- Email: [email protected]
Oxidation behavior of Fe–16Cr alloy interconnect for SOFC under hydrogen potential gradient
Solid State Ionics 168 (2004) 13 – 21 www.elsevier.com/locate/ssi
Oxidation behavior of Fe–16Cr alloy interconnect for SOFC under hydrogen potential gradient Hideto Kurokawa *, Kenichi Kawamura, Toshio Maruyama Department of Metallurgy and Ceramics Science, Graduate School of Science and Engineering, Tokyo Institute of Technology, 2-12-1 O-okayama, Meguro, Tokyo 152-8552, Japan Received 15 September 2003; received in revised form 9 February 2004; accepted 10 February 2004
Abstract Iron – chromium heat-resisting alloy is a candidate for interconnect of planar-type solid oxide fuel cells (SOFC). The oxidation kinetics of Fe – 16Cr alloy (SUS430) has been studied in an air/alloy/fuel (Ar – H2 – H2O gas mixture) atmosphere under a hydrogen potential gradient at 1073 K. The growth rates of chromia scale at both sides are almost the same, although oxygen partial pressures of both sides are quite different and a small amount of hydrogen permeates from the fuel side to the air side. The growth rate and microstructure of the scale are practically unaffected by hydrogen potential gradient in Fe – 16Cr alloy under the condition simulating SOFC at 1073 K. D 2004 Elsevier B.V. All rights reserved. PACS: 82.47.Ed Keywords: Solid oxide fuel cells; Alloy interconnect; High temperature oxidation; Hydrogen potential gradient; Fe – Cr steel; Chromium oxide scales
1. Introduction Solid oxide fuel cells (SOFC) has been extensively studied as an efficient power generation system. Since unit cell voltage is approximately 1 V, unit cells are connected electrically in series by interconnects in SOFC stacks to generate the desired power output. The property of the interconnect has to meet very severe demands as follows. Oxygen partial pressure of air side (cathodic condition) is about 2.1 104 Pa and that of fuel side (anodic condition) is 1.2 10 22 – 2.6 10 13 Pa when fuel gas is H2 – H2O gas mixture in the temperature range of 873 – 1273 K. Under the reducing and oxidative conditions, the interconnect has to posses the following properties: chemical stability both in the air side and the fuel side at operation temperatures of 873– 1273 K, high electronic conductivity, gas tightness, good machinability, and thermal expansion coefficient close to other ceramics
* Corresponding author. Tel.: +81-3-5734-3311; fax: +81-3-5734-2874. E-mail address: [email protected] (H. Kurokawa). 0167-2738/$ - see front matter D 2004 Elsevier B.V. All rights reserved. doi:10.1016/j.ssi.2004.02.008
components including yttria-stabilized zirconia electrolyte (g10 10 6 K 1). There are two major candidates for the interconnect of SOFC. The one is LaCrO3-based ceramic, and the other is Fe – Cr ferritic heat resisting alloy. The LaCrO3-based ceramic interconnect, as a matter of course, is stable at high temperature up to 1273 K, and its thermal expansion coefficient is similar to other ceramics components. Therefore LaCrO3-based ceramics have been studied extensively as the interconnect materials [1– 4]. Nowadays, alloy interconnect attracts a great deal of attention for commercial use of SOFC because of gastightness, machinability, and other advantages of alloys than ceramics [5] Alloy interconnect can be used below 1073 K where Fe – Cr alloys show good oxidation resistance [6]. Oxidation behavior of Fe –Cr alloys in reducing and oxidative conditions of SOFC operating atmospheres has been studied by several authors [7– 10]. In the previous work of our group, the oxidation kinetics of Fe –16Cr alloy (SUS430) has been studied under the conditions simulating the air side and the fuel side environments in SOFC, Ar –H2 –H2O fuel gas mixtures with the values of PH2/PH2O = 94/6 and 97/3, in the temperature range of 1023 – 1173 K [11]. The study
14
H. Kurokawa et al. / Solid State Ionics 168 (2004) 13–21
gas (Ar – H2 – H2O gas mixture). The objective of this study is to make clear the oxidation behavior of Fe –16Cr alloy (SUS430) in the air/alloy/fuel atmosphere under the hydrogen potential gradient.
2. Experimental
Fig. 1. Schematic illustration of chemical potential distribution in an Fe – Cr interconnect at 1073 K.
revealed that oxide scale formed on the alloy consists of chromia and Mn – Cr double oxide in both environments, and the scale growth rates in the air side and the fuel side are almost the same at 1073 K. Nakagawa et al. [12] have studied oxidation behavior of Fe –Cr ferritic steals (2.14 –12.12 wt.%Cr) for a heat exchanger in steam power generation. The oxidation condition is the steam/steel/air environment. They measured the amount of hydrogen permeated through the samples from the steam side to the air side. The growth rates of the iron oxides at the air side were increased much higher than those in air because hydrogen permeated from the steam side affects the oxidation behavior. The interconnect of SOFC is exposed simultaneously both to air and fuel atmospheres, in which hydrogen potential gradient develops as shown in Fig. 1. We have confirmed the permeation of hydrogen through chromia forming alloy from the fuel side to the air side in the atmosphere simulating SOFC [13]. The permeating hydrogen may affect the oxidation behavior of the air side. It is important to clarify the effect of hydrogen on oxidation behavior of Fe –Cr alloy interconnect. In the previous study [11], the samples were exposed to air or fuel conditions separately as shown in Fig. 2(A). In this study, the oxidation kinetics of SUS430 has been studied in an air/alloy/fuel atmosphere as shown in Fig. 2(B) where the samples were exposed to both air and fuel
SUS430 stainless steel was used in this study, and the chemical composition of the alloy is shown in Table 1. The samples were cut to the rectangular plates, 15 15 2 mm3, and were ground with SiC abrasive paper of #320– #2000 and were finally polished with 3 –4 Am diamond paste. After polishing, the surfaces of the samples were cleaned ultrasonically with acetone and distilled water. Experimental setup for oxidation under the air/alloy/ fuel atmosphere is illustrated in Fig. 3. The apparatus is equipped with gas flowmeters, a water vapor saturator, CaO-stabilized ZrO2 oxygen sensors, and an oxidation furnace. Samples were set at the isothermal zone of the furnace. The gas mixture of argon, hydrogen and water vapor with the value of PH2/PH2O = 97/3 was generated by flowing Ar – 20%H2 gas through the water vapor saturator at 273 K. The mixed gas introduced in the one side of the reaction tube, and dry air to the other side. Flow rates of these gases were adjusted to 1.67 10 6 m3 s 1 (100 ml/min). The sample and the tube were sealed by two Pyrex glass rings of 15 mm diameter and 1 – 2 mm in thickness. The oxygen partial pressure of PH2/PH2O gas mixture was monitored by CaO-stabilized ZrO2 oxygen sensors at points of inlet and outlet of the fuel side of the furnace. Samples were oxidized at 1073 K, for 86.4, 604.8, and 1080 ks. The reaction furnace was heated at a rate of 0.083 K s 1 (5 K min 1) to 1073 K and held for appropriate time, and was cooled at a rate of 0.083 K s 1 to room temperature. After that, the furnace was opened and the sample was picked out carefully for analysis. The samples after oxidation at 1073 K in the air/alloy/ fuel atmosphere were examined by X-ray diffraction (XRD). Phases of oxide scales of both sides (air and fuel) were investigated with Cu-Ka radiation at room temperature. The samples were fractured in liquid nitrogen to observe the fractured surfaces. Microstructure observation was carried out with an electron probe microanalyzer (EPMA) and a field emission scanning electron microscope (FE-SEM). Chemical composition of the samples was determined by energy dispersive X-
Table 1 Chemical composition of SUS430 alloy Element Fe Fig. 2. Schematic diagrams of oxidation apparatus.
