the stronger isotope effect in WCR sample compared to SSR sample. The total conductivity of the hydrated samples (Figs. 5 and 6) during the heating cycle shows an increase by more than one order of magnitude in comparison with the respective cooling cycle. At temperatures above 550 °C the conductivity values are similar for all the cycles. The higher value of conductivity obtained in the intermediate temperature range (150– 550 °C) for the heating cycles is due to the presence of a large number of protons within the samples. As mentioned earlier in experimental section, the hydrated samples were annealed at T ≈ 300 °C under wet atmosphere for several days to incorporate protons which will act as charge carriers. With increasing temperature the protons leave the sample and at higher temperatures, T N550 °C, the majority of protons has left the samples and hole conduction dominates. Dynamic TGA (Fig. 4) and temperature dependent infra-red (IR) spectroscopy [20] also indicate that in hydrated samples the protons start to leave the material ≈ 275 °C. Especially beyond 550 °C (Fig. 3) there is a very little mass loss observed
I. Ahmed et al. / Solid State Ionics 179 (2008) 1155–1160
Fig. 7. Comparison of total conductivity of BaZr0.5In0.5O3 − δ samples prepared by WCR and SSR.
1159
sample has a larger grain size compared to the WCR sample, implying that a lower number of grain boundaries are present in the sample. This lower number of grain boundaries resulted in lower grain boundary resistance. Since the total conductivity is the sum of bulk and grain boundary conductivity, a higher value of the total conductivity for the SSR sample is expected. It is notable that in the absence of protons (cooling cycle) the conductivity of both samples is very similar in the whole temperature range. The WCR sample under wet Ar shows a higher conductivity compared to that of the SSR sample, especially below 400 °C. This behaviour is, however, not yet fully understood. We speculate that i) the smaller grain size (which provides a larger surface area) or ii) a higher porosity (ρ ≈ 60% of theoretical) promotes water absorption. In a SSR sample, the grain size is large resulting in a smaller surface area and lower porosity (ρ ≈ 90% of theoretical). Similar results have been reported by Wang et al. [19]. 4. Conclusions
indicating that most of the protons have left the sample. In order to reflect this behaviour, impedance measurements were performed from 900 to 150 °C (cooling cycle) under identical inert atmosphere. The conductivity values were more than one order of magnitude lower than the values obtained during the heating cycle at a temperature below 300 °C (Figs. 5 and 6). Since the experiment was performed in an inert atmosphere, further proton uptake from the atmosphere was limited. Analogous behaviour has been reported previously for many systems [12,21–23]. The activation energies (150–300 °C) for the as-prepared samples (under wet Ar atmosphere) and for the heating cycle of hydrated samples (under dried Ar atmosphere) were in the range 0.54–0. 59 eV. These values are slightly higher than those typically found for the best proton conductors (0.4–0.5 eV) such as Y-doped BaZrO3 or BaCeO3 [24,25]. Islam et al. calculated [9] that the binding energies of hydroxyl-dopant pairs (OHOU M′Zr) were − 0.58 eV and − 0.26 eV, for M = In and Y, respectively. Therefore, slower proton mobility or higher activation energy in In-doped BaZrO3 is expected. Higher activation energy (0.7 eV, Fig. 6) for the as-prepared SSR sample under wet Ar atmosphere was obtained. The reason for this is not yet clear. A plausible