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High-temperature plastic deformation mechanisms of ytterbium-doped barium cerate proton conductor
Solid State Ionics 225 (2012) 286–290
Contents lists available at SciVerse ScienceDirect
Solid State Ionics journal homepage: www.elsevier.com/locate/ssi
High-temperature plastic deformation mechanisms of ytterbium-doped barium cerate proton conductor M. Jiménez-Melendo ⁎ Departamento de Física de la Materia Condensada and Instituto de Ciencia de Materiales (CSIC-Universidad de Sevilla). Aptdo. 1065. 41080 Sevilla, Spain
a r t i c l e
i n f o
Article history: Received 12 September 2011 Received in revised form 13 March 2012 Accepted 18 March 2012 Available online 10 April 2012 Keywords: Barium cerate Ytterbium Proton conductor Microstructure Grain boundary sliding
a b s t r a c t The enhanced proton conductivity exhibited by trivalent cation-doped barium cerate perovskites makes these materials excellent candidates for electrochemical applications, in particular as electrolytes for solid oxide fuel cells. These devices operate at elevated temperatures, where creep and other deformation processes influence the overall efficiency and lifetime. In this work, the high-temperature plastic deformation mechanisms of fine-grained 5 at.% Yb-doped BaCeO3 polycrystals produced by conventional solid-state reaction has been investigated by means of compressive tests at constant load between 1150 and 1250 °C in air. The creep curves show an unusual sigmoidal behavior, followed by extended steady states of deformation. Grain boundary sliding is the main deformation mechanism, characterized by a stress exponent n of 2, as found in other fine-grained superplastic ceramics and metals. © 2012 Elsevier B.V. All rights reserved.
however, protonic defects are created at expenses of oxygen vacancies by the reversible reaction:
1. Introduction Perovskite-structured oxides are attracting considerable interest in the last years because they exhibit a unique combination of proton, oxygen-ion and electronic conductivity at intermediate and high temperatures. These properties are of importance for practical applications in different high-temperature electrochemical fields: hydrogen sensors, steam electrolysis, steam concentrators, and, particularly, as electrolytes in solid state fuel cells (SOFC) [1–7]. These devices can electrochemically convert a variety of hydrocarbons (hydrogen, natural gas, methanol, etc.) into electricity with high efficiency and very low environmental emissions. Simple perovskite oxides of generic formula ABO3, where A 2+ and B 4+ denote A and B site cations, respectively (especially cerates and zirconates), seem to be particularly suitable for these applications because of their high proton conducting performance when doped with trivalent cations (such as Y 3+, Yb 3+, Nd 3+ and Gd 3+) [8–12]. Substitution of B + site cations by trivalent ions M 3+ results in the formation of extrinsic oxygen vacancies in the oxygen sublattice. In BaCeO3, the defect reaction is, written in Kröger–Vink notation: 0
x
::
M2 O3 →2MCe þ 3OO þ V O
ð1Þ
The oxygen vacancies VO·· are responsible for the oxygen ion conduction at elevated temperatures. In the presence of water vapor,
⁎ Tel.: + 34 954 550938; fax: + 34 954 612097. E-mail address: [email protected]. 0167-2738/$ – see front matter © 2012 Elsevier B.V. All rights reserved. doi:10.1016/j.ssi.2012.03.031
::
x
H 2 Oðg Þ þ V O þ OO →2ðOH Þ˙O
ð2Þ
The protonic defect is described as a hydroxyl group because the interstitial proton is closely associated with the neighboring oxygen ion. Although electronic holes are also formed, particularly at high oxygen partial pressures, their contribution to the overall electrical conductivity is not important at the SOFC operating temperatures (below 800 °C) due to the large band gaps of perovskite oxides. SOFCs and other perovskite oxides-based devices operate at temperatures where creep and other deformation processes (damage accumulation, cracking, microstructural evolution, etc.) may influence their overall performance and failure, along with other relevant effects such as chemical degradation of the electrolyte, electrode–electrolyte interdiffusion or electrode detachment. Numerous investigations are concerned with the processing, structure, defect chemistry and conductivity characteristics of doped barium cerates [8–16]. However, only a few studies have focused on their high-temperature mechanical properties [17–19], despite they are important for predicting the longterm operation of the material. In particular, the determination of the deformation mechanisms operating in creep conditions are still lacking. At the service temperature of proton conducting electrolytebased SOFCs (below 800 °C), creep rates are too low to be measured with confidence in the laboratory (the detection limit of the strain measurement system is about 10 − 7 s − 1). In order to determine the creep parameters of the material with acceptable accuracy, mechanical tests have to be carried out at higher temperatures. Once these
M. Jiménez-Melendo / Solid State Ionics 225 (2012) 286–290
287
parameters are measured, creep rates can be easily extrapolated to other conditions using the power law for creep [20]. Moreover, hightemperature plastic deformation is usually controlled by diffusion, which in turn controls other mass transport-assisted processes such as sintering and grain growth. Creep studies can thus render valuable information to devise optimum processing schedules. The aim of the present paper was to investigate the creep response of ytterbium-doped barium cerate in order to determine the plastic deformation mechanisms operating at high temperatures. Creep tests at constant load were performed in air under different experimental conditions of temperature and stress. Ytterbium was chosen as dopant because it has essentially the same ionic radius than cerium ions [21], and it is therefore expected that Yb ions reside almost exclusively in the Ce 4+-site, ensuring the maximum oxygen vacancy content for the ulterior proton uptake (Eqs. (1) and (2)).
size exponent, Q is the activation energy for flow and R is the gas constant.
