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Chromium vaporization from alumina-forming and aluminized alloys
Solid State Ionics 179 (2008) 2406–2415
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Solid State Ionics j o u r n a l h o m e p a g e : w w w. e l s ev i e r. c o m / l o c a t e / s s i
Chromium vaporization from alumina-forming and aluminized alloys M. Stanislowski, E. Wessel, T. Markus ⁎, L. Singheiser, W.J. Quadakkers Research Centre Jülich, Institute for Materials and Processes in Energy Systems (IEF-2), 52425 Jülich, Germany
a r t i c l e
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Article history: Received 15 March 2007 Received in revised form 2 September 2008 Accepted 11 September 2008 Keywords: Chromium vaporization Fe–Cr–Al alloys Aluminizing SOFC Degradation
a b s t r a c t The Cr vaporization from alumina-forming and aluminized alloys was investigated at temperatures between 800 °C and 1000 °C using the transpiration method. The growth of alumina scale was studied by XRD, SEM/ EDX, SNMS and TEM and discussed in the context of Cr vaporization. The experiments show that chromium vaporization from thermally grown alumina scales on alloys is about three orders of magnitude lower than that from pure chromia. The results indicate that Cr vaporization is governed by the content of dissolved Cr in the alumina scale. Bulk diffusion is identified as the major transport mechanism of Cr through the alumina scale. With regard to Cr vaporization as a major reason for long-term degradation in SOFCs the results show that alumina-forming-alloys and aluminized surfaces are promising for application in SOFC components without a current-conducting function. © 2008 Elsevier B.V. All rights reserved.
1. Introduction High-temperature alloys rely for their oxidation protection on the formation of a surface oxide scale of either Cr2O3, Al2O3 or SiO2 [1]. With respect to solid oxide fuel cell (SOFC) technology, the chromiaforming materials are being discussed for application in metalsupported SOFC designs and interconnector concepts because alumina and silica possess electronic conductivities which are far too small for most SOFC stack designs [2]. Alumina-forming alloys are being considered as candidate materials for heat exchangers, pumps, piping, casings and for high-temperature components without currentconducting function. In planar SOFC systems, Cr vaporization from metallic construction materials at high temperatures was identified as a major source of cell degradation. The gaseous Cr-species are reduced at the cathode to solid chromia and in this way inhibit the electrochemical processes [3–6]. In a previous paper the present authors reported on Cr vaporization from chromia- and spinel-forming steels [7]. This paper deals with alumina-forming alloys as well as aluminized chromiaforming alloys. Due to their high corrosion resistance and stability aluminaforming materials are used for a large number of applications such as catalyst supports in automotive applications or as furnace parts for temperatures up to 1300 °C [8]. Alumina-forming alloys usually contain about 20 wt.% chromium. [9]. The chromium is added because
⁎ Corresponding author. Tel.: +49 2461 614470; fax: +49 2461 613699. E-mail address: [email protected] (T. Markus). 0167-2738/$ – see front matter © 2008 Elsevier B.V. All rights reserved. doi:10.1016/j.ssi.2008.09.015
it promotes the formation of the protective alumina scale during the oxidation process. It is believed that chromium reduces the ingress of oxygen into the alloy thereby promoting oxidation of Al in form of an external α-Al2O3 scale instead of becoming oxidized internally. Chromia crystals further act as nucleation sites for α-alumina in the initial phase of the oxidation process [10]. Most investigations on alumina-forming materials address the oxidation behaviour at temperatures N1000 °C. Less publications deal with the oxidation behaviour in the temperature range between 800 °C and 950 °C which is of interest for SOFC applications [11–19]. Due to their high-Al contents, alumina-forming Fe- and Ni-based alloys possess poor room temperature ductility. Aluminizing improves the oxidation resistance of high-temperature alloys, for example in Nibase alloys for turbine-blades [20], without substantially affecting the mechanical properties of the core of the material. In the aluminizing process, the surface of Ni-, Fe-, Co- and Cr-based alloys is enriched with aluminium at temperatures between 700 °C and 1100 °C [21]. The diffusing Al forms intermetallic phases with the base material, e.g. βNiAl or FeAl [10,22–24]. At high temperatures the intermetallic phases form protective alumina scales. In this way, aluminizing offers the possibility of obtaining the protection of a thermally grown alumina scale to a variety of alloys without changing the mechanical properties in the core. In the literature, hardly any information is available on the vaporization of alumina-forming steels. Hilpert and Miller [25] determined the partial pressures and activities of Al, Cr and Fe/Ni in FeCrAl and NiCrAl alloys by Knudsen effusion mass spectrometry (KEMS). The vaporization of gaseous Al species from alumina was reviewed by Opila [26]. Chromium vaporization as well as the influence of metastable aluminas on the Cr vaporization from alumina-forming alloys has not yet been studied. However, in several
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Table 1 Chemical compositions of the investigated alloys determined by ICP-OES
Aluchrom YHf Aluchrom YB Kanthal AF PM 2000 Nicrofer 6025 HT Nicrofer 45 TM Ferrotherm 4828 Ducrolloy
Fe
Cr
Ni
Al
Si
Mn
Ti
Y
Zr
C
N
Others
74.7 74.2 74.5 76.5 9.3 23.9 65.8 5.5
20.2 20.5 21.2 18.6 25.0 27.3 19.2 92.9
0.18 0.14 0.19 0.02 62.6 45.9 11.3 –
5.7 5.2 5.5 5.7 2.3 0.1 – –
0.3 0.1 0.3 – 0.1 2.7 1.7 –
0.2 0.2 0.2 0.1 0.1 0.1 1.3 –
– 0.05 0.09 0.45 0.20 0.06 0.05 –
0.04 0.03 0.03 0.33 0.08 0.08 – 0.48
0.05 0.04 0.06 0.01 0.07 – – 0.02
0.024 0.022 0.017 0.014 0.17 0.053 0.048 0.010
0.005 0.005 0.014 0.006 0.104 – 0.046 0.014
Hf: 0.04 – Cu: 0.04 O: 0.254 – Ce: 0.05 P: 0.022 O: 0.449
technological applications alumina scales are regarded as protective coatings against Cr vaporization at high temperatures [27,28]. The aim of the present work was to measure the chromium release from alumina-forming alloys and aluminized surfaces by means of the transpiration method at temperatures between 800 °C and 1000 °C. The suitability of alumina scales or coatings for SOFC application concerning the retention of Cr will be discussed. 2. Experimental The following commercial alloys were investigated: the ferritic alumina-forming wrought alloys, Aluchrom YHf (material number 1.4767, ThyssenKrupp VDM, Werdohl, Germany), Aluchrom YB (material number 1.4767, ThyssenKrupp VDM, Werdohl, Germany), Kanthal AF (Kanthal AB, Hallstahammar, Sweden), the ferritic, alumina-forming oxide dispersion strengthened (ODS) alloy, PM 2000 (material number 1.4768, Plansee AG, Reutte, Austria), the high-Al Ni-based alumina-forming alloy, Nicrofer 6025 HT (material number 2.4633, ThyssenKrupp VDM, Werdohl, Germany). Additionally, two chromia-forming high-temperature alloys were used to test the suitability of an aluminizing treatment: the Ni-based alloy, Nicrofer 45 TM (material number 2.4889, ThyssenKrupp VDM, Werdohl, Germany) and the austenitic Fe-based alloy, Ferrotherm 4828 (material number 1.4828, Thyssen Schulte GmbH, Dortmund, Germany). As reference material for all the studies the pure chromiaforming Cr-based ODS alloy, Ducrolloy (Cr5Fe1Y2O3, Plansee AG, Reutte, Austria) was incorporated in the test programme. The chemical compositions of these alloys, determined by optical emission spectroscopy using an inductively coupled plasma (ICP-OES, IRISadvantage, Thermo Jarrell Ash, Franklin, USA), are listed in Table 1. For the vaporization experiments, samples with the dimensions 80 × 20 mm2 and a thickness of 2–3 mm were cut from plates by water jet cutting. The edges of the samples were rounded by spark erosion in order to avoid preferred sites for oxide scale cracking and/or spallation during the vaporization experiments. Details of the geometry of the samples are given in [7]. The samples were ground with SiC paper up to 1200 grit and ultrasonically cleaned in ethanol and acetone. Preoxidations were performed in a muffle furnace at 1200 °C for 24 h in air. In order to investigate Cr retention by aluminide coatings, some samples of alloys Nicrofer 45 TM and Ferrotherm 4828 were ground with SiC paper up to 1200 grit, ultrasonically cleaned in ethanol and acetone and aluminized using chemical vapor deposition (CVD) at 1080 °C by MTU Aero Engines GmbH, Munich, Germany [29]. The transpiration method was used to measure the chromium vaporization rates. The method is based on the principle that the vapor over the heated sample is carried away by a constant gas flow and collected in a condenser. The transported mass was subsequently determined by quantitative chemical analysis. The transport rate at a given temperature was calculated from the transported mass of Cr, the surface area of the sample and the duration of the experiment. The influence of the gas flow rate was eliminated in this work by setting the gas flow to rates sufficiently high to ensure the transported mass