Mass%
Cr
Mn
Si
Ni
Al
C
P
S
82.9 16.31 0.21 0.35 0.12 0.11 0.048 0.023 0.0006
H. Kurokawa et al. / Solid State Ionics 168 (2004) 13–21
15
Fig. 3. Schematic illustration of experimental setup for oxidation test of alloy interconnect.
ray spectroscopy (EDS). Oxidation rate was measured by average thickness of oxide scales.
3. Results 3.1. Scale morphology and composition Fig. 4 shows X-ray diffraction patterns of both sides of the sample oxidized at 1073 K for 86.4 ks. The scales on both sides of SUS430 are composed of Cr2O3, MnCr2O4 spinel and a small amount of FeCr2O4 spinel. Cristobalite was also identified in these X-ray diffraction patterns of both sides which peaks came from glass seal remaining on the sample surface. Secondary electron micrographs and X-ray maps of both sides of SUS430 alloy after oxidation at 1073 K for 86.4 ks are shown respectively in Figs. 5 and 6. In Fig. 5, oxygen map indicates that whole surface of the alloy
was covered with thin dense oxide and these element maps indicate that surface of the alloy was covered with mainly continuous Cr2O3 and dispersed MnCr2O4. Cuboidal oxide crystals, which contain Cr and Mn, were observed on the surface of the scale. Distribution of these crystals corresponded to the alloy grain boundaries on the air side. Based on the X-ray diffraction pattern, these cuboidal crystals are MnCr2O4. As is the case with the air side, surface of the fuel side was covered with mainly Cr2O3 including MnCr2O4. There is a network where manganese was detected strongly on the fuel side, meaning that MnCr2O4 was precipitated on the surface of the scale corresponding to the alloy grain boundaries. The size of Cr2O3 crystals on the fuel side was smaller than that of the air side, and planar or needle-like crystals were observed at the fuel side. The surfaces of the scales on both sides were observed with an FE-SEM after oxidation at 1073 K for 86.4 and 604.8 ks, and are shown in Fig. 7. Cuboidal MnCr2O4
Fig. 4. X-ray diffraction patterns of SUS430 after oxidation at 1073 K for 86.4 ks.
16
H. Kurokawa et al. / Solid State Ionics 168 (2004) 13–21
Fig. 5. Secondary electron micrographs and X-ray maps of a surface morphology (air side) of SUS430 after oxidation at 1073 K for 86.4 ks.
Fig. 6. Secondary electron micrographs and X-ray maps of a surface morphology (fuel side) of SUS430 after oxidation at 1073 K for 86.4 ks.
a
Fig. 7. Surface morphology of SUS430 after oxidation at 1073 K (a) in air for 86.4, (b) in air for 604.8 ks, (c) in fuel for 86.4 and (d) in fuel for 604.8 ks.
Fig. 8. Secondary electron micrographs and the results of EDS quantitative analysis of air side after oxidation at 1073 K for 604.8 ks.
H. Kurokawa et al. / Solid State Ionics 168 (2004) 13–21
crystals which precipitate on the Cr2O3 scale grew on the air side. Oxide whiskers on the top of the scale on the fuel side became larger with increasing the time of oxidation. Fig. 8 shows secondary electron micrographs and the results of EDS quantitative analysis on the surface and fractured surface of the scale at the air side after oxidation for 604.8 ks. The ratios of Fe, Cr and Mn (at.%) at several points are described in this figure. The values of (b) – (f) indicate that small crystals that cover surface thoroughly are mainly Cr2O3 including MnCr2O4 and FeCr2O4. At the point of (b), Fe in the alloy was detected because thickness of this scale is as thin as 1 Am. Iron in the alloy was also detected at the point of (e) because this point is very close to the scale/alloy boundary. The values of (a) and (g) indicate that cuboidal crystals which appear on the surface of the scale are MnCr2O4. Secondary electron micrographs and the results of EDS quantitative analysis on the surface and fractured surface of the scale at the fuel side after oxidation for 604.8 ks are shown in Fig. 9. The scale was composed of Cr2O3, MnCr2O4 and FeCr2O4. Oxide whisker with peculiar shape was observed on the top of the scale. Grain size of oxides is smaller than that of the air side. At the points of (a) and (b), Mn was detected strongly on the surface of the scale corresponding to the alloy grain boundaries. At the point of (g), Fe in the alloy was detected because this point is very close to the scale/alloy boundary. 3.2. Oxidation kinetics The oxidation kinetics of SUS430 was evaluated by average thickness of oxide scales on both sides. Average thickness and standard deviation of thickness were calculated with that values of 10 points selected randomly after tracing scales in fractured secondary electron micrographs. Scale thicknesses were plotted against oxidation time as
17
shown in Fig. 12. The kinetics was well described by the parabolic rate law [14] x2 ¼ 2kp t
ð1Þ
where x is the scale thickness, t is the oxidation time and kp is the parabolic rate constant. The parabolic rate constants kp of the air side and the fuel side were calculated. air side: kp ¼ 7:7 1019 ½m2 s1
ð2Þ
fuel side: kp ¼ 8:2 1019 ½m2 s1
ð3Þ
4. Discussion 4.1. Microstructure of the scale According to X-ray diffraction and microstructure analysis, the alloy was covered with dense oxide scale composed of Cr2O3, MnCr2O4 and a small amount of FeCr2O4. At the air side, MnCr2O4 crystals were observed on top of the scale layer which was mainly composed of Cr2O3. These crystals were precipitated on the surface of the scale corresponding to the alloy grain boundaries. At the fuel side, Cr2O3 whisker was observed on the top of the scale and MnCr2O4 was also precipitated on the surface of the scale corresponding to the alloy grain boundaries. This microstructure developed under the hydrogen potential gradient was in good agreement with that in the previous study [11] in which oxidation was carried out separately in the air side and the fuel side. Formation of oxides on the both sides can be explained by thermodynamic stability. Fig. 10 shows dissociation pressures of FeO, Cr2O3, MnO, FeCr2O4 and MnCr2O4 calculated from reported thermodynamic data [15]. The calculation took activities of metal elements into account
Fig. 9. Secondary electron micrographs and the results of EDS quantitative analysis of fuel side after oxidation at 1073 K for 604.8 ks.
18
H. Kurokawa et al. / Solid State Ionics 168 (2004) 13–21
Fig. 10. Equilibrium oxygen pressures of some metal/oxide and Ar – H2 – H2O gas mixture.