2. Experimental procedure
Powder diffractograms showed that both the calcined powders and the sintered pellets of 5 at.% Yb-doped barium cerate are formed exclusively by the orthorhombic perovskite phase, with space group
2.1. Material preparation Samples of polycrystalline BaCe0.95Yb0.05O3 − δ (i.e., 5 at.% Yb, where at.% refers to the cation content in Ce 4+-site) were produced via a solid-state reaction route. Appropriate amounts of commercial powders of BaCO3, CeO2 and Yb2O3 (Sigma–Aldrich, purity > 99.0%) were ball-milled in agate media for 1 h and then calcined at 1200 °C for 10 h in air. The resulting powders were reground again using the same procedure as before. Green compacts of 20 mm in diameter were obtained by room-temperature uniaxial and isostatic pressing at 150 and 210 MPa, respectively, which were sintered at 1500 °C for 10 h in air. The heating ramp was stopped at 600 °C for 2 h to eliminate water and possible hydroxides formed during intermediate stages. The final density of the samples, determined using Archimedes' method, was about 90–95% of the theoretical value of barium cerate (6360 kg/m 3). X-ray structure analyses were systematically performed on both the calcined powders and the sintered pellets in order to ensure the presence of the necessary orthorhombic crystalline phase. Powder X-ray diffraction patterns were obtained using a Bruker AXS D8 Advance X-ray diffractometer with Cu Kα radiation equipped with a scintillation detector in θ–2θ Bragg–Bentano configuration (X-ray Laboratory, CITIUS, University of Sevilla, Spain). Collected X-ray spectra were processed by the Le Bail refinement method [22] using the TOPAS 4.2 Bruker AXS software package.
2.3. Microstructural observations The microstructural characterization of as-fabricated and deformed cerates was carried out using scanning electron microscopy (SEM) (Microscopy Service, CITIUS, University of Sevilla, Spain). To reveal grain boundaries, longitudinal sections were cut from the samples and mechanically polished using up to 1-μm grade diamond paste, and then thermally etched at 1200 °C for 2 h in air. The relevant morphological parameters were measured by using a semiautomatic image analyzer. 3. Results and discussion
(a)
(b)
2.2. Mechanical tests Prismatic specimens of 5 × 3 × 3 mm in size were cut from the sintered pellets with a low-speed diamond saw and used for mechanical tests. Compression tests were carried out in air at temperatures T between 1150 and 1250 °C (0.77Tm b T b 0.82Tm, where Tm = 1850 K is the melting temperature of BaCeO3) under constant load at nominal stresses between 40 and 120 MPa. The specimens were sandwiched between SiC pads in order to reduce the friction with the alumina punching rams of the deformation machine. The recorded data, in_ stantaneous sample length l vs time t, were plotted as ε–ε curves on a semi-logarithmic graph, where ε is the true strain (=ln (lo/l), with lo the initial specimen length), and ε_ = dε/dt is the strain rate. The specimens were typically deformed to total strains of 50–60% for subsequent microstructural observations. The data were analyzed using the standard high-temperature power law for steady-state deformation [20]: n −p ε_ ¼ Aσ d expð−Q =RT Þ
ð3Þ
where A is a parameter depending on the deformation mechanism, d is the grain size, n is the stress exponent, p is the grain
(c)
1 μm Fig. 1. SEM micrographs at the same magnification of 5 at.% Yb-doped BaCeO3: (a) polished and thermally-etched cross-section of the as-received material; (b) fracture surface of the as-received material; (c) fracture surface of a deformed sample. Ultrafine grains are present in the as-received material (some of them marked by arrows in (b)), that grew during deformation.
M. Jiménez-Melendo / Solid State Ionics 225 (2012) 286–290
Pmcn. The crystallographic unit cell volume V5Yb = 339.79 Å 3 is practically identical to that reported in the literature for pure barium cerate Vpure = 339.86 Å 3 [14] and 340.19 Å 3 [11]. The effective ionic radius of Yb 3+ (R = 0.868 Å) is almost identical to that of Ce 4+ (R = 0.87 Å); these cation radii are, in turn, smaller than that of Ba 2+ (R = 1.35 Å) (radii are taken from Shannon [21]). Based on these values, the constancy of the unit cell volume with doping suggests that Yb 3+ has fully occupied the Ce 4+-site; a cell volume contraction would be expected if the Ba 2+-site were occupied by Yb 3+. Atomistic simulations in doped barium cerate [13] have shown that partitioning is unfavorable for small dopant ions, in agreement with the present results. Fig. 1a shows a SEM micrograph of a polished and thermally-etched cross-section of the as-received material. The grains exhibit an equiaxed shape, with an average grain size of 0.4 μm; the grain size distribution is consistent with a log normal law, as shown elsewhere [19]. A similar microstructure was reported for 5 at.% Y-doped BaCeO3 processed by the same route [18]. A detailed observation of the microstructure of the compounds on fresh fracture surfaces shows the presence of numerous ultrafine grains distributed over the boundaries and triple junctions of the grains (marked by arrows in Fig. 1b). These fine grains cannot be observed on conventional finished surfaces (as in Fig. 1a) because barium cerate hydrates easily during mechanical polishing, blurring the grain boundaries. Regarding the mechanical behavior of 5 at.% Yb-doped barium ce_ curves (usually referred as rate, Figs. 2 and 3 show representative ε–ε creep curves) obtained at 1185 and 1235 °C under different loads. The curves exhibit an unusual behavior when compared with other finegrained ceramics. Regardless of experimental conditions, there is systematically an increase in strain rate with strain at the beginning of the deformation until a maximum is reached (sigmoidal creep regime), after which ε_ decreases with a rather constant slope. This sec_ curves corresponds to a steady-state deformation ond stage of the ε–ε regime (secondary creep) under conditions of constant load due to the continuous increase in sample cross-section (and subsequent stress reduction) with strain. Despite the large strains attained, no appreciable barreling and/or cracking was observed in the samples after deformation, in agreement with the shape of the creep curves. The
10-3
STRAIN RATE (s-1)
5 at% Yb-BaCeO3 T = 1185ºC 10-4
.