of Cr to depend solely on the kinetics of the vaporization reaction at the surface of the sample (see Refs. [7,30]). The experiments were carried out with a constant flow of humidified air at temperatures between 800 °C and 1000 °C. The air flow was controlled by a flow meter, type 5850TR, supplied by Brooks, Veenendaal, The Netherlands. For the experiments the air flow rate was set to 1500 ml/min, relative to standard conditions (273 K, 101,325 Pa). The humidity of the air was adjusted by a bubble humidifier in combination with a condenser. The temperature of the condenser was controlled by a thermostat, type C10/K15, supplied by ThermoHaake, Karlsruhe, Germany. For the experiments, a humidity of p(H2O) = 1.88 · 103 Pa was adjusted, which corresponds to a relative humidity of 60% at 25 °C and standard air pressure. After the experiments, the Cr condensate was dissolved in HCl (Merck, Suprapur®) and quantitatively analysed by inductively coupled plasma mass spectrometry, ICP-MS (PerkinElmer, Norwalk, USA). Details of the experimental set-up and the used analysis method are given in Refs. [7,31]. After the vaporization experiments, the samples were characterized by X-ray diffraction (XRD, X'Pert MRD, Philips, The Netherlands, Cu Kα radiation) and scanning electron microscopy (SEM, LEO 440 and LEO 1530-Gemini, Cambridge, UK) coupled with energy dispersive Xray analysis (EDX, ISIS 300, Eynsham, UK). The concentration of the elements Al, Fe and Cr in the oxide scale was investigated as a function of the depth by electron-collision sputtered neutrals mass spectrometry (SNMS, Simslab 410a, Fisons, East Grinstead, UK). The surface was sputtered with O+2 ions at an energy of 10 keV and a current of 250 nA. A grid pattern of 420 × 420 μm2 and a gate of 170 × 170 μm2 were adjusted. The current of the ions 27Al+, 56Fe+ and 52Cr+ was determined as a function of the
Fig. 1. Comparison of the chromium vaporization rates of different chromia- and alumina-forming alloys at 900 °C in humid air after preoxidation for 24 h at 1200 °C in air. Ducrolloy was preoxidized at 900 °C for 100 h in air.
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Fig. 2. SEM images of cross section and plan view of a), b), Aluchrom YHf, and c), d), PM 2000, after oxidation at 1200 °C for 24 h in air.
sputtering time. An isotope correction was performed with respect to the frequency of the analysed isotopes. 3. Results 3.1. Alumina- and chromia-forming alloys 3.1.1. Comparison of Cr vaporization from different alloys Samples of the Aluchrom YHf, Aluchrom YB, Kanthal AF, PM 2000 and Nicrofer 6025 HT alloys were preoxidized at 1200 °C for 24 h in air to obtain a dense α-Al2O3 layer. The oxidized surfaces of the samples of the first four alloys looked bright grey whereas that of Nicrofer 6025 HT looked dark green with grey speckles. The Cr-based alloy Ducrolloy was preoxidized at 900 °C for 100 h in air. After oxidation it had a dark green colour. The chromium vaporization of the preoxidized samples was measured using the transpiration apparatus at 800 °C in air with a humidity of 1.88%. In order to obtain a condensate with a sufficiently high Cr content for the chemical analyses, the duration of the
measurements was 300 h for the ferritic alumina formers Aluchrom YHf, Aluchrom YB, Kanthal AF and PM 2000, and 40 h for Nicrofer 6025 HT and 33 h for Ducrolloy. Each sample was measured twice in succession. For the preoxidized samples no time-dependent vaporization behaviour was observed. The results of the vaporization experiments are shown in Fig. 1. After the transpiration tests XRD and SEM/EDX revealed the oxide scales on Aluchrom YHf, Aluchrom YB, Kanthal AF and PM 2000 to consist of α-Al2O3, whereas the oxide scale on Ducrolloy consisted of pure Cr2O3. Nicrofer 6025 HT formed a (Cr,Fe)2O3 scale with locally a 0.3 μm thick top layer of Cr–Ti-oxides, presumably (Cr,Mn)2Ti3O9. The Al in Nicrofer 6025 HT was internally oxidized to α-Al2O3. SEM images of Aluchrom YHf and PM 2000 after preoxidation at 1200 °C for 24 h in air are shown in Fig. 2. Fig. 3 shows SEM images of Nicrofer 6025 HT after oxidation at 1200 °C for 24 h in air. The alumina scales on Aluchrom YHf, Aluchrom YB, Kanthal AF and PM 2000 had thicknesses of 7.95 μm, 7.67 μm, 7.18 μm and 5.36 μm, respectively. Small surface nodules were visible in the backscattered
Fig. 3. SEM images of a) cross section and b) plan view of Nicrofer 6025 HT after oxidation at 1200 °C for 24 h in air.
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Fig. 4. Chromium vaporization rates as a function of time for non-preoxidized Aluchrom YHf, Nicrofer 6025 HT and Ducrolloy at 800 °C in humid air.
Fig. 6. Chromium vaporization rates as a function of time for Aluchrom YHf at 800 °C, 900 °C and 1000 °C in humid air, without preoxidation. The data points were fitted to a power law time dependence of the form y = a ∙ t b.
electron images of Aluchrom YHf and PM 2000 (see Fig. 2b and d). In the case of Aluchrom YHf they contained Zr and Y with impurities of Cr and in the case of PM 2000 of Y and Ti with minor amounts of Ca and Fe. Mg and Cr were found in the outer part of the alumina based scale on Aluchrom YHf. Mg as well as Ca is used for deoxidation in the fabrication of steels. Trace amounts of these elements may therefore be still present in the finished product. The Cr vaporization of non-preoxidized samples of Ducrolloy, Nicrofer 6025 HT and Aluchrom YHf was investigated as a function of time in order to correlate the Cr vaporization with the progress of growth of the oxide scale. Fig. 4 shows the Cr vaporization rates at 800 °C in humid air during oxidation tests up to 500 h. Fig. 5 shows SEM images of the alumina scales of Aluchrom YHf and Nicrofer 6025 HT after exposure at 800 °C for 10 h and 500 h, respectively. After oxidation for 10 h at 800 °C in air the alumina scale on Aluchrom YHf had a thickness of about 0.3 μm and consisted of α- and γ-alumina, according to XRD. The alumina scale on Nicrofer 6025 HT had a thickness of about 0.2 μm after oxidation for 500 h at 800 °C in air and consisted of α-alumina.
10 h and 1000 h at 900 °C and 1000 °C in humid air. The XRD patterns of Aluchrom YHf after oxidation at 800 °C for 10 h, 50 h, 150 h, and 1000 h in humid air are shown in Fig. 7. SEM images of the oxide scales formed on Aluchrom YHf after oxidation for 1000 h in humid air at 800 °C, 900 °C and 1000 °C are shown in Fig. 8. EDX line scans of the elements O, Al, Cr, Fe and Mg measured in metallographic cross sections after oxidation at 1000 °C for 1000 h in humid air are shown in Fig. 9. Mg is found in the outer part of the Al scale (Fig. 9). Fig. 10a–c shows SNMS depth profiles of the elements Fe, Al and Cr in the oxide scale on Aluchrom YHf after oxidation at 900 °C in air for 10 h, 100 h and 1000 h, respectively. Fig. 10a–c shows that there is a strong gradient in the Cr concentration in the alumina scale after exposures of 10 h and 100 h at 900 °C that disappears after 1000 h at 900 °C.