[16 –19]. The oxygen partial pressure of hydrogen and water vapor mixture with the value of PH2/PH2O = 97/3 is also shown. The standard Gibbs energy of formation for MnCr2O4 had not been reported so that the value was estimated by following manner. Gibbs free energy change of MnCr2O4 from MnO and Cr2O3 is assumed to be similar to that of FeCr2O4 from FeO and Cr2O3 [20]. This diagram elucidates that Cr2O3, FeCr2O4 and MnCr2O4 are stable even in the atmosphere of the fuel side and iron is oxidized at the air side but is not oxidized at the fuel side ( PH2/PH2O = 97/3). Iron oxide can be formed on the air side at the early stage of oxidation before the protective Cr2O3 scale is formed. The iron oxides react with the Cr2O3 scale and provide FeCr2O4 at the later stage. Fig. 11 shows schematic equilibrium oxygen pressure – composition diagram of Cr – Mn – O system at 1073 K,
Fig. 11. Schematic equilibrium oxygen pressures – composition diagram of Cr – Mn – O system.
which is estimated based on reported phase diagrams [21,22]. The equilibrium oxygen pressure of Cr/MnO/ MnCr2O4 at 1073 K is obtained from Fig. 10. This diagram indicates thermodynamically that MnCr2O4 spinel can appear on the alloy in spite of low Mn concentration. There were many cuboidal MnCr2O4 crystals on the scale of the air side. Molar ratios of Mn at the Cr2O3 grain boundaries at the points of (c) and (f) were larger than that at the inside of Cr2O3 grain at the points of (d) and (e) (see Fig. 8). This means that a slight amount of Mn in SUS430 (see Table 1) diffuses from the alloy to the surface of the scales and Mn diffusion at the grain boundary of oxide is faster than bulk diffusion. In case of oxidation of Fe – Cr based alloy including Mn, (Mn, Cr)3O4 spinel layer formed on top of the chromia scale after 3600 ks oxidation in air at 1073 K [9]. Cox et al. [23] showed diffusivities of metal ions in Cr2O3 should decrease in the order DMn>DFe>DNi>DCr by assuming that the metals diffuse as ions via Cr3 +-lattice sites in Cr2O3. From Auger Electron Spectroscopy measurements on growing Cr2O3 scale, Wild [24] concluded that Mn ions diffuses two orders of magnitude faster than Cr ions. Lobnig et al. [25] confirmed the very fast diffusion of Mn ions in Cr2O3 scales by a direct diffusion experiment. Because of the fast diffusion of Mn ions in oxide scales from alloy to the scale surface, MnCr2O4 spinel is formed on oxide scales mainly composed of Cr2O3. 4.2. Oxidation kinetics Since the oxidation kinetics of SUS430 obeyed the parabolic rate law as shown in Fig. 12, the growth of Cr2O3 scale was controlled by a diffusion mechanism. The parabolic rate constants of present and previous studies [11] are shown in Fig. 13. The values of kp were similar to the those previously obtained from separate oxidation experiments in the atmospheres of the air side and the fuel side of SOFC. Temperature dependence of the parabolic rate constant kp was expressed with the activation energy of 202.3 kJ mol 1 [11], which is in good agreement with that for self-diffusion of Cr cation in Cr2O3 [26]. This fact indicates
Fig. 12. Scale thickness of SUS430 after oxidation in air and fuel at 1073 K.
H. Kurokawa et al. / Solid State Ionics 168 (2004) 13–21
19
where DCr is the tracer diffusion coefficient of chromium in Cr2O3, R is the gas constant, T is the absolute temperature, PO2 is the oxygen partial pressure and P* is the standard pressure of 1.0 105 Pa. The flux of chromium in Cr2O3 scale is given by JCr ¼
DCr CCr dlCr RT dx
ð5Þ
where JCr is the flux of chromium, CCr is the concentration of chromium in Cr2O3 and lCr is the chemical potential of chromium. The Gibbs –Duhem relation gives 3 dlCr ¼ dlO2 4
ð6Þ
where lO2 is the chemical potential of oxygen. The flux of chromium in Cr2O3 is also written as Fig. 13. Logarithmic kp values for parabolic oxidation of SUS430.
that the oxidation kinetics of SUS430 alloy is controlled by an outward diffusion of Cr cations in the Cr2O3 scale. In this study, the parabolic rate constant kp of the fuel side is almost the same as that of the air side at 1073 K although the oxygen partial pressures of the both sides are quite different. Therefore, the parabolic rate constants of SUS430 alloy are independent of oxygen partial pressure. The diffusion coefficient of chromium interstitials considerably exceeds that of chromium vacancies at the temperature lower than 1373 K [11]. Positively ionized chromium interstitial is predominant point defect when Cr2O3 exhibits n-type behavior, and negatively ionized chromium vacancy is predominant point defect when Cr2O3 exhibits p-type behavior. Several authors reported n-type semiconducting behaviors of chromia at low oxygen partial pressures and low temperatures [27 – 29], although it is well known that chromia indicates p-type semiconducting behaviors in air at high temperatures. Oxygen partial pressure dependence of the parabolic rate constant kp can be estimated with diffusion coefficient of chromium ion in Cr2O3. Yurek [27] has compiled reported data of chromium tracer diffusion coefficients in single crystal of Cr2O3 [30,31], and has showed an approximate expression for the dependence of the chromium diffusion coefficient on temperature and oxygen partial pressure as the following equation PO2 3=16 PO2 3=16 DCr ½m2 s1 ¼ A þB P* P*
A ¼ 1:2 10
10
397½kJ mol1 exp RT
dx : dt
ð7Þ
According to Wagner’s theory [14], the parabolic rate constant is given by kp dx ¼ : dt x
ð8Þ
From Eqs. (4) – (8), the parabolic rate constant kp is arranged as o 3lO2 3l kp ¼ 4 Aexp exp O2 16RT 16RT I 3lO2 3loO2 3lO2 þ exp þ Bexp exp 16RT 16RT 16RT I 3lO2 exp ð9Þ 16RT where lOo 2 is the standard chemical potential of oxygen and lOI 2 is the chemical potential of oxygen at alloy/scale interface (I). Fig. 14 shows oxygen partial pressure de-
ð4Þ
502½kJ mol1 B ¼ 1:5 10 exp RT 2
JCr ¼ CCr
Fig. 14. Dependence of parabolic rate constant kp on oxygen partial pressure.