α =-2.5
.
α =-2.4
εo εo
σ =85MPa
10-5
σ =52MPa
10-6
0
10
20
30
40
50
STRAIN (%) Fig. 2. Creep curves for 5 at.% Yb-doped BaCeO3 deformed at 1185 °C and two different loads. The slopes α of the semi-logarithmic plots of the strain rate ε_ with strain ε under steady-state deformation (dashed lines) are shown. The comparison of the strain rates extrapolated to ε = 0 (initial microstructure), ε_ o , yields a stress exponent n = 2.0.
10-2
5 at% Yb-BaCeO3 T = 1235ºC STRAIN RATE (s-1)
288
.
10-3
εo
. ε
o
10-4
α =-2.2 σ =116MPa α =-2.3
n = 1.9 n = 2.2 10-5
σ = 59 MPa 0
10
20
70 MPa 59MPa 30
40
50
60
STRAIN (%) Fig. 3. Same as in Fig. 2 at T = 1235 °C. The comparison of ε_ o yields a value of n = 2.2. Two determinations of n by load changes between 59 and 70 MPa are also shown.
final cross-sectional area of the specimens measured after testing compared well with the value estimated from the curves, indicating that the specimens underwent a rather homogeneous deformation. These mechanical characteristics correlate well with the behavior reported for this compound [19] and 5 at.% Y-doped BaCeO3 [18] deformed under conditions of constant strain rate, as depicted in Fig. 4. It can be seen that the true stress σ vs true strain ε curves exhibit systematically a peak stress at the beginning of the deformation, followed by a controlled stress drop and then by an extensive secondary creep regime. The stress drop in these curves corresponds to a softening process, and the minimum stress is then equivalent to the _ maximum strain rate in the ε–ε plots. Such mechanical behavior is rather unusual in fine-grained ceramics. However, it has been widely reported in metals at hot deformation conditions [23,24]. In these materials, the peak stress, the softening stage and the subsequent steady-state flow resulted from a balance between dislocation work-hardening, dynamic recovery and dynamic recrystallization, which lead to a refinement of the average grain size. The origin of the yield drop in the cerates is discussed elsewhere [18], and was ascribed to the coarsening of the ultrafine grains after reaching the critical strain required to trigger grain growth; these new grains are able to participate then in the deformation process, resulting in an effective reduction of the average grain size. In the present creep tests, it is possible that the same mechanism is responsible for the initial increase in strain rate with strain according to Eq. (3). Once the final grain size distribution is established, a steady state of deformation is achieved, corresponding to the monotonic decrease of strain rate with strain. The coarsening of the ultrafine grains during deformation is illustrated in Fig. 1c, which shows the fracture surface of a crept sample. _ plots depicted in Figs. 2 and 3 allow the determination of The ε–ε the stress exponent n (Eq. (3)) in the steady-state regime, which gives fundamental information about the atomistic deformation mechanisms operating in the material [20]. Three independent methods were used to evaluate n: (i) by comparing the strain rates measured on different specimens deformed isothermally at different loads (conventional method); (ii) by load changes performed during a test on the same sample (differential method); and (iii) from the _ slopes of the ε–ε plots. In all the three methods, steady-state conditions have to be fulfilled in order to get meaningful values of the stress exponent.
M. Jiménez-Melendo / Solid State Ionics 225 (2012) 286–290
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microstructural modifications (except the coarsening of the ultrafine grains, Fig. 1c) observed in the strained samples with respect to undeformed ones. When plastic deformation is achieved by grain boundary sliding, diffusion is usually the rate-controlling mechanism [26]. The creep activation energy (Eq. (3)) can be thus identified with the diffusion energy of the slowest moving species in the compound along the fastest path. A value of Q = 420 kJ/mol has been determined from temperature changes in this work; it is somewhat higher than the value of Q = 350 kJ/mol reported from constant strain rates experiments [19]. Assuming that grain boundary sliding operates at lower temperatures, a deformation rate of 10 − 10 s − 1 can be estimated at 800 °C and 50 MPa from Eq. (3) with n = 2 and Q = 420 kJ/mol; such a value results in a relative length variation Δl/l of about 10 − 5 per day. Self-diffusion data are very limited for this system. Only the activation energy for oxygen bulk diffusion Q = 60 kJ/mol has been reported [15], which is much smaller than the value for creep. Although a definitive identification of the rate-controlling species cannot be drawn, it is most likely that cation diffusion would be the rate-limiting process, as suggested in other oxide perovskites where diffusional creep has been reported [17,28,29].
TRUE STRESS (MPa)
T = 1250 ºC 150
5 at% Yb - 2x10 -4 s-1
100
5 at% Y - 2x10 -4 s-1 5 at% Y - 8x10 -5 s-1
50
0
0
10
20
30
40
50
TRUE STRAIN (%) Fig. 4. True stress against true strain curves for fine-grained 5 at.% Yb- and 5 at.% Y-doped BaCeO3 at 1250 °C in air obtained at constant cross-head speed; the initial strain rates are shown.