explanation is that as the sample was dense (≈ 90% of theoretical) it takes longer time to reach equilibrium during the conductivity measurements. Similar findings have been found for Ba3Ca1.18Nb1.82O9 − δ (BCN 18) [10,19]. Typically, hole or oxide ion conductivities have higher activation energies and tend to dominate at higher temperature [26]. Therefore, the higher activation (0.85–0.97 eV) energy found at higher temperature, T N 550 °C for heating cycles and cooling cycles in the whole temperature range (150–1000 °C) reflects that the conductivity is not due to protons, but rather to hole- or oxide ion conduction. Similar findings have also been reported by Bohn et al. [27] for Y-doped BaZrO3. For comparison, from Fig. 7 it is clear that a hydrated SSR sample shows higher proton conductivity compared to that of a WCR sample during the heating cycle under a dried Ar atmosphere. From XRD and ESEM results we know that the SSR
By using a wet chemical route and a standard solid state synthesis route it was possible to obtain single-phase BaZr0.5In0.5O3 − δ samples. Analysis of XRPD indicated that all the samples were cubic perovskites. The WCR sample showed a smaller grain size compared to the SSR sample. Thermo-gravimetric results indicate that relatively larger number of protonic defects can be filled in the WCR sample upon hydration. Proton conduction in these materials was mainly dominating in the temperature range (150–550 °C). Higher proton conductivity was obtained for pre-hydrated SSR sample containing larger grain size compared to pre-hydrated WCR sample containing smaller grain size. Acknowledgements This work was supported by the Swedish Research Council, and the National Graduate School in Material Science, “Solid state proton conductivity: structure, dynamics and simulation”, at Chalmers University of Technology, Gothenburg, Sweden. References [1] M.J. Scholten, J. Schoonman, J.C. Van Miltenburg, H.A.J. Oonk, Solid State Ion. 61 (1993) 83. [2] T. Norby, P. Kofstad, Solid State Ion. 20 (1986) 169. [3] T. Norby, M. Wideroe, R. Glöckner, Y. Larring, Dalton Trans. 19 (2004) 3012. [4] F. Iguchi, T. Yamada, N. Sata, T. Tsurui, H. Yugami, Solid State Ion. 177 (2006) 2381. [5] S.M. Haile, G. Staneff, K.H. Ryu, J. Mater. Sci. 36 (2001) 1149. [6] I. Ahmed, C.S. Knee, M. Karlsson, S.-G. Eriksson, P.F. Henry, A. Matic, D. Engberg, L. Börjesson, J. Alloys Compd. 450 (2008) 103. [7] I.R. Gibson, G.P. Dransfield, J.T.S. Irvine, J. Mater. Sci. 33 (1998) 4297. [8] www.norecs.com. [9] R. Jenkins, R.L. Snyder. Introduction to X-ray Powder Diffractometry. ISBN:0-471-51339-3. [10] S. Valkenberg, H.G. Bohn, W. Schilling, Solid State Ion. 97 (1997) 511. [11] M.J. Verkerk, B.J. Middelhuis, A.J. Burggraaf, Solid State Ion. 6 (1982) 159. [12] I. Ahmed, S.-G. Eriksson, E. Ahlberg, C.S. Knee, P. Berastegui, L.-G. Johansson, H. Rundlöf, M. Karlsson, A. Matic, L. Börjesson, D. Engberg, Solid State Ion. 177 (2006) 1395.
1160 [13] [14] [15] [16] [17] [18] [19] [20] [21]
I. Ahmed et al. / Solid State Ionics 179 (2008) 1155–1160 A. Manthiram, J.F. Kuo, J.B. Goodenough, Solid State Ion. 62 (1993) 225. A.R. Potter, R.T. Baker, Solid State Ion. 177 (2006) 1917. H.G. Bohn, T. Schober, J. Am. Ceram. Soc. 83 (2000) 768. W. Wang, A.V. Virkar, J. Power Sources 142 (2005) 1. R. Haugsrud, T. Norby, Solid State Ion. 177 (2006) 1129. T. Norby, P. Kofstad, J. Am. Ceram. Soc. 67 (1984) 786. W. Wang, A.V. Virkar, J. Electrochem. Soc. 151 (2004) A1565. M. Karlsson manuscript in preparation. I. Ahmed, S.-G. Eriksson, E. Ahlberg, C.S. Knee, M. Karlsson, A. Matic, L. Börjesson, D. Engberg, Solid State Ion. 177 (2006) 2357.