20
H. Kurokawa et al. / Solid State Ionics 168 (2004) 13–21
Table 2 Change of the hydrogen flux through alloy interconnect Time[ks] Hydrogen flux [mol m 2 s 1]
0
36 4
1.0 10
5. Conclusions
360 5
2.9 10
3600 5
1.2 10
4.0 10 6
pendence of the parabolic rate constant kp at 1073 – 1473 K. At high oxygen pressure at 1473 K, the parabolic rate constant is increasing with increasing oxygen partial pressure. In this atmosphere, chromium vacancy is the predominant point defect responsible for diffusion of chromium. The predominant point defect changes from negatively ionized chromium vacancy to positively ionized chromium interstitial at lower oxygen pressure. On the other hand, the parabolic rate constant is independent of oxygen partial pressure over 20 orders of magnitude at 1073 K, and those of both the air side (o) and the fuel side (.) is almost the same. Diffusion of chromium in chromia scale was dominated by positively ionized chromium interstitial at 1073 K. The experimental values of kp of both sides at 1073 K (Fig. 13) are approximately six orders of magnitude larger than calculated one (Fig. 14). Lobnig et al. [25] have described that the chromium diffusion coefficients determined with chromium single crystals are several orders of magnitude smaller than those obtained with polycrystalline materials because of the faster grainboundary diffusion. Stubican and Carinci [32] reported that, although grain boundary diffusivity was much larger than volume diffusivity, the oxygen partial pressure dependence of both diffusivities was similar to each other in NiO and Fe3O4. The result of this study implies that the similar situation may appear in the chromium oxide scale. Hydrogen permeability of oxide scale which is mainly composed of Cr2O3 formed on SUS430 at 1073 K have been reported to be approximately 1.3 10 13 [mol m 1 s 1 Pa 1/2] [13]. Using this permeability, the amount of hydrogen permeated through the sample in this experimental condition (thickness of alloy: 2 mm, temperature: 1073 K, hydrogen partial pressure of the fuel side: 2.0 104 Pa) was calculated and shown in Table 2. Nakagawa et al. [12] have reported that hydrogen permeating through the Fe – Cr ferritic steals accelerated the growth rate of iron-oxide scale of the air side. In this study, although the samples of chromia forming alloy were exposed to the air/alloy/fuel atmosphere under the hydrogen potential gradient, surface morphology and microstructure of scales on both the air side and the fuel side are similar to those in previous study where samples were oxidized in both atmospheres separately at 1073 K. It is concluded that oxidation behavior of Fe –Cr alloy are practically unaffected by hydrogen potential gradient under the condition simulating SOFC at 1073 K.
The dense Cr2O3 scale including MnCr2O4 and a small amount of FeCr2O4 was observed on the surface of SUS430 after oxidation both in the air and fuel atmospheres of SOFC at 1073 K. The parabolic rate constant kp of the fuel side of the alloy is similar to that of the air side at 1073 K. The oxygen partial pressure dependence of the parabolic rate constant revealed that the growth rate of Cr2O3 scale is almost the same in various oxygen partial pressures ( PO 2g4.1 10 17 – 2.1 104 Pa) at 1073 K. Under the hydrogen potential gradient, the scale morphology of both the air side and the fuel side were almost similar to those which were oxidized in both atmospheres separately without hydrogen potential gradient. There is little effect of hydrogen potential gradient on oxidation behavior of Fe – Cr alloy in the atmosphere simulating SOFC.
Acknowledgements The authors thank Dr. Harumi Yokokawa, National Institute of Advanced Industrial Science and Technology, for valuable discussions of thermodynamics in spinel.
References [1] M. Dokiya, N. Sakai, N. Kawada, T. Yokokawa, T. Iwata, M. Mori, Intersoc. Energy Convers. Eng. 24 (1989) 1547. [2] J. Mizusaki, S. Yamauchi, K. Fueki, A. Ishikawa, Solid State Ion. 12 (1984) 119. [3] T.R. Armstrong, J.W. Stevenson, L.R. Pederson, P.E. Raney, J. Electrochem. Soc. 143 (1996) 2919. [4] I. Yasuda, M. Hishinuma, J. Electrochem. Soc. 143 (1996) 1583. [5] T. Kadowaki, T. Shiomitsu, E. Matsuda, H. Nakagawa, H. Tsuneizumi, T. Maruyama, Solid State Ion. 67 (1993) 65. [6] K. Huang, P.Y. Hou, J.B. Goodenough, Solid State Ion. 129 (2000) 237. [7] P. Kofstad, R. Bredesen, Solid State Ion. 52 (1992) 69. [8] K. Honegger, A. Plas, R. Diethelm, W. Glatz, in: H. Yokokawa, S.C. Singhal (Eds.), Solid Oxide Fuel Cells (SOFC), vol. VII, The Electrochemical Society, Pennington, NJ, 2001, p. 803, PV2001-16. [9] J.P. Abella´n, V. Shemet, F. Tietz, L. Singheiser, W.J. Quadakkers, A. Gil, in: H. Yokokawa, S.C. Singhal (Eds.), Solid Oxide Fuel Cells (SOFC), vol. VII, The Electrochemical Society, Pennington, NJ, 2001, p. 811, PV2001-16. [10] T. Horita, Y. Xiong, K. Yamaji, N. Sakai, H. Yokokawa, J. Electrochem. Soc. 150 (2003) A243. [11] T. Brylewski, M. Nanko, T. Maruyama, K. Przybylski, Solid State Ion. 143 (2001) 131. [12] K. Nakagawa, Y. Matsunaga, T. Yanagisawa, Mater. High Temp. 18 (2001) 51. [13] H. Kurokawa, Y. Oyama, K. Kawamura, T. Maruyama, J. Electrochem. Soc (submitted to). [14] C. Wagner, Z. Phys. Chem. 21 (1933) 25. [15] I. Barin, Thermochemical Data of Pure Substances, 3rd edition, VCH Verlagsgesellschaft, Weinheim, 1995. [16] Y. Jeannin, C. Mannerskantz, F.D. Richardson, Trans. Metall. Soc. AIME 227 (1963) 300.
H. Kurokawa et al. / Solid State Ionics 168 (2004) 13–21 [17] P.C. Lidster, H.B. Bell, Trans. Metall. Soc. AIME 245 (1969) 2273. [18] C.L. McCabe, R.G. Hudson, H.W. Paxton, Trans. Metall. Soc. AIME 212 (1958) 102. [19] W. Dresler, Trans. Inst. Min. Metall. 98 (1989) C61. [20] H. Yokokawa, (private communication). [21] B.-J. Lee, Metall. Trans., A, Phys. Metall. Mater. Sci. 24 (1993) 1919. [22] D.H. Speidel, Arnulf Muan, J. Am. Ceram. Soc. 46 (12) (1963) 578. [23] M.G.E. Cox, B. McEnanay, V.D. Scott, Phila. Mag. 26 (1972) 839. [24] R.K. Wild, Corros. Sci. 17 (1977) 87. [25] R.E. Lobnig, H.P. Schmidt, K. Hennesen, H.J. Grabke, Oxid. Met. 37 (1992) 81. [26] W.C. Hagel, A.W. Seybolt, J. Electrochem. Soc. 108 (1961) 1246.
21
[27] G.J. Yurek, Ann. Prog. Rept. (July 1, 1983 – June 30,1984), DOE No. DE-AC-02-79ER-10507, Dept of Energy, Gaithersburg, MD, Feb. 1985. [28] K. Przybylski, G.J. Yurek, Mat. Sci. Forum 43 (1989) 1. [29] E.W.A. Young, P.C.M. Stiphout, J.H.W. de Wit, J. Electrochem. Soc. 132 (4) (1989) 884. [30] A. Atkinson, R.I. Taylor, in: G. Simkovich, V.S. Stubican (Eds.), Transport in Nonstoichiometric Compounds, Nato ASI Series B, vol. 129, Plenum, New York, 1985, p. 285. [31] K. Hoshino, N.L. Peterson, J. Am. Ceram. Soc. 66 (1983) c202. [32] V.S. Stubican, L.R. Carinci, Z. Phys. Chem. 207 (1998) 215.