For the first listed method, the steady-state stages were extrapolated (dashed lines in Figs. 2 and 3) to ε = 0, i.e., to initial microstructure [25], in order to avoid the transitory creep regimes at the beginning of the tests. A comparison of the corresponding initial strain rates (denoted by ε_ o ) yields values of n = 2.0 (Fig. 2) and 2.2 (Fig. 3). A value of 2 has been systematically reported in finegrained superplastic ceramics and metals [26,27], where deformation is achieved by grain boundary sliding. It is also predicted by most theoretical models based on grain boundary sliding to explain superplasticity [26]. In this mechanism, the macroscopic deformation of the specimen is accommodated by the sliding of the grains over each other without appreciable deformation of the grains themselves, allowing very large strains (even several hundred per cent) to be achieved. The stress exponent has been also assessed by means of sudden load changes during the steady-state deformation of a sample. Fig. 3 shows two determinations of n by this technique, yielding again values of nearly 2. It is important to note that the maintenance of the strain rate levels after the positive and negative load changes indicates that no microstructural evolution took place during steadystate deformation. Finally, the stress exponent was deduced from the individual _ slopes α of the ε–ε plots. Assuming an ideal homogeneous deformation of the specimen, the instantaneous sample volume V = S.l (where S is the instantaneous cross-section) is equal to the initial volume Vo = So · lo. If Q is the constant applied load, the true stress can be expressed as σ = Q/S = Q · l/So · lo = σo exp(−ε) where the definition of true strain has been used. Then, the steady-state deformation rate ε_ is written by using Eq. (3): n n _ ε∝σ ¼ σ o expð−n:ε Þ
289
ð4Þ
_ ε plot is This equation demonstrates that the slope α of the ε– equal to −n. As shown in Figs. 2 and 3, values of n of about 2 were also measured by this method. It can be safely concluded, therefore, that the creep behavior of fine-grained Yb-doped barium cerate is characterized by a stress exponent n = 2, indicating that grain boundary sliding is the main deformation mechanism, as found in other superplastic materials [26,27]. This conclusion is in agreement with the absence of
4. Conclusions Polycrystals of 5 at.% Yb-doped barium cerate BaCe0.95Yb0.05O3 − δ have been produced by solid state reaction with a relative bulk density of 90–95%. A single orthorhombic perovskite phase was obtained after sintering at 1500 °C for 10 h. The microstructure is very homogeneous and equiaxed, with grains of average size of 0.4 μm. The orthorhombic unit cell volume is barely perturbed by doping, suggesting that dopant cations have exclusively occupied the Ce 4+site. Creep tests at constant load have been carried out in air at temperatures between 1150 and 1250 °C. The creep curves exhibit an initial sigmoidal regime with a maximum in strain rate, followed by an extensive steady state of deformation. A stress exponent n close to 2 has been estimated from different methods, indicating that grain boundary sliding is the main deformation mechanism; such a mechanism is responsible for the superplastic behavior of fine-grained ceramics and metals. Acknowledgment This work was supported by the Project no. MAT2009-13979-C03-01, Ministerio de Ciencia e Innovación, Spain. References [1] K.-D. Kreuer, Chem. Mater. 8 (1996) 610–641. [2] H. Iwahara, Solid State Ionics 86–88 (1996) 9–15. [3] S.C. Singhal, K. Kendall, High Temperature Solid Oxide Fuel Cell. Fundamentals, Design and Applications, Elsevier, Oxford, 2003. [4] N. Ito, in: T. Ishihara (Ed.), Perovskite Oxide for Solid Oxide Fuel Cells, Springer, 2009, pp. 65–94. [5] W.G. Coors, Solid State Ionics 178 (2007) 481–485. [6] J. Guan, S.E. Dorris, U. Balachandran, M. Liu, Solid State Ionics 100 (1997) 45–52. [7] C.H. Chen, H.J.M. Bouwmeester, R.H.E. van Doorn, H. Kruidhof, A.J. Burggraaf, Solid State Ionics 98 (1997) 7–13. [8] S. Imashuku, T. Uda, Y. Nose, K. Kishida, S. Harada, H. Inui, Y. Awakura, J. Electrochem. Soc. 155 (2008) B581–B586. [9] C. Zhang, H. Zhao, S. Zhai, Int. J. Hydrogen Energy 36 (2011) 3649–3657. [10] E. Gorbova, V. Maragou, D. Medvedev, A. Demin, P. Tsiakaras, J. Power Sources 181 (2008) 207–213. [11] N. Osman, I.A. Talib, H.A. Hamid, A.M. Jani, Ionics 14 (2008) 407–413. [12] J. Wu, L.P. Li, W.T.P. Espinosa, S.M. Haile, J. Mater. Res. 19 (2004) 2366–2376. [13] J. Wu, R.A. Davies, M.S. Islam, S.M. Haile, Chem. Mater. 17 (2005) 846–851. [14] K.S. Knight, N. Bonanos, Mater. Res. Bull. 30 (1995) 347–356. [15] M. Amsif, D. Marrero-López, A. Magrasó, J. Pena-Martínez, J.C. Ruiz-Morales, P. Núñez, J. Eur. Ceram. Soc. 29 (2009) 131–138. [16] J. Wu, S.M. Webb, S. Brennan, S.M. Haile, J. Appl. Phys. 97 (1–7) (2005) 054101. [17] E.T. Park, K.C. Goretta, A.R. de Arellano-López, J. Guan, U. Balachandran, S.E. Dorris, Solid State Ionics 117 (1999) 323–330.
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[18] C. Vaquero-Aguilar, M.J. López-Robledo, J. Martínez-Fernández, C. Real, M. Jiménez-Melendo, J. Eur. Ceram. Soc. 31 (2011) 1333–1338. [19] C. Vaquero-Aguilar, M. Jiménez-Melendo, J. Eur. Ceram. Soc. 31 (2011) 2671–2676. [20] J.P. Poirier, Creep of Crystals, Cambridge University Press, Cambridge, 1985. [21] R.D. Shannon, Acta Crystallogr. A32 (1976) 751–767. [22] A. Le Bail, H. Duroy, J.L. Fourquet, Mater. Res. Bull. 23 (1988) 447–452. [23] T. Sakai, J.J. Jonas, Acta Metall. 32 (1984) 189–209. [24] H.J. McQueen, N.D. Ryan, Mater. Sci. Eng., A 322 (2002) 43–63.