3.1.2. Temperature dependence of Cr vaporization on alumina-forming steels The Cr vaporization rates from non-preoxidized Aluchrom YHf were measured in humid air as a function of time for the temperatures 800 °C, 900 °C and 1000 °C. The results are shown in Fig. 6. On Aluchrom YHf, the phase α-Al2O3 and a spinel structure, probably MgAl2O4 were detected by XRD after oxidation for between
3.2. Aluminized alloys The Cr vaporization rates of aluminized samples of the Nicrofer 45 TM and Ferrotherm 4828 alloys were measured at 800 °C in humid air as a function of time without preoxidation. For comparison, the same measurements were carried out with non-aluminized samples of the same alloys. The results are shown in Fig. 11. According to XRD and SEM, the oxide scale on Nicrofer 45 TM consisted of a chromia scale on top of a thin silica layer. Ferrotherm 4828 formed a chromia scale with an outer (Cr,Mn)-spinel layer which tended to spall off during the vaporization experiments. The aluminized samples formed well-
Fig. 5. SEM images of a) Aluchrom YHf and b) Nicrofer 6025 HT, after oxidation at 800 °C in air for 10 h and 500 h, respectively.
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Fig. 7. XRD patterns of Aluchrom YHf after oxidation at 800 °C in air for 10 h, 50 h, 150 h, and 1000 h.
adherent, dark grey alumina scales already after very short exposure times. The aluminized surfaces were investigated by XRD and SEM/EDX. Images of the microstructure of the aluminized samples are shown in Fig. 12. During the aluminizing process an Al-enriched zone was formed near the surface of Nicrofer 45 TM with a thickness of about 15 µm consisting mainly of β-NiAl and FeAl. Beneath the Al enrichment zone an interdiffusion zone with a width of about 10 µm was identified by SEM. The surface of the aluminized Nicrofer 45 TM was decorated with small precipitates, especially near the grain boundaries. EDX analyses showed that these inclusions consisted of Cr- and Fe-rich phases. After oxidation for 100 h at 800 °C in humid air, the formation of a surface scale of α- and γ-alumina with a thickness of 1.2–1.8 μm was observed. The Al-enriched zone in Ferrotherm 4828 consisted mainly of NiAl, FeAl and other Fe aluminides and had a width of about 20 μm. A deep interdiffusion zone with a width of about 150 μm was present beneath the aluminized zone. After oxidation at 800 °C for 100 h in humid air a surface scale consisting of α- and γ-alumina was observed (Fig. 12).
Fig. 8. SEM cross sections and plan views of the oxide scale of Aluchrom YHf after oxidation for 1000 h in air at a), b), 800 °C, c), d), 900 °C, and e), f), 1000 °C.
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4. Discussion Of the preoxidized samples in Fig. 1, the highest Cr vaporization rate is, as expected, observed for Ducrolloy (Cr5Fe1Y2O3), which forms a surface scale of virtually pure chromia. The Cr vaporization rate for the Ni-based alloy, Nicrofer 6025 HT, preoxidized for 24 h at 1200 °C in air, is similar to that of Ducrolloy. This can be explained by the formation of a chromium oxide scale on top of the alloy (see Fig. 3a). The aluminium present in this alloy was oxidized internally and did not form an external alumina scale. The surface of the Cr oxide was covered locally with a 200–300 nm thick layer of (Cr, Mn)2Ti3O9 on top of the iron/chromium oxide scale. This Ti–Cr-rich layer may reduce Cr vaporization but does not form an effective barrier layer. After preoxidation for 24 h at 1200 °C in air the chromium vaporization rates of the alumina formers Aluchrom YHf, Aluchrom YB, Kanthal AF and PM 2000 are about three orders of magnitude lower
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than that of Ducrolloy. These alloys are all of the Fecralloy type with the nominal composition (in mass%) Fe20Cr5Al + (Y, Zr, Hf) and form dense surface scales of α-Al2O3. Only minor differences in the Cr vaporization rates within this group of alloys were found. The alumina scale on the ODS alloy PM 2000 is significantly thinner and contains fewer pores than that of the wrought alloys Aluchrom YHf, Aluchrom YB and Kanthal AF. According to the literature, the growth of the reactive element (RE) doped alumina scale on the alloy is mainly governed by inward oxygen transport via grain boundaries [32]. Minor contents of reactive elements, such as Y, Zr, Ce, La, and Hf [33] block the grain boundaries and suppress the outward diffusion of cations [32,34]. PM 2000 contains the highest RE content of the investigated Fecralloy type materials (see Table 1). This may explain the small thickness of the alumina scale. Another possible explanation for the differences in oxidation rate might be the shape in which the RE prevails in the alloy (in metallic form or in oxide dispersion), the amount of RE and the type
Fig. 9. EDX line scans of the elements O, Al, Cr, Fe, and Mg through the alumina scale of Aluchrom YHf after annealing at 1000 °C for 1000 h in air.
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that of Aluchrom YHf, Kanthal AF or Aluchrom YB. In this way, a smaller surface area is exposed to the environment, which leads to lower vaporization rates compared to samples with rough oxide surfaces. The RE content leads to the formation of RE-rich nodules on the surface of the alumina scale of PM 2000 and Aluchrom YHf (Fig. 2b and d). The formation of such surface nodules is also known from the literature [13,32]. Their formation and growth is attributed to the outward transport of REs and minor alloying additions (e.g. Mg, Ti) via grain boundaries [34]. EDX analyses showed that these nodules contained Y and Ti with small amounts of Ca and Mg, depending on the type of alloy. Small concentrations of Cr were detected in the nodules and also in alumina grains at the surface. The presence of Cr in nodules on the surface of alumina-forming alloys was also reported in the literature [9,37,38]. A high concentration of Mg was detected in the outer part of the alumina scale of Aluchrom YHf, probably in the form of Al–Mg spinel, which corresponds to the observations by Dimyati et al. [39]. The presence of Cr in the nodules as well as on the surface of the alumina grains and the small difference in Cr vaporization between PM 2000 and Aluchrom YHf, Kanthal AF and Aluchrom YB, indicates that the transport of Cr by grain boundaries is not the dominating factor for Cr vaporization. The transpiration experiments with non-preoxidized samples in Fig. 4 show the dependence of the Cr vaporization rate on the initial growth of the oxide scale. The Cr vaporization rates of Ducrolloy in Fig. 4 decrease within the first 50 h. This can be explained by the fact that initially the formation of Cr(g) contributes to the vaporization process whereas after a dense scale has formed, the overall vaporization process is governed by vaporization processes at the scale/gas interface [40]. After the surface is completely covered with a dense chromia scale the vaporization rate is approximately constant. In contrast to this, the Cr vaporization rate of Nicrofer 6025 HT decreases steadily with increasing time. After 500 h of oxidation at 800 °C in air the Cr vaporization rates are reduced by 90% compared to the pure chromia scale of Ducrolloy. This can be explained by the formation of a slowly-growing alumina scale which reaches a thickness of 200–300 nm after oxidation for 500 h at 800 °C in air (see Fig. 5b). After oxidation of the same alloy for 24 h at 1200 °C in air the aluminium was oxidized internally (see Fig. 3). This different oxidation behaviour illustrates that the Nicrofer 6025 HT alloy cannot be considered as a “real” alumina-forming alloy. The formation of the alumina scale on Ni-base alloys with low aluminium contents, such as Nicrofer 6025 HT, depends strongly on the oxidation conditions [1]. The thickness of the alumina scale on Nicrofer 6025 HT after oxidation for 500 h at 800 °C in air is comparable to the thickness of the alumina scale on Aluchrom YHf after 10 h at 800 °C in air (see Fig. 5a).
Fig. 10. SNMS depth profiles of the elements Fe, Al and Cr in the oxide scale on Aluchrom YHf after oxidation at 900 °C in humid air for a), 10 h, b), 100 h and c), 1000 h.
of RE [35]. Also inclusions of REs such as Hf and especially Zr, which prevail e.g. in Aluchrom YHf, tend to induce microporosity in the alumina scale [36]. This might be an additional rapid pathway for Cr from the alloy to the oxide surface. Apart from the thinner, denser oxide scale the slightly lower Cr vaporization rates of PM 2000 may be affected by the smoother morphology of the oxide surface compared to
Fig. 11. Chromium vaporization rates as a function of time at 800 °C in humid air for non-preoxidized and aluminized samples of Nicrofer 45 TM and Ferrotherm 4828.
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Fig. 12. SEM cross sections and plan views of the oxide scale of Nicrofer 45 TM, a), b), after aluminizing, c), d), after aluminizing and annealing at 800 °C for 100 h in air, and, e), f), of Ferrotherm 4828 after aluminizing and annealing at 800 °C for 100 h in air.