Oxidation behavior of Fe–16Cr alloy interconnect for SOFC under hydrogen potential gradient Hideto Kurokawa *, Kenichi Kawamura, Toshio Maruyama Department of Metallurgy and Ceramics Science, Graduate School of Science and Engineering, Tokyo Institute of Technology, 2-12-1 O-okayama, Meguro, Tokyo 152-8552, Japan Received 15 September 2003; received in revised form 9 February 2004; accepted 10 February 2004
Abstract Iron – chromium heat-resisting alloy is a candidate for interconnect of planar-type solid oxide fuel cells (SOFC). The oxidation kinetics of Fe – 16Cr alloy (SUS430) has been studied in an air/alloy/fuel (Ar – H2 – H2O gas mixture) atmosphere under a hydrogen potential gradient at 1073 K. The growth rates of chromia scale at both sides are almost the same, although oxygen partial pressures of both sides are quite different and a small amount of hydrogen permeates from the fuel side to the air side. The growth rate and microstructure of the scale are practically unaffected by hydrogen potential gradient in Fe – 16Cr alloy under the condition simulating SOFC at 1073 K. D 2004 Elsevier B.V. All rights reserved. PACS: 82.47.Ed Keywords: Solid oxide fuel cells; Alloy interconnect; High temperature oxidation; Hydrogen potential gradient; Fe – Cr steel; Chromium oxide scales
1. Introduction Solid oxide fuel cells (SOFC) has been extensively studied as an efficient power generation system. Since unit cell voltage is approximately 1 V, unit cells are connected electrically in series by interconnects in SOFC stacks to generate the desired power output. The property of the interconnect has to meet very severe demands as follows. Oxygen partial pressure of air side (cathodic condition) is about 2.1 104 Pa and that of fuel side (anodic condition) is 1.2 10 22 – 2.6 10 13 Pa when fuel gas is H2 – H2O gas mixture in the temperature range of 873 – 1273 K. Under the reducing and oxidative conditions, the interconnect has to posses the following properties: chemical stability both in the air side and the fuel side at operation temperatures of 873– 1273 K, high electronic conductivity, gas tightness, good machinability, and thermal expansion coefficient close to other ceramics
* Corresponding author. Tel.: +81-3-5734-3311; fax: +81-3-5734-2874. E-mail address: [email protected] (H. Kurokawa). 0167-2738/$ - see front matter D 2004 Elsevier B.V. All rights reserved. doi:10.1016/j.ssi.2004.02.008
components including yttria-stabilized zirconia electrolyte (g10 10 6 K 1). There are two major candidates for the interconnect of SOFC. The one is LaCrO3-based ceramic, and the other is Fe – Cr ferritic heat resisting alloy. The LaCrO3-based ceramic interconnect, as a matter of course, is stable at high temperature up to 1273 K, and its thermal expansion coefficient is similar to other ceramics components. Therefore LaCrO3-based ceramics have been studied extensively as the interconnect materials [1– 4]. Nowadays, alloy interconnect attracts a great deal of attention for commercial use of SOFC because of gastightness, machinability, and other advantages of alloys than ceramics [5] Alloy interconnect can be used below 1073 K where Fe – Cr alloys show good oxidation resistance [6]. Oxidation behavior of Fe –Cr alloys in reducing and oxidative conditions of SOFC operating atmospheres has been studied by several authors [7– 10]. In the previous work of our group, the oxidation kinetics of Fe –16Cr alloy (SUS430) has been studied under the conditions simulating the air side and the fuel side environments in SOFC, Ar –H2 –H2O fuel gas mixtures with the values of PH2/PH2O = 94/6 and 97/3, in the temperature range of 1023 – 1173 K [11]. The study
14
H. Kurokawa et al. / Solid State Ionics 168 (2004) 13–21
gas (Ar – H2 – H2O gas mixture). The objective of this study is to make clear the oxidation behavior of Fe –16Cr alloy (SUS430) in the air/alloy/fuel atmosphere under the hydrogen potential gradient.
2. Experimental
Fig. 1. Schematic illustration of chemical potential distribution in an Fe – Cr interconnect at 1073 K.
revealed that oxide scale formed on the alloy consists of chromia and Mn – Cr double oxide in both environments, and the scale growth rates in the air side and the fuel side are almost the same at 1073 K. Nakagawa et al. [12] have studied oxidation behavior of Fe –Cr ferritic steals (2.14 –12.12 wt.%Cr) for a heat exchanger in steam power generation. The oxidation condition is the steam/steel/air environment. They measured the amount of hydrogen permeated through the samples from the steam side to the air side. The growth rates of the iron oxides at the air side were increased much higher than those in air because hydrogen permeated from the steam side affects the oxidation behavior. The interconnect of SOFC is exposed simultaneously both to air and fuel atmospheres, in which hydrogen potential gradient develops as shown in Fig. 1. We have confirmed the permeation of hydrogen through chromia forming alloy from the fuel side to the air side in the atmosphere simulating SOFC [13]. The permeating hydrogen may affect the oxidation behavior of the air side. It is important to clarify the effect of hydrogen on oxidation behavior of Fe –Cr alloy interconnect. In the previous study [11], the samples were exposed to air or fuel conditions separately as shown in Fig. 2(A). In this study, the oxidation kinetics of SUS430 has been studied in an air/alloy/fuel atmosphere as shown in Fig. 2(B) where the samples were exposed to both air and fuel
SUS430 stainless steel was used in this study, and the chemical composition of the alloy is shown in Table 1. The samples were cut to the rectangular plates, 15 15 2 mm3, and were ground with SiC abrasive paper of #320– #2000 and were finally polished with 3 –4 Am diamond paste. After polishing, the surfaces of the samples were cleaned ultrasonically with acetone and distilled water. Experimental setup for oxidation under the air/alloy/ fuel atmosphere is illustrated in Fig. 3. The apparatus is equipped with gas flowmeters, a water vapor saturator, CaO-stabilized ZrO2 oxygen sensors, and an oxidation furnace. Samples were set at the isothermal zone of the furnace. The gas mixture of argon, hydrogen and water vapor with the value of PH2/PH2O = 97/3 was generated by flowing Ar – 20%H2 gas through the water vapor saturator at 273 K. The mixed gas introduced in the one side of the reaction tube, and dry air to the other side. Flow rates of these gases were adjusted to 1.67 10 6 m3 s 1 (100 ml/min). The sample and the tube were sealed by two Pyrex glass rings of 15 mm diameter and 1 – 2 mm in thickness. The oxygen partial pressure of PH2/PH2O gas mixture was monitored by CaO-stabilized ZrO2 oxygen sensors at points of inlet and outlet of the fuel side of the furnace. Samples were oxidized at 1073 K, for 86.4, 604.8, and 1080 ks. The reaction furnace was heated at a rate of 0.083 K s 1 (5 K min 1) to 1073 K and held for appropriate time, and was cooled at a rate of 0.083 K s 1 to room temperature. After that, the furnace was opened and the sample was picked out carefully for analysis. The samples after oxidation at 1073 K in the air/alloy/ fuel atmosphere were examined by X-ray diffraction (XRD). Phases of oxide scales of both sides (air and fuel) were investigated with Cu-Ka radiation at room temperature. The samples were fractured in liquid nitrogen to observe the fractured surfaces. Microstructure observation was carried out with an electron probe microanalyzer (EPMA) and a field emission scanning electron microscope (FE-SEM). Chemical composition of the samples was determined by energy dispersive X-
Table 1 Chemical composition of SUS430 alloy Element Fe Fig. 2. Schematic diagrams of oxidation apparatus.