[25] A. Bravo-León, J. Ye, M. Jiménez Melendo, A. Domínguez Rodríguez, Scr. Metall. Mater. 29 (1993) 1639–1643. [26] T.G. Nieh, J. Wadsworth, O.D. Sherby, Superplasticity in Metals and Ceramics, Cambridge University Press, Cambridge, 1997. [27] M. Jiménez-Melendo, A. Domínguez-Rodríguez, A. Bravo-León, J. Am. Ceram. Soc. 81 (2008) 2761–2776. [28] J. Wolfenstine, T.R. Armstrong, W.J. Weber, M.A. Boling-Risser, K.C. Goretta, J.L. Routbort, J. Mater. Res. 11 (1996) 657–662. [29] G. Majkic, L.T. Wheeler, K. Salama, Solid State Ionics 146 (2002) 393–404.
Contents lists available at SciVerse ScienceDirect
Solid State Ionics journal homepage: www.elsevier.com/locate/ssi
High-temperature plastic deformation mechanisms of ytterbium-doped barium cerate proton conductor M. Jiménez-Melendo ⁎ Departamento de Física de la Materia Condensada and Instituto de Ciencia de Materiales (CSIC-Universidad de Sevilla). Aptdo. 1065. 41080 Sevilla, Spain
a r t i c l e
i n f o
Article history: Received 12 September 2011 Received in revised form 13 March 2012 Accepted 18 March 2012 Available online 10 April 2012 Keywords: Barium cerate Ytterbium Proton conductor Microstructure Grain boundary sliding
a b s t r a c t The enhanced proton conductivity exhibited by trivalent cation-doped barium cerate perovskites makes these materials excellent candidates for electrochemical applications, in particular as electrolytes for solid oxide fuel cells. These devices operate at elevated temperatures, where creep and other deformation processes influence the overall efficiency and lifetime. In this work, the high-temperature plastic deformation mechanisms of fine-grained 5 at.% Yb-doped BaCeO3 polycrystals produced by conventional solid-state reaction has been investigated by means of compressive tests at constant load between 1150 and 1250 °C in air. The creep curves show an unusual sigmoidal behavior, followed by extended steady states of deformation. Grain boundary sliding is the main deformation mechanism, characterized by a stress exponent n of 2, as found in other fine-grained superplastic ceramics and metals. © 2012 Elsevier B.V. All rights reserved.
however, protonic defects are created at expenses of oxygen vacancies by the reversible reaction:
1. Introduction Perovskite-structured oxides are attracting considerable interest in the last years because they exhibit a unique combination of proton, oxygen-ion and electronic conductivity at intermediate and high temperatures. These properties are of importance for practical applications in different high-temperature electrochemical fields: hydrogen sensors, steam electrolysis, steam concentrators, and, particularly, as electrolytes in solid state fuel cells (SOFC) [1–7]. These devices can electrochemically convert a variety of hydrocarbons (hydrogen, natural gas, methanol, etc.) into electricity with high efficiency and very low environmental emissions. Simple perovskite oxides of generic formula ABO3, where A 2+ and B 4+ denote A and B site cations, respectively (especially cerates and zirconates), seem to be particularly suitable for these applications because of their high proton conducting performance when doped with trivalent cations (such as Y 3+, Yb 3+, Nd 3+ and Gd 3+) [8–12]. Substitution of B + site cations by trivalent ions M 3+ results in the formation of extrinsic oxygen vacancies in the oxygen sublattice. In BaCeO3, the defect reaction is, written in Kröger–Vink notation: 0
x
::
M2 O3 →2MCe þ 3OO þ V O
ð1Þ
The oxygen vacancies VO·· are responsible for the oxygen ion conduction at elevated temperatures. In the presence of water vapor,
⁎ Tel.: + 34 954 550938; fax: + 34 954 612097. E-mail address: [email protected]. 0167-2738/$ – see front matter © 2012 Elsevier B.V. All rights reserved. doi:10.1016/j.ssi.2012.03.031
::
x
H 2 Oðg Þ þ V O þ OO →2ðOH Þ˙O
ð2Þ
The protonic defect is described as a hydroxyl group because the interstitial proton is closely associated with the neighboring oxygen ion. Although electronic holes are also formed, particularly at high oxygen partial pressures, their contribution to the overall electrical conductivity is not important at the SOFC operating temperatures (below 800 °C) due to the large band gaps of perovskite oxides. SOFCs and other perovskite oxides-based devices operate at temperatures where creep and other deformation processes (damage accumulation, cracking, microstructural evolution, etc.) may influence their overall performance and failure, along with other relevant effects such as chemical degradation of the electrolyte, electrode–electrolyte interdiffusion or electrode detachment. Numerous investigations are concerned with the processing, structure, defect chemistry and conductivity characteristics of doped barium cerates [8–16]. However, only a few studies have focused on their high-temperature mechanical properties [17–19], despite they are important for predicting the longterm operation of the material. In particular, the determination of the deformation mechanisms operating in creep conditions are still lacking. At the service temperature of proton conducting electrolytebased SOFCs (below 800 °C), creep rates are too low to be measured with confidence in the laboratory (the detection limit of the strain measurement system is about 10 − 7 s − 1). In order to determine the creep parameters of the material with acceptable accuracy, mechanical tests have to be carried out at higher temperatures. Once these
M. Jiménez-Melendo / Solid State Ionics 225 (2012) 286–290
287
parameters are measured, creep rates can be easily extrapolated to other conditions using the power law for creep [20]. Moreover, hightemperature plastic deformation is usually controlled by diffusion, which in turn controls other mass transport-assisted processes such as sintering and grain growth. Creep studies can thus render valuable information to devise optimum processing schedules. The aim of the present paper was to investigate the creep response of ytterbium-doped barium cerate in order to determine the plastic deformation mechanisms operating at high temperatures. Creep tests at constant load were performed in air under different experimental conditions of temperature and stress. Ytterbium was chosen as dopant because it has essentially the same ionic radius than cerium ions [21], and it is therefore expected that Yb ions reside almost exclusively in the Ce 4+-site, ensuring the maximum oxygen vacancy content for the ulterior proton uptake (Eqs. (1) and (2)).
size exponent, Q is the activation energy for flow and R is the gas constant.