The vaporization rates of these two alloys were very similar after these oxidation times (see Fig. 4). After 100 h of oxidation at 800 °C in air the alumina scale on Aluchrom YHf reached a thickness of 0.6–0.9 μm and the Cr vaporization rates were more than two orders of magnitude lower than that of Nicrofer 6025 HT after 500 h at 800 °C. The Cr vaporization rates of Aluchrom YHf were more than 3 orders of magnitude lower than those of Ducrolloy. Similar vaporization rates were measured for the preoxidized samples (see Fig. 1). This demonstrates the high Cr retention capability of alumina. The types of alumina phases formed after oxidation at 800 °C were identified by XRD as α-alumina, in the case of Nicrofer 6025 HT, and as a mixture of α- and γ-alumina, in the case of Aluchrom YHf. In the latter case, the alumina scale consisted of a layered structure consisting of an outer layer of metastable γ-alumina, which formed characteristic platelets, and an inner layer of stable α-alumina (see Figs. 5a and 8a). Similar results were reported by other authors [10,37,41]. The formation of the layered structure can be explained by the growth of metastable aluminas by outward cation diffusion [16,42] and the growth of α-alumina by inward oxygen diffusion [43].
After 150 h of oxidation at 800 °C in air the metastable θ-alumina phase was detected on Aluchrom YHf by XRD (see Fig. 7), next to γand α-alumina. In the literature several papers are available dealing with the metastable alumina phases formed at lower temperatures. These range from θ-alumina [12,13,15,32,44] to γ-alumina [16,45] and combinations of the metastable phases of γ-, δ- and θ-alumina [19]. In the present work, it was found that the θ-, γ- and α-alumina phases co-exist after oxidation at 800 °C for 1000 h in air and that the amount of metastable phases increased with time (see Fig. 7). It is generally assumed that metastable aluminas are formed at the beginning of the oxidation process and are replaced by stable α-alumina at times ranging from minutes to hours depending on the temperature [12,23]. The sequence of the phase transformation ranges is given in the literature beginning with metastable γ-alumina to δ- and θ-alumina to finally stable α-alumina [23,46]. The formation and morphology of the metastable aluminas are influenced by several factors such as small deviations in the chemical composition [15], pre-existing oxides [10], presence of water vapor [47], surface defects such as grain boundaries [16] or previous cold deformation [32].
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Fig. 6 shows the influence of temperature on the Cr vaporization rate of non-preoxidized Aluchrom YHf. The oxide scales formed at 900 °C and 1000 °C consisted exclusively of α-alumina. Even after relatively short oxidation times of 10 h at 900 °C and 1000 °C no metastable alumina phases could be detected by XRD. This is in contrast to the oxidation test at 800 °C where the oxide scale always consisted of γ-, θ-, and α-alumina. Accordingly, there must be a transition temperature between 800 °C and 900 °C above in which only α-alumina is formed. In the literature, this transition temperature is given as 900 °C [13], 950 °C [48] or 1050 °C [49]. This transition temperature will, however, depend on exposure time. The Cr vaporization rates at 1000 °C are more than one order of magnitude higher than at 800 °C and a factor of 3–4 higher than at 900 °C. The time dependence of the Cr vaporization rates at 1000 °C and 900 °C are comparable but differ from that at 800 °C. At 800 °C the Cr vaporization rates show a stronger decrease with increasing time. During the initial stage of oxidation the Cr vaporization strongly depends on the growth of the alumina scale. The time dependence of the Cr vaporization rates was fitted to a power law time dependence of the form y = a ∙ tb. The exponents of the best fits to the curves at 900 °C and 1000 °C were b = −0.67 and b = −0.69, respectively. For the curve of the Cr vaporization rates at 800 °C an exponent of b = −1.01 was obtained for the best fit (see Fig. 6). These results indicate that the kinetics of the Cr vaporization at 900 °C and 1000 °C are similar but differ from those at 800 °C. This is in agreement with the observations of other authors who measured a fundamental change in oxidation kinetics with the temperature increasing from 800 °C to 900 °C [12,50]. In the literature, this effect was attributed to the formation of metastable aluminas at lower temperatures [16], sometimes accompanied by crack formation caused by the transformation of metastable to stable alumina [41]. After exposures of 1000 h at 1000 °C, 900 °C and 800 °C in air the SEM images of the cross sections of the alumina scales in Fig. 8 showed alumina scales with thicknesses of 3.5–4.2 μm, 1.4–2.6 μm and 0.9– 2.6 μm, respectively. It is noteworthy that the alumina scales formed during oxidation at 800 °C and 900 °C within 1000 h have similar thicknesses. This can be explained by the formation of different types of alumina phases. It is well known from the literature that metastable aluminas grow faster than α-alumina [12,13,22]. Metastable aluminas are cation-deficient and are therefore considered to possess a lower Cr retention capability than α-alumina [27,28]. This cannot be confirmed by transpiration experiments performed in the work. This might be explained by the formation of a dense α-alumina layer under the metastable γ-/θ-alumina scale, as already mentioned. In Fig. 6, the Cr vaporization rates measured at 1000 °C after the same exposure time are shown to be a factor 4–5 higher than those at 900 °C. The alumina scale formed on Aluchrom YHf after an oxidation of 1000 h at 1000 °C is about twice as thick as that which was formed at 900 °C (see Fig. 8). Vaporization tests [51] with preoxidized samples of Aluchrom YHf and PM 2000 at 800 °C, 900 °C and 1000 °C showed that the Cr vaporization rates follow an Arrhenius-type temperature dependence. According to the Hertz–Langmuir-equation the vaporization rates are proportional to the partial pressures. In our experiments we did not directly measure the partial pressures but rather the quantities of vaporized elements per time and area. As the most abundant volatile Crspecies are of the one-atomic form CrO2(OH)2, CrO2(OH) and CrO3 the measured vaporization rates are directly proportional to the partial pressures. Therefore one can follow in this context: 1 pðCrÞ e exp − T
ð1Þ
This type of temperature dependence is valid for the vaporization rates as well as for the diffusion. The question remains of how the Cr content in the alumina scale is influenced by oxide growth. It is known that Cr2O3 and Fe2O3 are
formed in the initial stage of oxidation on the surface of ferritic alumina-forming alloys, before the dense Al2O3 layer is formed [10,11,14,15]. After a few minutes or hours (depending on temperature) the transient oxides are replaced by Al2O3. It is proposed that Cr2O3 crystals act as nucleation sites for α-Al2O3 [10]. The initially high Cr concentration in the oxide scale thus drops significantly within the first few hours of oxidation as was shown by Quadakkers et al. [11,52]. The SNMS measurements after oxidation at 900 °C in air for 10 h and 100 h and 1000 h in Fig. 10a–c shows that a gradient of Cr exists within the alumina scale. With increasing oxidation time the gradient flattens and after 1000 h of oxidation at 900 °C the concentration of Cr at the surface is larger by about a factor of 2 than after 10 h or 100 h. This will lead to Cr vaporization rates that are higher by a factor of 2 after long oxidation times. The decrease of the Cr gradient in the alumina scale after 1000 h of oxidation can be explained by the decrease of the scale thickening rate of the alumina scale with increasing time and thickness of the scale. The concentration gradient of Cr is then flattened by the slower diffusion transport. The transpiration experiments with aluminized samples in Fig. 11 show that surface modification by aluminizing offers a similar protection against Cr vaporization of chromia-forming alloys as found for the alumina-forming steels. The measured Cr vaporization rates of non-preoxidized samples of the aluminized Ni-base alloy Nicrofer 45 TM and the aluminized austenitic steel Ferrotherm 4828 are comparable to that of Aluchrom YHf. After annealing at 800 °C for 100 h in air the alumina scales grown on the aluminized samples have a thickness of 1.2–1.8 μm, which is about twice the thickness of the alumina scale on Aluchrom YHf after the same oxidation time (see Fig. 12). Grabke [24] and Prasanna et al. [13] reported that minor additions of Y suppress the formation of θ-alumina and encourage the formation of α-alumina. In contrast to Aluchrom YHf, the aluminized samples contain no Y. The faster growth of the alumina scale on the aluminized alloys can therefore be explained by the slower phase transformation from rapid-growing metastable γ- and θ-aluminas to slowly-growing stable α-alumina [12,13,22]. Another explanation might be, that the rate of nucleation of α-Al2O3 depends on the underlying metallic matrix [53]. For aluminized Ferrotherm 4828 it was observed that the morphology of the alumina scale formed on the alloy grains differed from that near the alloy grain boundaries (see Fig. 12). This can be attributed to a slower diffusion of aluminium in the austenitic grains compared to the diffusion in the ferritic grains and along the grain boundaries. The more rapid aluminium diffusion in the ferritic grains is also responsible for the deeper ingress of aluminium in Ferrotherm 4828 than in Nicrofer 45 TM. The narrow Al-diffusion zone of 30–200 μm of the aluminized alloys is expected to not substantially affect the mechanical properties of the bulk of the alloy. Due to the presence of the aluminium reservoir in the surface-near regions the alumina scales on aluminized alloys grow and are self-healing and are therefore more durable and reliable than e.g. ceramic coatings. Aluminizing thus offers an effective and attractive measure to obtain protection against Cr vaporization by a thermally grown and self-healing oxide scale. 5. Conclusions The measurement of Cr vaporization from alumina scales on alumina-forming and aluminized alloys by the transpiration method shows that Cr vaporization rates from alumina scales are about three orders of magnitude lower than those from pure chromia scales. It can be assumed that the kinetics of the Cr release is governed by the scale growth. At 1000 °C and 900 °C the kinetics of the scale growth are similar. At 800 °C a different scale growth kinetics could be identified from the results of Cr transpiration experiments which can be explained by the formation of metastable γ- and θ-aluminas. The chromium retention of metastable aluminas was comparable to that of stable α-alumina.