Mass%
Cr
Mn
Si
Ni
Al
C
P
S
82.9 16.31 0.21 0.35 0.12 0.11 0.048 0.023 0.0006
H. Kurokawa et al. / Solid State Ionics 168 (2004) 13–21
15
Fig. 3. Schematic illustration of experimental setup for oxidation test of alloy interconnect.
ray spectroscopy (EDS). Oxidation rate was measured by average thickness of oxide scales.
3. Results 3.1. Scale morphology and composition Fig. 4 shows X-ray diffraction patterns of both sides of the sample oxidized at 1073 K for 86.4 ks. The scales on both sides of SUS430 are composed of Cr2O3, MnCr2O4 spinel and a small amount of FeCr2O4 spinel. Cristobalite was also identified in these X-ray diffraction patterns of both sides which peaks came from glass seal remaining on the sample surface. Secondary electron micrographs and X-ray maps of both sides of SUS430 alloy after oxidation at 1073 K for 86.4 ks are shown respectively in Figs. 5 and 6. In Fig. 5, oxygen map indicates that whole surface of the alloy
was covered with thin dense oxide and these element maps indicate that surface of the alloy was covered with mainly continuous Cr2O3 and dispersed MnCr2O4. Cuboidal oxide crystals, which contain Cr and Mn, were observed on the surface of the scale. Distribution of these crystals corresponded to the alloy grain boundaries on the air side. Based on the X-ray diffraction pattern, these cuboidal crystals are MnCr2O4. As is the case with the air side, surface of the fuel side was covered with mainly Cr2O3 including MnCr2O4. There is a network where manganese was detected strongly on the fuel side, meaning that MnCr2O4 was precipitated on the surface of the scale corresponding to the alloy grain boundaries. The size of Cr2O3 crystals on the fuel side was smaller than that of the air side, and planar or needle-like crystals were observed at the fuel side. The surfaces of the scales on both sides were observed with an FE-SEM after oxidation at 1073 K for 86.4 and 604.8 ks, and are shown in Fig. 7. Cuboidal MnCr2O4
Fig. 4. X-ray diffraction patterns of SUS430 after oxidation at 1073 K for 86.4 ks.
16
H. Kurokawa et al. / Solid State Ionics 168 (2004) 13–21
Fig. 5. Secondary electron micrographs and X-ray maps of a surface morphology (air side) of SUS430 after oxidation at 1073 K for 86.4 ks.
Fig. 6. Secondary electron micrographs and X-ray maps of a surface morphology (fuel side) of SUS430 after oxidation at 1073 K for 86.4 ks.
a
Fig. 7. Surface morphology of SUS430 after oxidation at 1073 K (a) in air for 86.4, (b) in air for 604.8 ks, (c) in fuel for 86.4 and (d) in fuel for 604.8 ks.
Fig. 8. Secondary electron micrographs and the results of EDS quantitative analysis of air side after oxidation at 1073 K for 604.8 ks.
H. Kurokawa et al. / Solid State Ionics 168 (2004) 13–21
crystals which precipitate on the Cr2O3 scale grew on the air side. Oxide whiskers on the top of the scale on the fuel side became larger with increasing the time of oxidation. Fig. 8 shows secondary electron micrographs and the results of EDS quantitative analysis on the surface and fractured surface of the scale at the air side after oxidation for 604.8 ks. The ratios of Fe, Cr and Mn (at.%) at several points are described in this figure. The values of (b) – (f) indicate that small crystals that cover surface thoroughly are mainly Cr2O3 including MnCr2O4 and FeCr2O4. At the point of (b), Fe in the alloy was detected because thickness of this scale is as thin as 1 Am. Iron in the alloy was also detected at the point of (e) because this point is very close to the scale/alloy boundary. The values of (a) and (g) indicate that cuboidal crystals which appear on the surface of the scale are MnCr2O4. Secondary electron micrographs and the results of EDS quantitative analysis on the surface and fractured surface of the scale at the fuel side after oxidation for 604.8 ks are shown in Fig. 9. The scale was composed of Cr2O3, MnCr2O4 and FeCr2O4. Oxide whisker with peculiar shape was observed on the top of the scale. Grain size of oxides is smaller than that of the air side. At the points of (a) and (b), Mn was detected strongly on the surface of the scale corresponding to the alloy grain boundaries. At the point of (g), Fe in the alloy was detected because this point is very close to the scale/alloy boundary. 3.2. Oxidation kinetics The oxidation kinetics of SUS430 was evaluated by average thickness of oxide scales on both sides. Average thickness and standard deviation of thickness were calculated with that values of 10 points selected randomly after tracing scales in fractured secondary electron micrographs. Scale thicknesses were plotted against oxidation time as
17
shown in Fig. 12. The kinetics was well described by the parabolic rate law [14] x2 ¼ 2kp t
ð1Þ
where x is the scale thickness, t is the oxidation time and kp is the parabolic rate constant. The parabolic rate constants kp of the air side and the fuel side were calculated. air side: kp ¼ 7:7 1019 ½m2 s1
ð2Þ
fuel side: kp ¼ 8:2 1019 ½m2 s1
ð3Þ
4. Discussion 4.1. Microstructure of the scale According to X-ray diffraction and microstructure analysis, the alloy was covered with dense oxide scale composed of Cr2O3, MnCr2O4 and a small amount of FeCr2O4. At the air side, MnCr2O4 crystals were observed on top of the scale layer which was mainly composed of Cr2O3. These crystals were precipitated on the surface of the scale corresponding to the alloy grain boundaries. At the fuel side, Cr2O3 whisker was observed on the top of the scale and MnCr2O4 was also precipitated on the surface of the scale corresponding to the alloy grain boundaries. This microstructure developed under the hydrogen potential gradient was in good agreement with that in the previous study [11] in which oxidation was carried out separately in the air side and the fuel side. Formation of oxides on the both sides can be explained by thermodynamic stability. Fig. 10 shows dissociation pressures of FeO, Cr2O3, MnO, FeCr2O4 and MnCr2O4 calculated from reported thermodynamic data [15]. The calculation took activities of metal elements into account
Fig. 9. Secondary electron micrographs and the results of EDS quantitative analysis of fuel side after oxidation at 1073 K for 604.8 ks.
18
H. Kurokawa et al. / Solid State Ionics 168 (2004) 13–21
Fig. 10. Equilibrium oxygen pressures of some metal/oxide and Ar – H2 – H2O gas mixture.