[22] I. Ahmed, E. Ahlberg, S.-G. Eriksson, C.S. Knee, M. Karlsson, A. Matic, L. Börjesson, Ceram. Eng. Sci. Proc. 27 (2006) 105. [23] I. Ahmed, S.-G. Eriksson, E. Ahlberg, C.S. Knee, H. Götlind, L.-G. Johansson, M. Karlsson, A. Matic, L. Börjesson, Solid State Ion. 178 (2007) 515. [24] K.D. Kreuer, Solid State Ion. 125 (1999) 285. [25] H. Iwahara, T. Yajima, T. Hibino, K. Ozaki, H. Suzuki, Solid State Ion. 61 (1993) 65. [26] N. Bonanos, Solid State Ion. 53-56 (1992) 967. [27] H.G. Bohn, T. Schober, J. Am. Ceram. Soc. 83 (2000) 768.
Solid State Ionics 179 (2008) 1155 – 1160 www.elsevier.com/locate/ssi
Influence of microstructure on electrical properties in BaZr0.5In0.5O3 − δ proton conductor Istaq Ahmed a,⁎, Sten-G. Eriksson a , Elisabet Ahlberg b , Christopher S. Knee b a
Department of Chemical and Biological Engineering, Chalmers University of Technology, SE-412 96, Göteborg, Sweden b Department of Chemistry, Göteborg University, SE-412 96 Göteborg, Sweden Received 9 July 2007; received in revised form 13 February 2008; accepted 20 February 2008
Abstract Two different preparative routes, one a wet chemical route (WCR) and the other a traditional solid state sintering route (SSR), have been used to synthesise the oxygen deficient perovskite BaZr0.5In0.5O3 − δ. Analysis of X-ray powder diffraction data showed that both phases adopt cubic crystal structures. Environmental scanning electron microscope (ESEM) images showed that a smaller grain size was obtained for the WCR sample. Dynamic thermo-gravimetric analysis indicated significant mass losses for hydrated samples compared to their respective dried samples. Proton conductivity was investigated for hydrated and as-prepared phases under dry and wet atmospheres, respectively. The hydrated sample prepared via the solid state sintering route displayed a higher proton conductivity under dry conditions, reflecting the presence of fewer grain boundaries in this sample. © 2008 Elsevier B.V. All rights reserved. Keywords: X-ray diffraction; ESEM; Proton conductor; Perovskite; Grain boundaries
1. Introduction High proton conductivity and chemical stability are highly desirable characteristics for a material to be used as an electrolyte for intermediate temperature range Solid Oxide Fuel Cells (SOFC). In this respect, zirconates are chemically more stable compared to cerates [1]. Protons accumulated at grain boundaries facilitate the transport process, for example, in Y2O3 [2]. In contrast, proton transports in the grain boundaries of perovskite structured oxides are relatively slower. This adverse effect is believed to be a greater problem for zirconates compared to cerates [3]. In order to improve the grain boundary conductivity, it is necessary to understand how their properties (for example, grain size, orientation, segregation, etc.) influence the proton conductivity. Iguchi et al. [4] have investigated the influence of grain structure on electrical conductivity on BaZr0.95Y0.05O3 and it was reported that proton conductivity was not affected by the grain structure. However, Haile et al. [5] previously reported that the grain boundary of BaCe0.85Gd0.15O3 was strongly influenced
⁎ Corresponding author. E-mail address: [email protected] (I. Ahmed). 0167-2738/$ - see front matter © 2008 Elsevier B.V. All rights reserved. doi:10.1016/j.ssi.2008.02.031
by the grain size. It was shown that the number of grain boundaries was dependent on the average grain size, i.e., samples sintered at higher temperature showed larger grains, fewer grain boundaries and a lower total grain boundary resistance. Indeed, the influence of grain size on electrical properties may vary for different systems. In this work we have chosen to investigate the heavily acceptor doped perovskite oxide, BaZr0.5In0.5O3 − δ. Structural proton positions in this material have already been located at low temperature by Ahmed et al. [6]. The grain structure and electrical conductivity were investigated in terms of two different preparative routes, namely, by a traditional solid state sintering route (SSR) and by a wet chemical route (WCR). The combined techniques of X-ray powder diffraction (XRPD), dynamic thermogravimetric analysis (TGA), environmental scanning electron microscopy (ESEM) and electrochemical measurements (impedance analysis) have been used in this study. 2. Experimental The material BaZr0.5In0.5O3 − δ was prepared by both a wet chemical route (WCR) and a solid state sintering route (SSR). In the case of the WCR, appropriate amounts of Ba(C2H3O2)2,