M. Stanislowski et al. / Solid State Ionics 179 (2008) 2406–2415
The Cr retention of aluminized surfaces of Ni- and Fe-base alloys was proven to be as efficient as that of alumina-forming steels. Thus, aluminizing offers the possibility to reduce chromium vaporization from high-temperature alloys by about two to three orders of magnitude compared to chromia-forming materials. This makes aluminizing a very interesting method for the protection of metal components in SOFC systems which do not have a current-conducting function. Acknowledgements The authors are grateful to their colleagues at Research Centre Jülich for their assistance with the work. Special thanks are due to Mr. H. J. Penkalla and Ms. D. Esser (IWV-2) for TEM/FIB studies, Mr. U. Breuer (ZCH) for SNMS measurements and Mr. P. Lersch (IWV-2) for XRD measurements. ThyssenKrupp VDM, Werdohl (Germany), Plansee AG, Reutte (Austria) and Kanthal AB, Hallstahammar (Sweden) are acknowledged for providing the materials. A special acknowledgment is given to Klaus Hilpert who suddenly and unexpectedly died in retirement. We acknowledge him for his valuable contributions to this paper. References [1] W.J. Quadakkers, J. Piron-Abellàn, V. Shemet, L. Singheiser, Mater. High Temp. 20 (2003) 115. [2] W.A. Meulenberg, A. Gil, E. Wessel, H.P. Buchkremer, D. Stöver, Oxid. Metals 57 (2002) 1. [3] D. Das, M. Miller, H. Nickel, K. Hilpert, in: U. Bossel (Ed.), Proceedings of the 1st European Solid Oxide Fuel Cell Forum, Lucerne, Switzerland, October 3–7, (1994) 703 (EPFL). [4] S. Taniguchi, M. Kadowaki, H. Kawamura, T. Yasuo, Y. Akiyama, Y. Miyake, T. Saitoh, J. Power Sources 55 (1995) 73. [5] K. Hilpert, D. Das, M. Miller, D.H. Peck, R. Weiß, J. Electrochem. Soc. 143 (1996) 3642. [6] Y. Matsuzaki, I. Yasuda, Solid State Ionics 132 (2000) 271. [7] M. Stanislowski, E. Wessel, K. Hilpert, T. Markus, L. Singheiser, J. Electrochem. Soc. 154 (2007) A295. [8] I. Itoh, M. Fukuya, R. Hisatomi, H. Morimoto, K. Ohmura, H. Tanaka, F. Fudanki, M. Arakawa, Nippon Steel Tech. Rep. 64 (1995) 69. [9] G.C. Wood, Oxid. Metals 2 (1970) 11. [10] M.W. Brumm, H.J. Grabke, Corr. Sci. 33 (1992) 1677. [11] W.J. Quadakkers, A. Elschner, W. Speier, H. Nickel, Appl. Surf. Sci. 52 (1991) 271. [12] W.J. Quadakkers, K. Schmidt, H. Grübmeier, E. Wallura, Mater. High Temp. 10 (1992) 23. [13] K.M.N. Prasanna, A.S. Khanna, R. Chandra, W.J. Quadakkers, Oxid. Metals 46 (1996) 465.
2415
[14] J. Jedliński, A. Glazkov, M. Konopka, G. Borchardt, E. Tscherkasova, M. Bronfin, M. Nocuń, Appl. Surf. Sci. 103 (1996) 205. [15] D. Naumenko, W.J. Quadakkers, A. Galerie, Y. Wouters, S. Jourdain, Mater. High Temp. 20 (2003) 287. [16] H. El Kadiri, R. Molins, Y. Bienvenu, M.F. Horstemeyer, Oxid. Metals 64 (2005) 63. [17] R. Molins, A.M. Huntz, Mater. Sci. Forum 461–464 (2004) 29. [18] H. El Kadiri, R. Molins, Y. Bienvenu, Mater. Sci. Forum 461–464 (2004) 1107. [19] G. Barthomé, E. N'Dah, Y. Wouters, A. Galerie, Mater. Corr. 56 (2005) 389. [20] R. Streiff, J. Phys. IV, C9 3 (1993) 17. [21] H. Simon, M. Thoma, Angewandte Oberflächentechnik für metallische Werkstoffe, Hanser Verlag, München, 1985. [22] G.C. Rybicki, J.L. Smialek, Oxid. Metals 31 (1989) 275. [23] J. Doychak, J.L. Smialek, T.E. Mitchell, Met. Trans. A 20A (1989) 499. [24] H.J. Grabke, Mater. Sci. Forum 251–254 (1997) 149. [25] K. Hilpert, M. Miller, Z. Metallk. 83 (1992) 739. [26] E.J. Opila, Mater. Sci. Forum 461–464 (2004) 765. [27] W.J. Quadakkers, F. Baumanns and H. Nickel, German Patent No. DE 44 10 711 C1 (1994). [28] W.J. Quadakkers, German Patent No. DE 195 47 699 A1 (1995). [29] H. Pillhöfer, M. Thoma, H. Walter and P. Adam, German Patent No. DE 41 19 967 C1 (1991). [30] H.K. Hofmeister, R.v. Haeseler, O. Glemser, Z. Elektrochem. 64 (1960) 513. [31] M. Stanislowski, U. Seeling, D.-H. Peck, S.-K. Woo, L. Singheiser, K. Hilpert, Solid State Ionics 176 (2005) 2523. [32] W.J. Quadakkers, Werkst. Korr. 41 (1990) 659. [33] W.J. Quadakkers, L. Singheiser, Mater. Sci. Forum 369–372 (2001) 77. [34] W.J. Quadakkers, H. Holzbrecher, K.G. Briefs, H. Beske, Oxid. Met. 32 (1989) 67. [35] B. Pint, J. Am. Ceram. Soc. 86 (2003) 686. [36] E. Wessel, V. Kochubey, D. Naumenko, L. Niewolak, L. Singheiser, W.J. Quadakkers, Scripta Mat. 51 (2004) 987. [37] P.T. Moseley, K.R. Hyde, B.A. Bellamy, G. Tappin, Corr. Sci. 24 (1984) 547. [38] M.J. Bennett, R. Newton, J.R. Nicholls, Mater. High Temp. 20 (2003) 347. [39] A. Dimyati, H.J. Penkalla, P. Untoro, D. Naumenko, W.J. Quadakkers, J. Mayer, Z. Metallkd. 94 (2003) 180. [40] P. Kofstad, J. Phys. Chem. Solids 44 (1983) 129. [41] W.J. Quadakkers, D. Naumenko, E. Wessel, V. Kochubey, L. Singheiser, Oxid. Metals 61 (2004) 17. [42] B.A. Pint, J.R. Martin, L.W. Hobbs, Solid State Ionics 78 (1995) 99. [43] J. Jedlinski, R. Borchardt, Oxid. Metals 36 (1991) 318. [44] W.C. Hagel, Corrosion 21 (1965) 316. [45] L. Marechal, B. Lesage, A.-M. Huntz, R. Molins, Mater. High Temp. 20 (2003) 295. [46] K. Wefers, G.M. Bell, Techn. Paper, vol. 19, Alcoa Res. Lab., 1972. [47] J.R. Regina, J.N. DuPont, A.R. Marder, Oxid. Metals 61 (2004) 69. [48] J.A. Nychka, D.R. Clarke, Oxid. Metals 63 (2005) 325. [49] H. Al-Badairy, D. Naumenko, J. Le Coze, G.J. Tatlock, W.J. Quadakkers, Mater. High Temp. 20 (2003) 405. [50] M.J. Bennett, R. Newton, J.R. Nicholls, H. Al-Badairy, G.J. Tatlock, Mater. Sci. Forum 461–464 (2004) 463. [51] M. Stanislowski, PhD thesis, RWTH Aachen University, Germany (2006). [52] W.J. Quadakkers, A. Elschner, H. Holzbrecher, K. Schmidt, W. Speier, H. Nickel, Mikrochim. Acta 107 (1992) 197. [53] D. Clemens, V. Vosberg, L.W. Hobbs, U. Breuer, W.J. Quadakkers, H. Nickel, J. Anal. Chem. 355 (1996) 703.