2. Experimental procedure
Powder diffractograms showed that both the calcined powders and the sintered pellets of 5 at.% Yb-doped barium cerate are formed exclusively by the orthorhombic perovskite phase, with space group
2.1. Material preparation Samples of polycrystalline BaCe0.95Yb0.05O3 − δ (i.e., 5 at.% Yb, where at.% refers to the cation content in Ce 4+-site) were produced via a solid-state reaction route. Appropriate amounts of commercial powders of BaCO3, CeO2 and Yb2O3 (Sigma–Aldrich, purity > 99.0%) were ball-milled in agate media for 1 h and then calcined at 1200 °C for 10 h in air. The resulting powders were reground again using the same procedure as before. Green compacts of 20 mm in diameter were obtained by room-temperature uniaxial and isostatic pressing at 150 and 210 MPa, respectively, which were sintered at 1500 °C for 10 h in air. The heating ramp was stopped at 600 °C for 2 h to eliminate water and possible hydroxides formed during intermediate stages. The final density of the samples, determined using Archimedes' method, was about 90–95% of the theoretical value of barium cerate (6360 kg/m 3). X-ray structure analyses were systematically performed on both the calcined powders and the sintered pellets in order to ensure the presence of the necessary orthorhombic crystalline phase. Powder X-ray diffraction patterns were obtained using a Bruker AXS D8 Advance X-ray diffractometer with Cu Kα radiation equipped with a scintillation detector in θ–2θ Bragg–Bentano configuration (X-ray Laboratory, CITIUS, University of Sevilla, Spain). Collected X-ray spectra were processed by the Le Bail refinement method [22] using the TOPAS 4.2 Bruker AXS software package.
2.3. Microstructural observations The microstructural characterization of as-fabricated and deformed cerates was carried out using scanning electron microscopy (SEM) (Microscopy Service, CITIUS, University of Sevilla, Spain). To reveal grain boundaries, longitudinal sections were cut from the samples and mechanically polished using up to 1-μm grade diamond paste, and then thermally etched at 1200 °C for 2 h in air. The relevant morphological parameters were measured by using a semiautomatic image analyzer. 3. Results and discussion
(a)
(b)
2.2. Mechanical tests Prismatic specimens of 5 × 3 × 3 mm in size were cut from the sintered pellets with a low-speed diamond saw and used for mechanical tests. Compression tests were carried out in air at temperatures T between 1150 and 1250 °C (0.77Tm b T b 0.82Tm, where Tm = 1850 K is the melting temperature of BaCeO3) under constant load at nominal stresses between 40 and 120 MPa. The specimens were sandwiched between SiC pads in order to reduce the friction with the alumina punching rams of the deformation machine. The recorded data, in_ stantaneous sample length l vs time t, were plotted as ε–ε curves on a semi-logarithmic graph, where ε is the true strain (=ln (lo/l), with lo the initial specimen length), and ε_ = dε/dt is the strain rate. The specimens were typically deformed to total strains of 50–60% for subsequent microstructural observations. The data were analyzed using the standard high-temperature power law for steady-state deformation [20]: n −p ε_ ¼ Aσ d expð−Q =RT Þ
ð3Þ
where A is a parameter depending on the deformation mechanism, d is the grain size, n is the stress exponent, p is the grain
(c)
1 μm Fig. 1. SEM micrographs at the same magnification of 5 at.% Yb-doped BaCeO3: (a) polished and thermally-etched cross-section of the as-received material; (b) fracture surface of the as-received material; (c) fracture surface of a deformed sample. Ultrafine grains are present in the as-received material (some of them marked by arrows in (b)), that grew during deformation.
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Pmcn. The crystallographic unit cell volume V5Yb = 339.79 Å 3 is practically identical to that reported in the literature for pure barium cerate Vpure = 339.86 Å 3 [14] and 340.19 Å 3 [11]. The effective ionic radius of Yb 3+ (R = 0.868 Å) is almost identical to that of Ce 4+ (R = 0.87 Å); these cation radii are, in turn, smaller than that of Ba 2+ (R = 1.35 Å) (radii are taken from Shannon [21]). Based on these values, the constancy of the unit cell volume with doping suggests that Yb 3+ has fully occupied the Ce 4+-site; a cell volume contraction would be expected if the Ba 2+-site were occupied by Yb 3+. Atomistic simulations in doped barium cerate [13] have shown that partitioning is unfavorable for small dopant ions, in agreement with the present results. Fig. 1a shows a SEM micrograph of a polished and thermally-etched cross-section of the as-received material. The grains exhibit an equiaxed shape, with an average grain size of 0.4 μm; the grain size distribution is consistent with a log normal law, as shown elsewhere [19]. A similar microstructure was reported for 5 at.% Y-doped BaCeO3 processed by the same route [18]. A detailed observation of the microstructure of the compounds on fresh fracture surfaces shows the presence of numerous ultrafine grains distributed over the boundaries and triple junctions of the grains (marked by arrows in Fig. 1b). These fine grains cannot be observed on conventional finished surfaces (as in Fig. 1a) because barium cerate hydrates easily during mechanical polishing, blurring the grain boundaries. Regarding the mechanical behavior of 5 at.% Yb-doped barium ce_ curves (usually referred as rate, Figs. 2 and 3 show representative ε–ε creep curves) obtained at 1185 and 1235 °C under different loads. The curves exhibit an unusual behavior when compared with other finegrained ceramics. Regardless of experimental conditions, there is systematically an increase in strain rate with strain at the beginning of the deformation until a maximum is reached (sigmoidal creep regime), after which ε_ decreases with a rather constant slope. This sec_ curves corresponds to a steady-state deformation ond stage of the ε–ε regime (secondary creep) under conditions of constant load due to the continuous increase in sample cross-section (and subsequent stress reduction) with strain. Despite the large strains attained, no appreciable barreling and/or cracking was observed in the samples after deformation, in agreement with the shape of the creep curves. The
10-3
STRAIN RATE (s-1)
5 at% Yb-BaCeO3 T = 1185ºC 10-4
.