1156
I. Ahmed et al. / Solid State Ionics 179 (2008) 1155–1160
Zr(C5H8O2)4 and In(C2H3O2)3, were dissolved in hot acetic acid with a stirrer. Then the solution was heated slowly to 150– 250 °C in order to evaporate the acid. After evaporating the acid the solution formed a black solid. This black solid was ground in an agate mortar and heated to 400 °C for 2 h followed by another heating at 800 °C for 8 h. A yellow powder formed as a product which finally was ball milled at 500 rpm for 15 min and then pelletized and sintered at 1100 °C in air for 24 h. Regarding the solid state sintered material it was prepared in the same way as reported in Ref. [6] with a final sintering temperature of 1500 °C. After the final sintering the density of the samples were determined by dimension measurement. Calculated densities were ≈ 60% for WCR samples (both for hydrated and as-prepared), and ≈ 84% for hydrated and ≈ 90% for asprepared SSR samples. Measurement of the density by an immersion technique is accurate only when the samples have densities N 95%, as lower-density samples may allow water to enter the pores [7]. Dried samples were prepared by annealing the as-prepared samples at 950 °C under ≈ 10− 6 mbar pressure overnight. The hydration reaction was carried out at ≈300 °C under humid atmosphere (N2 saturated with water vapour at 76.2 °C corresponding to p(H2O) ≈ 0.40 atm) for several days. The X-ray diffraction measurements were carried out at ambient temperature using a Bruker AXS D8 ADVANCE VARIO powder diffractometer (CuKα1 = 1.54058 Å). Microstructure analysis has been performed in Electro Scan 2020 Environmental Scanning Electron Microscope (ESEM). The microscope was operated at 20 kV for secondary electron imaging. Dynamic TGA was carried out on dried and hydrated samples during heating run. The heating run was performed from ambient temperature to 850 °C with a NETZSCH STA 409 PC
at a rate of 1.5 °C/min in a stream of nominally dry O2 and Ar gas at a flow rate of 20 ml/min. The impedance was measured from 1 MHz to 1 Hz using a Solatron 1260 frequency response analyser in the stand alone mode. The sine wave amplitude was 1 V rms. The conductivity cell used was a ProboStat™ from Norwegian Electro Ceramics AS (NorECs) [8]. Approximately 0.7–0.8 cm2 of the oxide electrode surface was covered with conducting platinum paste and platinum grid to assure good ohmic contacts. The covered surface area was measured by an optical microscope. For the asprepared sample the impedance measurement was performed in porous AL23 tube under nominally dry and wet (saturated at 22 °C with water or heavy water vapour or p(H2O / D2O) = 0.0261 atm) Ar atmosphere on cooling from 1000 to 150 °C in 50 °C steps. After reaching the desired temperature, 1 h elapsed before the impedance spectra was recorded. For hydrated samples impedance measurements were performed from 150 to 900 °C (heating cycle) and down from 900 to 150 °C (cooling cycle) under dried Ar atmosphere with 10 min elapsed before the impedance was recorded. In this case the Ar gas was dried by flowing through two beds of P2O5. In order to limit any water up take from the ambient atmospheres, two silica tubes were used as extra protection during the experiments. 3. Results and discussion Fig. 1 shows the room temperature XRPD patterns of dried and hydrated samples. Indexing of these patterns indicates that all samples possess cubic symmetry of space group Pm-3m. Peak broadening for the WCR sample clearly shows a smaller grain size compared to the SSR sample (see inset a) in Fig. 1). The Scherrer equation [9] was used to calculate the average grain size of the samples. Nano-sized (≈ 46 nm) grains were
Fig. 1. X-ray powder diffraction patterns for BaZr0.5In0.5O3 − δ samples. Insets show the dependence of the width of the 110 Bragg peak on the preparative route a) and hydration level b) and c). Inset d) shows the 111 Bragg peak on hydrated and dried WCR samples, respectively.