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Solid State Ionics j o u r n a l h o m e p a g e : w w w. e l s ev i e r. c o m / l o c a t e / s s i
Chromium vaporization from alumina-forming and aluminized alloys M. Stanislowski, E. Wessel, T. Markus ⁎, L. Singheiser, W.J. Quadakkers Research Centre Jülich, Institute for Materials and Processes in Energy Systems (IEF-2), 52425 Jülich, Germany
a r t i c l e
i n f o
Article history: Received 15 March 2007 Received in revised form 2 September 2008 Accepted 11 September 2008 Keywords: Chromium vaporization Fe–Cr–Al alloys Aluminizing SOFC Degradation
a b s t r a c t The Cr vaporization from alumina-forming and aluminized alloys was investigated at temperatures between 800 °C and 1000 °C using the transpiration method. The growth of alumina scale was studied by XRD, SEM/ EDX, SNMS and TEM and discussed in the context of Cr vaporization. The experiments show that chromium vaporization from thermally grown alumina scales on alloys is about three orders of magnitude lower than that from pure chromia. The results indicate that Cr vaporization is governed by the content of dissolved Cr in the alumina scale. Bulk diffusion is identified as the major transport mechanism of Cr through the alumina scale. With regard to Cr vaporization as a major reason for long-term degradation in SOFCs the results show that alumina-forming-alloys and aluminized surfaces are promising for application in SOFC components without a current-conducting function. © 2008 Elsevier B.V. All rights reserved.
1. Introduction High-temperature alloys rely for their oxidation protection on the formation of a surface oxide scale of either Cr2O3, Al2O3 or SiO2 [1]. With respect to solid oxide fuel cell (SOFC) technology, the chromiaforming materials are being discussed for application in metalsupported SOFC designs and interconnector concepts because alumina and silica possess electronic conductivities which are far too small for most SOFC stack designs [2]. Alumina-forming alloys are being considered as candidate materials for heat exchangers, pumps, piping, casings and for high-temperature components without currentconducting function. In planar SOFC systems, Cr vaporization from metallic construction materials at high temperatures was identified as a major source of cell degradation. The gaseous Cr-species are reduced at the cathode to solid chromia and in this way inhibit the electrochemical processes [3–6]. In a previous paper the present authors reported on Cr vaporization from chromia- and spinel-forming steels [7]. This paper deals with alumina-forming alloys as well as aluminized chromiaforming alloys. Due to their high corrosion resistance and stability aluminaforming materials are used for a large number of applications such as catalyst supports in automotive applications or as furnace parts for temperatures up to 1300 °C [8]. Alumina-forming alloys usually contain about 20 wt.% chromium. [9]. The chromium is added because
⁎ Corresponding author. Tel.: +49 2461 614470; fax: +49 2461 613699. E-mail address: [email protected] (T. Markus). 0167-2738/$ – see front matter © 2008 Elsevier B.V. All rights reserved. doi:10.1016/j.ssi.2008.09.015
it promotes the formation of the protective alumina scale during the oxidation process. It is believed that chromium reduces the ingress of oxygen into the alloy thereby promoting oxidation of Al in form of an external α-Al2O3 scale instead of becoming oxidized internally. Chromia crystals further act as nucleation sites for α-alumina in the initial phase of the oxidation process [10]. Most investigations on alumina-forming materials address the oxidation behaviour at temperatures N1000 °C. Less publications deal with the oxidation behaviour in the temperature range between 800 °C and 950 °C which is of interest for SOFC applications [11–19]. Due to their high-Al contents, alumina-forming Fe- and Ni-based alloys possess poor room temperature ductility. Aluminizing improves the oxidation resistance of high-temperature alloys, for example in Nibase alloys for turbine-blades [20], without substantially affecting the mechanical properties of the core of the material. In the aluminizing process, the surface of Ni-, Fe-, Co- and Cr-based alloys is enriched with aluminium at temperatures between 700 °C and 1100 °C [21]. The diffusing Al forms intermetallic phases with the base material, e.g. βNiAl or FeAl [10,22–24]. At high temperatures the intermetallic phases form protective alumina scales. In this way, aluminizing offers the possibility of obtaining the protection of a thermally grown alumina scale to a variety of alloys without changing the mechanical properties in the core. In the literature, hardly any information is available on the vaporization of alumina-forming steels. Hilpert and Miller [25] determined the partial pressures and activities of Al, Cr and Fe/Ni in FeCrAl and NiCrAl alloys by Knudsen effusion mass spectrometry (KEMS). The vaporization of gaseous Al species from alumina was reviewed by Opila [26]. Chromium vaporization as well as the influence of metastable aluminas on the Cr vaporization from alumina-forming alloys has not yet been studied. However, in several
M. Stanislowski et al. / Solid State Ionics 179 (2008) 2406–2415
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Table 1 Chemical compositions of the investigated alloys determined by ICP-OES
Aluchrom YHf Aluchrom YB Kanthal AF PM 2000 Nicrofer 6025 HT Nicrofer 45 TM Ferrotherm 4828 Ducrolloy
Fe
Cr
Ni
Al
Si
Mn
Ti
Y
Zr
C
N
Others
74.7 74.2 74.5 76.5 9.3 23.9 65.8 5.5
20.2 20.5 21.2 18.6 25.0 27.3 19.2 92.9
0.18 0.14 0.19 0.02 62.6 45.9 11.3 –
5.7 5.2 5.5 5.7 2.3 0.1 – –
0.3 0.1 0.3 – 0.1 2.7 1.7 –
0.2 0.2 0.2 0.1 0.1 0.1 1.3 –
– 0.05 0.09 0.45 0.20 0.06 0.05 –
0.04 0.03 0.03 0.33 0.08 0.08 – 0.48
0.05 0.04 0.06 0.01 0.07 – – 0.02
0.024 0.022 0.017 0.014 0.17 0.053 0.048 0.010
0.005 0.005 0.014 0.006 0.104 – 0.046 0.014
Hf: 0.04 – Cu: 0.04 O: 0.254 – Ce: 0.05 P: 0.022 O: 0.449
technological applications alumina scales are regarded as protective coatings against Cr vaporization at high temperatures [27,28]. The aim of the present work was to measure the chromium release from alumina-forming alloys and aluminized surfaces by means of the transpiration method at temperatures between 800 °C and 1000 °C. The suitability of alumina scales or coatings for SOFC application concerning the retention of Cr will be discussed. 2. Experimental The following commercial alloys were investigated: the ferritic alumina-forming wrought alloys, Aluchrom YHf (material number 1.4767, ThyssenKrupp VDM, Werdohl, Germany), Aluchrom YB (material number 1.4767, ThyssenKrupp VDM, Werdohl, Germany), Kanthal AF (Kanthal AB, Hallstahammar, Sweden), the ferritic, alumina-forming oxide dispersion strengthened (ODS) alloy, PM 2000 (material number 1.4768, Plansee AG, Reutte, Austria), the high-Al Ni-based alumina-forming alloy, Nicrofer 6025 HT (material number 2.4633, ThyssenKrupp VDM, Werdohl, Germany). Additionally, two chromia-forming high-temperature alloys were used to test the suitability of an aluminizing treatment: the Ni-based alloy, Nicrofer 45 TM (material number 2.4889, ThyssenKrupp VDM, Werdohl, Germany) and the austenitic Fe-based alloy, Ferrotherm 4828 (material number 1.4828, Thyssen Schulte GmbH, Dortmund, Germany). As reference material for all the studies the pure chromiaforming Cr-based ODS alloy, Ducrolloy (Cr5Fe1Y2O3, Plansee AG, Reutte, Austria) was incorporated in the test programme. The chemical compositions of these alloys, determined by optical emission spectroscopy using an inductively coupled plasma (ICP-OES, IRISadvantage, Thermo Jarrell