[16 –19]. The oxygen partial pressure of hydrogen and water vapor mixture with the value of PH2/PH2O = 97/3 is also shown. The standard Gibbs energy of formation for MnCr2O4 had not been reported so that the value was estimated by following manner. Gibbs free energy change of MnCr2O4 from MnO and Cr2O3 is assumed to be similar to that of FeCr2O4 from FeO and Cr2O3 [20]. This diagram elucidates that Cr2O3, FeCr2O4 and MnCr2O4 are stable even in the atmosphere of the fuel side and iron is oxidized at the air side but is not oxidized at the fuel side ( PH2/PH2O = 97/3). Iron oxide can be formed on the air side at the early stage of oxidation before the protective Cr2O3 scale is formed. The iron oxides react with the Cr2O3 scale and provide FeCr2O4 at the later stage. Fig. 11 shows schematic equilibrium oxygen pressure – composition diagram of Cr – Mn – O system at 1073 K,
Fig. 11. Schematic equilibrium oxygen pressures – composition diagram of Cr – Mn – O system.
which is estimated based on reported phase diagrams [21,22]. The equilibrium oxygen pressure of Cr/MnO/ MnCr2O4 at 1073 K is obtained from Fig. 10. This diagram indicates thermodynamically that MnCr2O4 spinel can appear on the alloy in spite of low Mn concentration. There were many cuboidal MnCr2O4 crystals on the scale of the air side. Molar ratios of Mn at the Cr2O3 grain boundaries at the points of (c) and (f) were larger than that at the inside of Cr2O3 grain at the points of (d) and (e) (see Fig. 8). This means that a slight amount of Mn in SUS430 (see Table 1) diffuses from the alloy to the surface of the scales and Mn diffusion at the grain boundary of oxide is faster than bulk diffusion. In case of oxidation of Fe – Cr based alloy including Mn, (Mn, Cr)3O4 spinel layer formed on top of the chromia scale after 3600 ks oxidation in air at 1073 K [9]. Cox et al. [23] showed diffusivities of metal ions in Cr2O3 should decrease in the order DMn>DFe>DNi>DCr by assuming that the metals diffuse as ions via Cr3 +-lattice sites in Cr2O3. From Auger Electron Spectroscopy measurements on growing Cr2O3 scale, Wild [24] concluded that Mn ions diffuses two orders of magnitude faster than Cr ions. Lobnig et al. [25] confirmed the very fast diffusion of Mn ions in Cr2O3 scales by a direct diffusion experiment. Because of the fast diffusion of Mn ions in oxide scales from alloy to the scale surface, MnCr2O4 spinel is formed on oxide scales mainly composed of Cr2O3. 4.2. Oxidation kinetics Since the oxidation kinetics of SUS430 obeyed the parabolic rate law as shown in Fig. 12, the growth of Cr2O3 scale was controlled by a diffusion mechanism. The parabolic rate constants of present and previous studies [11] are shown in Fig. 13. The values of kp were similar to the those previously obtained from separate oxidation experiments in the atmospheres of the air side and the fuel side of SOFC. Temperature dependence of the parabolic rate constant kp was expressed with the activation energy of 202.3 kJ mol 1 [11], which is in good agreement with that for self-diffusion of Cr cation in Cr2O3 [26]. This fact indicates
Fig. 12. Scale thickness of SUS430 after oxidation in air and fuel at 1073 K.
H. Kurokawa et al. / Solid State Ionics 168 (2004) 13–21
19
where DCr is the tracer diffusion coefficient of chromium in Cr2O3, R is the gas constant, T is the absolute temperature, PO2 is the oxygen partial pressure and P* is the standard pressure of 1.0 105 Pa. The flux of chromium in Cr2O3 scale is given by JCr ¼
DCr CCr dlCr RT dx
ð5Þ
where JCr is the flux of chromium, CCr is the concentration of chromium in Cr2O3 and lCr is the chemical potential of chromium. The Gibbs –Duhem relation gives 3 dlCr ¼ dlO2 4
ð6Þ
where lO2 is the chemical potential of oxygen. The flux of chromium in Cr2O3 is also written as Fig. 13. Logarithmic kp values for parabolic oxidation of SUS430.
that the oxidation kinetics of SUS430 alloy is controlled by an outward diffusion of Cr cations in the Cr2O3 scale. In this study, the parabolic rate constant kp of the fuel side is almost the same as that of the air side at 1073 K although the oxygen partial pressures of the both sides are quite different. Therefore, the parabolic rate constants of SUS430 alloy are independent of oxygen partial pressure. The diffusion coefficient of chromium interstitials considerably exceeds that of chromium vacancies at the temperature lower than 1373 K [11]. Positively ionized chromium interstitial is predominant point defect when Cr2O3 exhibits n-type behavior, and negatively ionized chromium vacancy is predominant point defect when Cr2O3 exhibits p-type behavior. Several authors reported n-type semiconducting behaviors of chromia at low oxygen partial pressures and low temperatures [27 – 29], although it is well known that chromia indicates p-type semiconducting behaviors in air at high temperatures. Oxygen partial pressure dependence of the parabolic rate constant kp can be estimated with diffusion coefficient of chromium ion in Cr2O3. Yurek [27] has compiled reported data of chromium tracer diffusion coefficients in single crystal of Cr2O3 [30,31], and has showed an approximate expression for the dependence of the chromium diffusion coefficient on temperature and oxygen partial pressure as the following equation PO2 3=16 PO2 3=16 DCr ½m2 s1 ¼ A þB P* P*
A ¼ 1:2 10
10
397½kJ mol1 exp RT
dx : dt
ð7Þ
According to Wagner’s theory [14], the parabolic rate constant is given by kp dx ¼ : dt x
ð8Þ
From Eqs. (4) – (8), the parabolic rate constant kp is arranged as o 3lO2 3l kp ¼ 4 Aexp exp O2 16RT 16RT I 3lO2 3loO2 3lO2 þ exp þ Bexp exp 16RT 16RT 16RT I 3lO2 exp ð9Þ 16RT where lOo 2 is the standard chemical potential of oxygen and lOI 2 is the chemical potential of oxygen at alloy/scale interface (I). Fig. 14 shows oxygen partial pressure de-
ð4Þ
502½kJ mol1 B ¼ 1:5 10 exp RT 2
JCr ¼ CCr
Fig. 14. Dependence of parabolic rate constant kp on oxygen partial pressure.