I. Ahmed et al. / Solid State Ionics 179 (2008) 1155–1160
Fig. 2. ESEM images of BaZr0.5In0.5O3 −
obtained for the WCR sample, which was approximately five times smaller than the average crystallite size of the SSR sample. Similar findings have been reported previously [4,5,10,11]. The lower sintering temperature for the WCR sample caused the smaller grain size. In order to grow bigger grains a higher surface energy is required. This higher surface energy is provided externally by sintering the samples at higher temperatures. An expansion of the unit cell after hydration was observed, Fig. 1, insets b) and c) for SSR and WCR samples, respectively. The indexed unit cell parameters (a) have been determined to be 4.1922(12) Å, 4.23(0.05) Å, 4.1976(1) Å and 4.2376(4) Å for WCR dried and hydrated and SSR dried and hydrated samples, respectively. This expansion, which is generally common for UU hydrated samples, is a result of filling oxygen vacancies (VO ) by U hydroxyl groups (OHO) in the presence of water vapour according to the following reaction, H2 OðgÞ þ OOx þ V UUO ↔2OHOU
ð1Þ
After hydration both WCR and SSR samples show two sets of diffraction lines. One set occurs close to the 2θ angles obtained for the dried sample which is non-hydrated phase whilst the other set of diffraction peaks occurs at lower 2θ angles due to the hydrated phase. This is commonly observed for heavily doped hydrated samples and indicates that it is not a result of a lowering of symmetry after hydration (see splitting of the 111 reflection on hydrated WCR sample, inset d) in Fig. 1). We have previously observed similar results from both X-ray and neutron powder diffraction analysis of In3+ doped BaZrO3 samples [6,12]. ESEM images (Fig. 2) clearly showed the smaller grain size for the sample prepared by the WCR, which is consistent with XRD results. No secondary phase or specific grain boundary orientation was observed from the images. Significant mass losses occurred for all hydrated samples (Fig. 3) from the TGA experiments. During heating the samples mass loss occurred beyond T ≈ 275 °C. The maximum mass losses for hydrated WCR and SSR samples were ≈ 1.53% and 1.32%, respectively. These results suggest that approximately 97% and 83% of the theoretically possible protonic defects [OHU] were filled during the hydration process. Similar results
δ
1157
samples prepared by a) SSR and b) WCR.
for the SSR sample have been reported under humid atmospheres [12,13]. The TGA results show that approximately full hydration was observed for hydrated WCR sample which contradicts the diffraction results (see insets c) and d) in Fig. 1) where the amount of non-hydrated phase is larger compared to the hydrated phase. It is likely that after hydration of the WCR sample, most of the protons are incorporated in the grain-boundary region. Only a small fraction of the incorporated protons are present in the bulk region corresponding to the ones we detected from the diffraction experiment and most of the incorporated protons during the hydration process in WCR sample were not observed in the diffraction study. In addition, the WCR sample shows more gradual mass loss compared to the SSR sample. This may be a result of “trapped protons” at grain boundaries which require higher energy to leave from the sample. Both XRD and SEM results show that the WCR sample on average has a smaller grain size which indicates that WCR sample have a larger number of grain boundaries compared to the SSR sample. Higher mass loss for the WCR sample may be due to the larger surface area which probably facilitates the hydration reaction (Eq. (1)). The dried SSR sample showed very little mass loss during the de-hydration process. It is likely that some protons remain in the dry sample.