Ash, Franklin, USA), are listed in Table 1. For the vaporization experiments, samples with the dimensions 80 × 20 mm2 and a thickness of 2–3 mm were cut from plates by water jet cutting. The edges of the samples were rounded by spark erosion in order to avoid preferred sites for oxide scale cracking and/or spallation during the vaporization experiments. Details of the geometry of the samples are given in [7]. The samples were ground with SiC paper up to 1200 grit and ultrasonically cleaned in ethanol and acetone. Preoxidations were performed in a muffle furnace at 1200 °C for 24 h in air. In order to investigate Cr retention by aluminide coatings, some samples of alloys Nicrofer 45 TM and Ferrotherm 4828 were ground with SiC paper up to 1200 grit, ultrasonically cleaned in ethanol and acetone and aluminized using chemical vapor deposition (CVD) at 1080 °C by MTU Aero Engines GmbH, Munich, Germany [29]. The transpiration method was used to measure the chromium vaporization rates. The method is based on the principle that the vapor over the heated sample is carried away by a constant gas flow and collected in a condenser. The transported mass was subsequently determined by quantitative chemical analysis. The transport rate at a given temperature was calculated from the transported mass of Cr, the surface area of the sample and the duration of the experiment. The influence of the gas flow rate was eliminated in this work by setting the gas flow to rates sufficiently high to ensure the transported mass
of Cr to depend solely on the kinetics of the vaporization reaction at the surface of the sample (see Refs. [7,30]). The experiments were carried out with a constant flow of humidified air at temperatures between 800 °C and 1000 °C. The air flow was controlled by a flow meter, type 5850TR, supplied by Brooks, Veenendaal, The Netherlands. For the experiments the air flow rate was set to 1500 ml/min, relative to standard conditions (273 K, 101,325 Pa). The humidity of the air was adjusted by a bubble humidifier in combination with a condenser. The temperature of the condenser was controlled by a thermostat, type C10/K15, supplied by ThermoHaake, Karlsruhe, Germany. For the experiments, a humidity of p(H2O) = 1.88 · 103 Pa was adjusted, which corresponds to a relative humidity of 60% at 25 °C and standard air pressure. After the experiments, the Cr condensate was dissolved in HCl (Merck, Suprapur®) and quantitatively analysed by inductively coupled plasma mass spectrometry, ICP-MS (PerkinElmer, Norwalk, USA). Details of the experimental set-up and the used analysis method are given in Refs. [7,31]. After the vaporization experiments, the samples were characterized by X-ray diffraction (XRD, X'Pert MRD, Philips, The Netherlands, Cu Kα radiation) and scanning electron microscopy (SEM, LEO 440 and LEO 1530-Gemini, Cambridge, UK) coupled with energy dispersive Xray analysis (EDX, ISIS 300, Eynsham, UK). The concentration of the elements Al, Fe and Cr in the oxide scale was investigated as a function of the depth by electron-collision sputtered neutrals mass spectrometry (SNMS, Simslab 410a, Fisons, East Grinstead, UK). The surface was sputtered with O+2 ions at an energy of 10 keV and a current of 250 nA. A grid pattern of 420 × 420 μm2 and a gate of 170 × 170 μm2 were adjusted. The current of the ions 27Al+, 56Fe+ and 52Cr+ was determined as a function of the
Fig. 1. Comparison of the chromium vaporization rates of different chromia- and alumina-forming alloys at 900 °C in humid air after preoxidation for 24 h at 1200 °C in air. Ducrolloy was preoxidized at 900 °C for 100 h in air.
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Fig. 2. SEM images of cross section and plan view of a), b), Aluchrom YHf, and c), d), PM 2000, after oxidation at 1200 °C for 24 h in air.
sputtering time. An isotope correction was performed with respect to the frequency of the analysed isotopes. 3. Results 3.1. Alumina- and chromia-forming alloys 3.1.1. Comparison of Cr vaporization from different alloys Samples of the Aluchrom YHf, Aluchrom YB, Kanthal AF, PM 2000 and Nicrofer 6025 HT alloys were preoxidized at 1200 °C for 24 h in air to obtain a dense α-Al2O3 layer. The oxidized surfaces of the samples of the first four alloys looked bright grey whereas that of Nicrofer 6025 HT looked dark green with grey speckles. The Cr-based alloy Ducrolloy was preoxidized at 900 °C for 100 h in air. After oxidation it had a dark green colour. The chromium vaporization of the preoxidized samples was measured using the transpiration apparatus at 800 °C in air with a humidity of 1.88%. In order to obtain a condensate with a sufficiently high Cr content for the chemical analyses, the duration of the
measurements was 300 h for the ferritic alumina formers Aluchrom YHf, Aluchrom YB, Kanthal AF and PM 2000, and 40 h for Nicrofer 6025 HT and 33 h for Ducrolloy. Each sample was measured twice in succession. For the preoxidized samples no time-dependent vaporization behaviour was observed. The results of the vaporization experiments are shown in Fig. 1. After the transpiration tests XRD and SEM/EDX revealed the oxide scales on Aluchrom YHf, Aluchrom YB, Kanthal AF and PM 2000 to consist of α-Al2O3, whereas the oxide scale on Ducrolloy consisted of pure Cr2O3. Nicrofer 6025 HT formed a (Cr,Fe)2O3 scale with locally a 0.3 μm thick top layer of Cr–Ti-oxides, presumably (Cr,Mn)2Ti3O9. The Al in Nicrofer 6025 HT was internally oxidized to α-Al2O3. SEM images of Aluchrom YHf and PM 2000 after preoxidation at 1200 °C for 24 h in air are shown in Fig. 2. Fig. 3 shows SEM images of Nicrofer 6025 HT after oxidation at 1200 °C for 24 h in air. The alumina scales on Aluchrom YHf, Aluchrom YB, Kanthal AF and PM 2000 had thicknesses of 7.95 μm, 7.67 μm, 7.18 μm and 5.36 μm, respectively. Small surface nodules were visible in the backscattered
Fig. 3. SEM images of a) cross section and b) plan view of Nicrofer 6025 HT after oxidation at 1200 °C for 24 h in air.
M. Stanislowski et al. / Solid State Ionics 179 (2008) 2406–2415
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Fig. 4. Chromium vaporization rates as a function of time for non-preoxidized Aluchrom YHf, Nicrofer 6025 HT and Ducrolloy at 800 °C in humid air.
Fig. 6. Chromium vaporization rates as a function of time for Aluchrom YHf at 800 °C, 900 °C and 1000 °C in humid air, without preoxidation. The data points were fitted to a power law time dependence of the form y = a ∙ t b.
electron images of Aluchrom YHf and PM 2000 (see Fig. 2b and d). In the case of Aluchrom YHf they contained Zr and Y with impurities of Cr and in the case of PM 2000 of Y and Ti with minor amounts of Ca and Fe. Mg and Cr were found in the outer part of the alumina based scale on Aluchrom YHf. Mg as well as Ca is used for deoxidation in the fabrication of steels. Trace amounts of these elements may therefore be still present in the finished product. The Cr vaporization of non-preoxidized samples of Ducrolloy, Nicrofer 6025 HT and Aluchrom YHf was investigated as a function of time in order to correlate the Cr vaporization with the progress of growth of the oxide scale. Fig. 4 shows the Cr vaporization rates at 800 °C in humid air during oxidation tests up to 500 h. Fig. 5 shows SEM images of the alumina scales of Aluchrom YHf and Nicrofer 6025 HT after exposure at 800 °C for 10 h and 500 h, respectively. After oxidation for 10 h at 800 °C in air the alumina scale on Aluchrom YHf had a thickness of about 0.3 μm and consisted of α- and γ-alumina, according to XRD. The alumina scale on Nicrofer 6025 HT had a thickness of about 0.2 μm after oxidation for 500 h at 800 °C in air and consisted of α-alumina.