α =-2.5
.
α =-2.4
εo εo
σ =85MPa
10-5
σ =52MPa
10-6
0
10
20
30
40
50
STRAIN (%) Fig. 2. Creep curves for 5 at.% Yb-doped BaCeO3 deformed at 1185 °C and two different loads. The slopes α of the semi-logarithmic plots of the strain rate ε_ with strain ε under steady-state deformation (dashed lines) are shown. The comparison of the strain rates extrapolated to ε = 0 (initial microstructure), ε_ o , yields a stress exponent n = 2.0.
10-2
5 at% Yb-BaCeO3 T = 1235ºC STRAIN RATE (s-1)
288
.
10-3
εo
. ε
o
10-4
α =-2.2 σ =116MPa α =-2.3
n = 1.9 n = 2.2 10-5
σ = 59 MPa 0
10
20
70 MPa 59MPa 30
40
50
60
STRAIN (%) Fig. 3. Same as in Fig. 2 at T = 1235 °C. The comparison of ε_ o yields a value of n = 2.2. Two determinations of n by load changes between 59 and 70 MPa are also shown.
final cross-sectional area of the specimens measured after testing compared well with the value estimated from the curves, indicating that the specimens underwent a rather homogeneous deformation. These mechanical characteristics correlate well with the behavior reported for this compound [19] and 5 at.% Y-doped BaCeO3 [18] deformed under conditions of constant strain rate, as depicted in Fig. 4. It can be seen that the true stress σ vs true strain ε curves exhibit systematically a peak stress at the beginning of the deformation, followed by a controlled stress drop and then by an extensive secondary creep regime. The stress drop in these curves corresponds to a softening process, and the minimum stress is then equivalent to the _ maximum strain rate in the ε–ε plots. Such mechanical behavior is rather unusual in fine-grained ceramics. However, it has been widely reported in metals at hot deformation conditions [23,24]. In these materials, the peak stress, the softening stage and the subsequent steady-state flow resulted from a balance between dislocation work-hardening, dynamic recovery and dynamic recrystallization, which lead to a refinement of the average grain size. The origin of the yield drop in the cerates is discussed elsewhere [18], and was ascribed to the coarsening of the ultrafine grains after reaching the critical strain required to trigger grain growth; these new grains are able to participate then in the deformation process, resulting in an effective reduction of the average grain size. In the present creep tests, it is possible that the same mechanism is responsible for the initial increase in strain rate with strain according to Eq. (3). Once the final grain size distribution is established, a steady state of deformation is achieved, corresponding to the monotonic decrease of strain rate with strain. The coarsening of the ultrafine grains during deformation is illustrated in Fig. 1c, which shows the fracture surface of a crept sample. _ plots depicted in Figs. 2 and 3 allow the determination of The ε–ε the stress exponent n (Eq. (3)) in the steady-state regime, which gives fundamental information about the atomistic deformation mechanisms operating in the material [20]. Three independent methods were used to evaluate n: (i) by comparing the strain rates measured on different specimens deformed isothermally at different loads (conventional method); (ii) by load changes performed during a test on the same sample (differential method); and (iii) from the _ slopes of the ε–ε plots. In all the three methods, steady-state conditions have to be fulfilled in order to get meaningful values of the stress exponent.
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200
microstructural modifications (except the coarsening of the ultrafine grains, Fig. 1c) observed in the strained samples with respect to undeformed ones. When plastic deformation is achieved by grain boundary sliding, diffusion is usually the rate-controlling mechanism [26]. The creep activation energy (Eq. (3)) can be thus identified with the diffusion energy of the slowest moving species in the compound along the fastest path. A value of Q = 420 kJ/mol has been determined from temperature changes in this work; it is somewhat higher than the value of Q = 350 kJ/mol reported from constant strain rates experiments [19]. Assuming that grain boundary sliding operates at lower temperatures, a deformation rate of 10 − 10 s − 1 can be estimated at 800 °C and 50 MPa from Eq. (3) with n = 2 and Q = 420 kJ/mol; such a value results in a relative length variation Δl/l of about 10 − 5 per day. Self-diffusion data are very limited for this system. Only the activation energy for oxygen bulk diffusion Q = 60 kJ/mol has been reported [15], which is much smaller than the value for creep. Although a definitive identification of the rate-controlling species cannot be drawn, it is most likely that cation diffusion would be the rate-limiting process, as suggested in other oxide perovskites where diffusional creep has been reported [17,28,29].
TRUE STRESS (MPa)
T = 1250 ºC 150
5 at% Yb - 2x10 -4 s-1
100
5 at% Y - 2x10 -4 s-1 5 at% Y - 8x10 -5 s-1
50
0
0
10
20
30
40
50
TRUE STRAIN (%) Fig. 4. True stress against true strain curves for fine-grained 5 at.% Yb- and 5 at.% Y-doped BaCeO3 at 1250 °C in air obtained at constant cross-head speed; the initial strain rates are shown.