20
H. Kurokawa et al. / Solid State Ionics 168 (2004) 13–21
Table 2 Change of the hydrogen flux through alloy interconnect Time[ks] Hydrogen flux [mol m 2 s 1]
0
36 4
1.0 10
5. Conclusions
360 5
2.9 10
3600 5
1.2 10
4.0 10 6
pendence of the parabolic rate constant kp at 1073 – 1473 K. At high oxygen pressure at 1473 K, the parabolic rate constant is increasing with increasing oxygen partial pressure. In this atmosphere, chromium vacancy is the predominant point defect responsible for diffusion of chromium. The predominant point defect changes from negatively ionized chromium vacancy to positively ionized chromium interstitial at lower oxygen pressure. On the other hand, the parabolic rate constant is independent of oxygen partial pressure over 20 orders of magnitude at 1073 K, and those of both the air side (o) and the fuel side (.) is almost the same. Diffusion of chromium in chromia scale was dominated by positively ionized chromium interstitial at 1073 K. The experimental values of kp of both sides at 1073 K (Fig. 13) are approximately six orders of magnitude larger than calculated one (Fig. 14). Lobnig et al. [25] have described that the chromium diffusion coefficients determined with chromium single crystals are several orders of magnitude smaller than those obtained with polycrystalline materials because of the faster grainboundary diffusion. Stubican and Carinci [32] reported that, although grain boundary diffusivity was much larger than volume diffusivity, the oxygen partial pressure dependence of both diffusivities was similar to each other in NiO and Fe3O4. The result of this study implies that the similar situation may appear in the chromium oxide scale. Hydrogen permeability of oxide scale which is mainly composed of Cr2O3 formed on SUS430 at 1073 K have been reported to be approximately 1.3 10 13 [mol m 1 s 1 Pa 1/2] [13]. Using this permeability, the amount of hydrogen permeated through the sample in this experimental condition (thickness of alloy: 2 mm, temperature: 1073 K, hydrogen partial pressure of the fuel side: 2.0 104 Pa) was calculated and shown in Table 2. Nakagawa et al. [12] have reported that hydrogen permeating through the Fe – Cr ferritic steals accelerated the growth rate of iron-oxide scale of the air side. In this study, although the samples of chromia forming alloy were exposed to the air/alloy/fuel atmosphere under the hydrogen potential gradient, surface morphology and microstructure of scales on both the air side and the fuel side are similar to those in previous study where samples were oxidized in both atmospheres separately at 1073 K. It is concluded that oxidation behavior of Fe –Cr alloy are practically unaffected by hydrogen potential gradient under the condition simulating SOFC at 1073 K.
The dense Cr2O3 scale including MnCr2O4 and a small amount of FeCr2O4 was observed on the surface of SUS430 after oxidation both in the air and fuel atmospheres of SOFC at 1073 K. The parabolic rate constant kp of the fuel side of the alloy is similar to that of the air side at 1073 K. The oxygen partial pressure dependence of the parabolic rate constant revealed that the growth rate of Cr2O3 scale is almost the same in various oxygen partial pressures ( PO 2g4.1 10 17 – 2.1 104 Pa) at 1073 K. Under the hydrogen potential gradient, the scale morphology of both the air side and the fuel side were almost similar to those which were oxidized in both atmospheres separately without hydrogen potential gradient. There is little effect of hydrogen potential gradient on oxidation behavior of Fe – Cr alloy in the atmosphere simulating SOFC.
Acknowledgements The authors thank Dr. Harumi Yokokawa, National Institute of Advanced Industrial Science and Technology, for valuable discussions of thermodynamics in spinel.
References [1] M. Dokiya, N. Sakai, N. Kawada, T. Yokokawa, T. Iwata, M. Mori, Intersoc. Energy Convers. Eng. 24 (1989) 1547. [2] J. Mizusaki, S. Yamauchi, K. Fueki, A. Ishikawa, Solid State Ion. 12 (1984) 119. [3] T.R. Armstrong, J.W. Stevenson, L.R. Pederson, P.E. Raney, J. Electrochem. Soc. 143 (1996) 2919. [4] I. Yasuda, M. Hishinuma, J. Electrochem. Soc. 143 (1996) 1583. [5] T. Kadowaki, T. Shiomitsu, E. Matsuda, H. Nakagawa, H. Tsuneizumi, T. Maruyama, Solid State Ion. 67 (1993) 65. [6] K. Huang, P.Y. Hou, J.B. Goodenough, Solid State Ion. 129 (2000) 237. [7] P. Kofstad, R. Bredesen, Solid State Ion. 52 (1992) 69. [8] K. Honegger, A. Plas, R. Diethelm, W. Glatz, in: H. Yokokawa, S.C. Singhal (Eds.), Solid Oxide Fuel Cells (SOFC), vol. VII, The Electrochemical Society, Pennington, NJ, 2001, p. 803, PV2001-16. [9] J.P. Abella´n, V. Shemet, F. Tietz, L. Singheiser, W.J. Quadakkers, A. Gil, in: H. Yokokawa, S.C. Singhal (Eds.), Solid Oxide Fuel Cells (SOFC), vol. VII, The Electrochemical Society, Pennington, NJ, 2001, p. 811, PV2001-16. [10] T. Horita, Y. Xiong, K. Yamaji, N. Sakai, H. Yokokawa, J. Electrochem. Soc. 150 (2003) A243. [11] T. Brylewski, M. Nanko, T. Maruyama, K. Przybylski, Solid State Ion. 143 (2001) 131. [12] K. Nakagawa, Y. Matsunaga, T. Yanagisawa, Mater. High Temp. 18 (2001) 51. [13] H. Kurokawa, Y. Oyama, K. Kawamura, T. Maruyama, J. Electrochem. Soc (submitted to). [14] C. Wagner, Z. Phys. Chem. 21 (1933) 25. [15] I. Barin, Thermochemical Data of Pure Substances, 3rd edition, VCH Verlagsgesellschaft, Weinheim, 1995. [16] Y. Jeannin, C. Mannerskantz, F.D. Richardson, Trans. Metall. Soc. AIME 227 (1963) 300.
H. Kurokawa et al. / Solid State Ionics 168 (2004) 13–21 [17] P.C. Lidster, H.B. Bell, Trans. Metall. Soc. AIME 245 (1969) 2273. [18] C.L. McCabe, R.G. Hudson, H.W. Paxton, Trans. Metall. Soc. AIME 212 (1958) 102. [19] W. Dresler, Trans. Inst. Min. Metall. 98 (1989) C61. [20] H. Yokokawa, (private communication). [21] B.-J. Lee, Metall. Trans., A, Phys. Metall. Mater. Sci. 24 (1993) 1919. [22] D.H. Speidel, Arnulf Muan, J. Am. Ceram. Soc. 46 (12) (1963) 578. [23] M.G.E. Cox, B. McEnanay, V.D. Scott, Phila. Mag. 26 (1972) 839. [24] R.K. Wild, Corros. Sci. 17 (1977) 87. [25] R.E. Lobnig, H.P. Schmidt, K. Hennesen, H.J. Grabke, Oxid. Met. 37 (1992) 81. [26] W.C. Hagel, A.W. Seybolt, J. Electrochem. Soc. 108 (1961) 1246.
21
[27] G.J. Yurek, Ann. Prog. Rept. (July 1, 1983 – June 30,1984), DOE No. DE-AC-02-79ER-10507, Dept of Energy, Gaithersburg, MD, Feb. 1985. [28] K. Przybylski, G.J. Yurek, Mat. Sci. Forum 43 (1989) 1. [29] E.W.A. Young, P.C.M. Stiphout, J.H.W. de Wit, J. Electrochem. Soc. 132 (4) (1989) 884. [30] A. Atkinson, R.I. Taylor, in: G. Simkovich, V.S. Stubican (Eds.), Transport in Nonstoichiometric Compounds, Nato ASI Series B, vol. 129, Plenum, New York, 1985, p. 285. [31] K. Hoshino, N.L. Peterson, J. Am. Ceram. Soc. 66 (1983) c202. [32] V.S. Stubican, L.R. Carinci, Z. Phys. Chem. 207 (1998) 215.