Fig. 3. Comparison of TGA results under different gas atmospheres for BaZr0.5In0.5O3 − δ samples prepared by both SSR and WCR.
1158
I. Ahmed et al. / Solid State Ionics 179 (2008) 1155–1160
Fig. 4. Complex plane plots at 300 °C for different samples together with the equivalent circuit used to extract the conductivity data. Fig. 6. Total conductivity of the BaZr0.5In0.5O3 − δ sample prepared via SSR.
Infra-red (IR) spectroscopy on dried BaZr0.25In0.75O3 − δ sample showed that it is difficult to remove all the protons from the sample [12]. In contrast, the dried WCR sample showed a much greater mass loss during the heating run. This behaviour is discussed together with conductivity results in the text below. Typical impedance spectra are shown in Fig. 4 together with the equivalent circuit used in the data analysis. Some data sets could not be evaluated at higher temperatures N 700 °C and also at lower temperatures (especially for hydrated samples during cooling) b 250 °C. At higher temperatures mainly electrode polarization (capacitance value ≤ 10− 6 F cm− 2) was observed, whilst at lower temperatures high electrode resistance was observed, which limited our ability to extract the conductivity data reliably. The total conductivity for as-prepared samples measured under wet Ar atmosphere (Figs. 5 and 6) showed a higher conductivity compared to the measurements under dry Ar atmosphere and the effect is stronger at lower temperatures. This is typical for proton conducting materials [3,10,13–17]. Conductivity depend both on the concentration (c) of charge carriers and on their mobility (µ). In a wet atmosphere, the protons are incorporated according to Eq. (1), hence, the concentration of
Fig. 5. Total conductivity of the BaZr0.5In0.5O3 − δ sample prepared via WCR.
charge carriers (protons) is higher compared to the concentration in dry atmosphere. It is surprising that the as-prepared samples under dried atmosphere show some proton conducting behaviour. This is probably due to 1) significant number of protons incorporated during sample synthesis time, since the sample was cooled down under ambient atmosphere, 2) water can diffuse from ambient atmosphere through porous Alumina tube during the electrical measurement. The latter case has also been observed by Norby et al. [18]. Total conductivity (Figs. 5 and 6) measured under heavy water (D2O) atmosphere was lower, especially in the intermediate temperature range (150–550 °C), compared to a normal water (H2O) atmosphere. This is due to the isotope effect, which in principle confirms that the material is a proton conductor. This effect is stronger for the WCR sample, particularly in the low temperature range. As stated in Experimental section that as-prepared WCR sample is more porous (ρtheoretical ≈ 60%) than that of SSR sample (ρtheoretical ≈ 90%). Therefore, it is relatively easier to establish equilibrium in WCR compared to SSR sample especially in the low temperature region [10,19]. This may explain the stronger isotope effect in WCR sample compared to SSR sample. The total conductivity of the hydrated samples (Figs. 5 and 6) during the heating cycle shows an increase by more than one order of magnitude in comparison with the respective cooling cycle. At temperatures above 550 °C the conductivity values are similar for all the cycles. The higher value of conductivity obtained in the intermediate temperature range (150– 550 °C) for the heating cycles is due to the presence of a large number of protons within the samples. As mentioned earlier in experimental section, the hydrated samples were annealed at T ≈ 300 °C under wet atmosphere for several days to incorporate protons which will act as charge carriers. With increasing temperature the protons leave the sample and at higher temperatures, T N550 °C, the majority of protons has left the samples and hole conduction dominates. Dynamic TGA (Fig. 4) and temperature dependent infra-red (IR) spectroscopy [20] also indicate that in hydrated samples the protons start to leave the material ≈ 275 °C. Especially beyond 550 °C (Fig. 3) there is a very little mass loss observed
I. Ahmed et al. / Solid State Ionics 179 (2008) 1155–1160
Fig. 7. Comparison of total conductivity of BaZr0.5In0.5O3 − δ samples prepared by WCR and SSR.