For the first listed method, the steady-state stages were extrapolated (dashed lines in Figs. 2 and 3) to ε = 0, i.e., to initial microstructure [25], in order to avoid the transitory creep regimes at the beginning of the tests. A comparison of the corresponding initial strain rates (denoted by ε_ o ) yields values of n = 2.0 (Fig. 2) and 2.2 (Fig. 3). A value of 2 has been systematically reported in finegrained superplastic ceramics and metals [26,27], where deformation is achieved by grain boundary sliding. It is also predicted by most theoretical models based on grain boundary sliding to explain superplasticity [26]. In this mechanism, the macroscopic deformation of the specimen is accommodated by the sliding of the grains over each other without appreciable deformation of the grains themselves, allowing very large strains (even several hundred per cent) to be achieved. The stress exponent has been also assessed by means of sudden load changes during the steady-state deformation of a sample. Fig. 3 shows two determinations of n by this technique, yielding again values of nearly 2. It is important to note that the maintenance of the strain rate levels after the positive and negative load changes indicates that no microstructural evolution took place during steadystate deformation. Finally, the stress exponent was deduced from the individual _ slopes α of the ε–ε plots. Assuming an ideal homogeneous deformation of the specimen, the instantaneous sample volume V = S.l (where S is the instantaneous cross-section) is equal to the initial volume Vo = So · lo. If Q is the constant applied load, the true stress can be expressed as σ = Q/S = Q · l/So · lo = σo exp(−ε) where the definition of true strain has been used. Then, the steady-state deformation rate ε_ is written by using Eq. (3): n n _ ε∝σ ¼ σ o expð−n:ε Þ
289
ð4Þ
_ ε plot is This equation demonstrates that the slope α of the ε– equal to −n. As shown in Figs. 2 and 3, values of n of about 2 were also measured by this method. It can be safely concluded, therefore, that the creep behavior of fine-grained Yb-doped barium cerate is characterized by a stress exponent n = 2, indicating that grain boundary sliding is the main deformation mechanism, as found in other superplastic materials [26,27]. This conclusion is in agreement with the absence of
4. Conclusions Polycrystals of 5 at.% Yb-doped barium cerate BaCe0.95Yb0.05O3 − δ have been produced by solid state reaction with a relative bulk density of 90–95%. A single orthorhombic perovskite phase was obtained after sintering at 1500 °C for 10 h. The microstructure is very homogeneous and equiaxed, with grains of average size of 0.4 μm. The orthorhombic unit cell volume is barely perturbed by doping, suggesting that dopant cations have exclusively occupied the Ce 4+site. Creep tests at constant load have been carried out in air at temperatures between 1150 and 1250 °C. The creep curves exhibit an initial sigmoidal regime with a maximum in strain rate, followed by an extensive steady state of deformation. A stress exponent n close to 2 has been estimated from different methods, indicating that grain boundary sliding is the main deformation mechanism; such a mechanism is responsible for the superplastic behavior of fine-grained ceramics and metals. Acknowledgment This work was supported by the Project no. MAT2009-13979-C03-01, Ministerio de Ciencia e Innovación, Spain. References [1] K.-D. Kreuer, Chem. Mater. 8 (1996) 610–641. [2] H. Iwahara, Solid State Ionics 86–88 (1996) 9–15. [3] S.C. Singhal, K. Kendall, High Temperature Solid Oxide Fuel Cell. Fundamentals, Design and Applications, Elsevier, Oxford, 2003. [4] N. Ito, in: T. Ishihara (Ed.), Perovskite Oxide for Solid Oxide Fuel Cells, Springer, 2009, pp. 65–94. [5] W.G. Coors, Solid State Ionics 178 (2007) 481–485. [6] J. Guan, S.E. Dorris, U. Balachandran, M. Liu, Solid State Ionics 100 (1997) 45–52. [7] C.H. Chen, H.J.M. Bouwmeester, R.H.E. van Doorn, H. Kruidhof, A.J. Burggraaf, Solid State Ionics 98 (1997) 7–13. [8] S. Imashuku, T. Uda, Y. Nose, K. Kishida, S. Harada, H. Inui, Y. Awakura, J. Electrochem. Soc. 155 (2008) B581–B586. [9] C. Zhang, H. Zhao, S. Zhai, Int. J. Hydrogen Energy 36 (2011) 3649–3657. [10] E. Gorbova, V. Maragou, D. Medvedev, A. Demin, P. Tsiakaras, J. Power Sources 181 (2008) 207–213. [11] N. Osman, I.A. Talib, H.A. Hamid, A.M. Jani, Ionics 14 (2008) 407–413. [12] J. Wu, L.P. Li, W.T.P. Espinosa, S.M. Haile, J. Mater. Res. 19 (2004) 2366–2376. [13] J. Wu, R.A. Davies, M.S. Islam, S.M. Haile, Chem. Mater. 17 (2005) 846–851. [14] K.S. Knight, N. Bonanos, Mater. Res. Bull. 30 (1995) 347–356. [15] M. Amsif, D. Marrero-López, A. Magrasó, J. Pena-Martínez, J.C. Ruiz-Morales, P. Núñez, J. Eur. Ceram. Soc. 29 (2009) 131–138. [16] J. Wu, S.M. Webb, S. Brennan, S.M. Haile, J. Appl. Phys. 97 (1–7) (2005) 054101. [17] E.T. Park, K.C. Goretta, A.R. de Arellano-López, J. Guan, U. Balachandran, S.E. Dorris, Solid State Ionics 117 (1999) 323–330.
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