1159
sample has a larger grain size compared to the WCR sample, implying that a lower number of grain boundaries are present in the sample. This lower number of grain boundaries resulted in lower grain boundary resistance. Since the total conductivity is the sum of bulk and grain boundary conductivity, a higher value of the total conductivity for the SSR sample is expected. It is notable that in the absence of protons (cooling cycle) the conductivity of both samples is very similar in the whole temperature range. The WCR sample under wet Ar shows a higher conductivity compared to that of the SSR sample, especially below 400 °C. This behaviour is, however, not yet fully understood. We speculate that i) the smaller grain size (which provides a larger surface area) or ii) a higher porosity (ρ ≈ 60% of theoretical) promotes water absorption. In a SSR sample, the grain size is large resulting in a smaller surface area and lower porosity (ρ ≈ 90% of theoretical). Similar results have been reported by Wang et al. [19]. 4. Conclusions
indicating that most of the protons have left the sample. In order to reflect this behaviour, impedance measurements were performed from 900 to 150 °C (cooling cycle) under identical inert atmosphere. The conductivity values were more than one order of magnitude lower than the values obtained during the heating cycle at a temperature below 300 °C (Figs. 5 and 6). Since the experiment was performed in an inert atmosphere, further proton uptake from the atmosphere was limited. Analogous behaviour has been reported previously for many systems [12,21–23]. The activation energies (150–300 °C) for the as-prepared samples (under wet Ar atmosphere) and for the heating cycle of hydrated samples (under dried Ar atmosphere) were in the range 0.54–0. 59 eV. These values are slightly higher than those typically found for the best proton conductors (0.4–0.5 eV) such as Y-doped BaZrO3 or BaCeO3 [24,25]. Islam et al. calculated [9] that the binding energies of hydroxyl-dopant pairs (OHOU M′Zr) were − 0.58 eV and − 0.26 eV, for M = In and Y, respectively. Therefore, slower proton mobility or higher activation energy in In-doped BaZrO3 is expected. Higher activation energy (0.7 eV, Fig. 6) for the as-prepared SSR sample under wet Ar atmosphere was obtained. The reason for this is not yet clear. A plausible explanation is that as the sample was dense (≈ 90% of theoretical) it takes longer time to reach equilibrium during the conductivity measurements. Similar findings have been found for Ba3Ca1.18Nb1.82O9 − δ (BCN 18) [10,19]. Typically, hole or oxide ion conductivities have higher activation energies and tend to dominate at higher temperature [26]. Therefore, the higher activation (0.85–0.97 eV) energy found at higher temperature, T N 550 °C for heating cycles and cooling cycles in the whole temperature range (150–1000 °C) reflects that the conductivity is not due to protons, but rather to hole- or oxide ion conduction. Similar findings have also been reported by Bohn et al. [27] for Y-doped BaZrO3. For comparison, from Fig. 7 it is clear that a hydrated SSR sample shows higher proton conductivity compared to that of a WCR sample during the heating cycle under a dried Ar atmosphere. From XRD and ESEM results we know that the SSR
By using a wet chemical route and a standard solid state synthesis route it was possible to obtain single-phase BaZr0.5In0.5O3 − δ samples. Analysis of XRPD indicated that all the samples were cubic perovskites. The WCR sample showed a smaller grain size compared to the SSR sample. Thermo-gravimetric results indicate that relatively larger number of protonic defects can be filled in the WCR sample upon hydration. Proton conduction in these materials was mainly dominating in the temperature range (150–550 °C). Higher proton conductivity was obtained for pre-hydrated SSR sample containing larger grain size compared to pre-hydrated WCR sample containing smaller grain size. Acknowledgements This work was supported by the Swedish Research Council, and the National Graduate School in Material Science, “Solid state proton conductivity: structure, dynamics and simulation”, at Chalmers University of Technology, Gothenburg, Sweden. References [1] M.J. Scholten, J. Schoonman, J.C. Van Miltenburg, H.A.J. Oonk, Solid State Ion. 61 (1993) 83. [2] T. Norby, P. Kofstad, Solid State Ion. 20 (1986) 169. [3] T. Norby, M. Wideroe, R. Glöckner, Y. Larring, Dalton Trans. 19 (2004) 3012. [4] F. Iguchi, T. Yamada, N. Sata, T. Tsurui, H. Yugami, Solid State Ion. 177 (2006) 2381. [5] S.M. Haile, G. Staneff, K.H. Ryu, J. Mater. Sci. 36 (2001) 1149. [6] I. Ahmed, C.S. Knee, M. Karlsson, S.-G. Eriksson, P.F. Henry, A. Matic, D. Engberg, L. Börjesson, J. Alloys Compd. 450 (2008) 103. [7] I.R. Gibson, G.P. Dransfield, J.T.S. Irvine, J. Mater. Sci. 33 (1998) 4297. [8] www.norecs.com. [9] R. Jenkins, R.L. Snyder. Introduction to X-ray Powder Diffractometry. ISBN:0-471-51339-3. [10] S. Valkenberg, H.G. Bohn, W. Schilling, Solid State Ion. 97 (1997) 511. [11] M.J. Verkerk, B.J. Middelhuis, A.J. Burggraaf, Solid State Ion. 6 (1982) 159. [12] I. Ahmed, S.-G. Eriksson, E. Ahlberg, C.S. Knee, P. Berastegui, L.-G. Johansson, H. Rundlöf, M. Karlsson, A. Matic, L. Börjesson, D. Engberg, Solid State Ion. 177 (2006) 1395.
1160 [13] [14] [15] [16] [17] [18] [19] [20] [21]
I. Ahmed et al. / Solid State Ionics 179 (2008) 1155–1160 A. Manthiram, J.F. Kuo, J.B. Goodenough, Solid State Ion. 62 (1993) 225. A.R. Potter, R.T. Baker, Solid State Ion. 177 (2006) 1917. H.G. Bohn, T. Schober, J. Am. Ceram. Soc. 83 (2000) 768. W. Wang, A.V. Virkar, J. Power Sources 142 (2005) 1. R. Haugsrud, T. Norby, Solid State Ion. 177 (2006) 1129. T. Norby, P. Kofstad, J. Am. Ceram. Soc. 67 (1984) 786. W. Wang, A.V. Virkar, J. Electrochem. Soc. 151 (2004) A1565. M. Karlsson manuscript in preparation. I. Ahmed, S.-G. Eriksson, E. Ahlberg, C.S. Knee, M. Karlsson, A. Matic, L. Börjesson, D. Engberg, Solid State Ion. 177 (2006) 2357.
[22] I. Ahmed, E. Ahlberg, S.-G. Eriksson, C.S. Knee, M. Karlsson, A. Matic, L. Börjesson, Ceram. Eng. Sci. Proc. 27 (2006) 105. [23] I. Ahmed, S.-G. Eriksson, E. Ahlberg, C.S. Knee, H. Götlind, L.-G. Johansson, M. Karlsson, A. Matic, L. Börjesson, Solid State Ion. 178 (2007) 515. [24] K.D. Kreuer, Solid State Ion. 125 (1999) 285. [25] H. Iwahara, T. Yajima, T. Hibino, K. Ozaki, H. Suzuki, Solid State Ion. 61 (1993) 65. [26] N. Bonanos, Solid State Ion. 53-56 (1992) 967. [27] H.G. Bohn, T. Schober, J. Am. Ceram. Soc. 83 (2000) 768.