10 h and 1000 h at 900 °C and 1000 °C in humid air. The XRD patterns of Aluchrom YHf after oxidation at 800 °C for 10 h, 50 h, 150 h, and 1000 h in humid air are shown in Fig. 7. SEM images of the oxide scales formed on Aluchrom YHf after oxidation for 1000 h in humid air at 800 °C, 900 °C and 1000 °C are shown in Fig. 8. EDX line scans of the elements O, Al, Cr, Fe and Mg measured in metallographic cross sections after oxidation at 1000 °C for 1000 h in humid air are shown in Fig. 9. Mg is found in the outer part of the Al scale (Fig. 9). Fig. 10a–c shows SNMS depth profiles of the elements Fe, Al and Cr in the oxide scale on Aluchrom YHf after oxidation at 900 °C in air for 10 h, 100 h and 1000 h, respectively. Fig. 10a–c shows that there is a strong gradient in the Cr concentration in the alumina scale after exposures of 10 h and 100 h at 900 °C that disappears after 1000 h at 900 °C.
3.1.2. Temperature dependence of Cr vaporization on alumina-forming steels The Cr vaporization rates from non-preoxidized Aluchrom YHf were measured in humid air as a function of time for the temperatures 800 °C, 900 °C and 1000 °C. The results are shown in Fig. 6. On Aluchrom YHf, the phase α-Al2O3 and a spinel structure, probably MgAl2O4 were detected by XRD after oxidation for between
3.2. Aluminized alloys The Cr vaporization rates of aluminized samples of the Nicrofer 45 TM and Ferrotherm 4828 alloys were measured at 800 °C in humid air as a function of time without preoxidation. For comparison, the same measurements were carried out with non-aluminized samples of the same alloys. The results are shown in Fig. 11. According to XRD and SEM, the oxide scale on Nicrofer 45 TM consisted of a chromia scale on top of a thin silica layer. Ferrotherm 4828 formed a chromia scale with an outer (Cr,Mn)-spinel layer which tended to spall off during the vaporization experiments. The aluminized samples formed well-
Fig. 5. SEM images of a) Aluchrom YHf and b) Nicrofer 6025 HT, after oxidation at 800 °C in air for 10 h and 500 h, respectively.
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Fig. 7. XRD patterns of Aluchrom YHf after oxidation at 800 °C in air for 10 h, 50 h, 150 h, and 1000 h.
adherent, dark grey alumina scales already after very short exposure times. The aluminized surfaces were investigated by XRD and SEM/EDX. Images of the microstructure of the aluminized samples are shown in Fig. 12. During the aluminizing process an Al-enriched zone was formed near the surface of Nicrofer 45 TM with a thickness of about 15 µm consisting mainly of β-NiAl and FeAl. Beneath the Al enrichment zone an interdiffusion zone with a width of about 10 µm was identified by SEM. The surface of the aluminized Nicrofer 45 TM was decorated with small precipitates, especially near the grain boundaries. EDX analyses showed that these inclusions consisted of Cr- and Fe-rich phases. After oxidation for 100 h at 800 °C in humid air, the formation of a surface scale of α- and γ-alumina with a thickness of 1.2–1.8 μm was observed. The Al-enriched zone in Ferrotherm 4828 consisted mainly of NiAl, FeAl and other Fe aluminides and had a width of about 20 μm. A deep interdiffusion zone with a width of about 150 μm was present beneath the aluminized zone. After oxidation at 800 °C for 100 h in humid air a surface scale consisting of α- and γ-alumina was observed (Fig. 12).
Fig. 8. SEM cross sections and plan views of the oxide scale of Aluchrom YHf after oxidation for 1000 h in air at a), b), 800 °C, c), d), 900 °C, and e), f), 1000 °C.
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4. Discussion Of the preoxidized samples in Fig. 1, the highest Cr vaporization rate is, as expected, observed for Ducrolloy (Cr5Fe1Y2O3), which forms a surface scale of virtually pure chromia. The Cr vaporization rate for the Ni-based alloy, Nicrofer 6025 HT, preoxidized for 24 h at 1200 °C in air, is similar to that of Ducrolloy. This can be explained by the formation of a chromium oxide scale on top of the alloy (see Fig. 3a). The aluminium present in this alloy was oxidized internally and did not form an external alumina scale. The surface of the Cr oxide was covered locally with a 200–300 nm thick layer of (Cr, Mn)2Ti3O9 on top of the iron/chromium oxide scale. This Ti–Cr-rich layer may reduce Cr vaporization but does not form an effective barrier layer. After preoxidation for 24 h at 1200 °C in air the chromium vaporization rates of the alumina formers Aluchrom YHf, Aluchrom YB, Kanthal AF and PM 2000 are about three orders of magnitude lower
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than that of Ducrolloy. These alloys are all of the Fecralloy type with the nominal composition (in mass%) Fe20Cr5Al + (Y, Zr, Hf) and form dense surface scales of α-Al2O3. Only minor differences in the Cr vaporization rates within this group of alloys were found. The alumina scale on the ODS alloy PM 2000 is significantly thinner and contains fewer pores than that of the wrought alloys Aluchrom YHf, Aluchrom YB and Kanthal AF. According to the literature, the growth of the reactive element (RE) doped alumina scale on the alloy is mainly governed by inward oxygen transport via grain boundaries [32]. Minor contents of reactive elements, such as Y, Zr, Ce, La, and Hf [33] block the grain boundaries and suppress the outward diffusion of cations [32,34]. PM 2000 contains the highest RE content of the investigated Fecralloy type materials (see Table 1). This may explain the small thickness of the alumina scale. Another possible explanation for the differences in oxidation rate might be the shape in which the RE prevails in the alloy (in metallic form or in oxide dispersion), the amount of RE and the type
Fig. 9. EDX line scans of the elements O, Al, Cr, Fe, and Mg through the alumina scale of Aluchrom YHf after annealing at 1000 °C for 1000 h in air.
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that of Aluchrom YHf, Kanthal AF or Aluchrom YB. In this way, a smaller surface area is exposed to the environment, which leads to lower vaporization rates compared to samples with rough oxide surfaces. The RE content leads to the formation of RE-rich nodules on the surface of the alumina scale of PM 2000 and Aluchrom YHf (Fig. 2b and d). The formation of such surface nodules is also known from the literature [13,32]. Their formation and growth is attributed to the outward transport of REs and minor alloying additions (e.g. Mg, Ti) via grain boundaries [34]. EDX analyses showed that these nodules contained Y and Ti with small amounts of Ca and Mg, depending on the type of alloy. Small concentrations of Cr were detected in the nodules and also in alumina grains at the surface. The presence of Cr in nodules on the surface of alumina-forming alloys was also reported in the literature [9,37,38]. A high concentration of Mg was detected in the outer part of the alumina scale of Aluchrom YHf, probably in the form of Al–Mg spinel, which corresponds to the observations by Dimyati et al. [39]. The presence of Cr in the nodules as well as on the surface of the alumina grains and the small difference in Cr vaporization between PM 2000 and Aluchrom YHf, Kanthal AF and Aluchrom YB, indicates that the transport of Cr by grain boundaries is not the dominating factor for Cr vaporization. The transpiration experiments with non-preoxidized samples in Fig. 4 show the dependence of the Cr vaporization rate on the initial growth of the oxide scale. The Cr vaporization rates of Ducrolloy in Fig. 4 decrease within the first 50 h. This can be explained by the fact that initially the formation of Cr(g) contributes to the vaporization process whereas after a dense scale has formed, the overall vaporization process is governed by vaporization processes at the scale/gas interface [40]. After the surface is completely covered with a dense chromia scale the vaporization rate is approximately constant. In contrast to this, the Cr vaporization rate of Nicrofer 6025 HT decreases steadily with increasing time. After 500 h of oxidation at 800 °C in air the Cr vaporization rates are reduced by 90% compared to the pure chromia scale of Ducrolloy. This can be explained by the formation of a slowly-growing alumina scale which reaches a thickness of 200–300 nm after oxidation for 500 h at 800 °C in air (see Fig. 5b). After oxidation of the same alloy for 24 h at 1200 °C in air the aluminium was oxidized internally (see Fig. 3). This different oxidation behaviour illustrates that the Nicrofer 6025 HT alloy cannot be considered as a “real” alumina-forming alloy. The formation of the alumina scale on Ni-base alloys with low aluminium contents, such as Nicrofer 6025 HT, depends strongly on the oxidation conditions [1]. The thickness of the alumina scale on Nicrofer 6025 HT after oxidation for 500 h at 800 °C in air is comparable to the thickness of the alumina scale on Aluchrom YHf after 10 h at 800 °C in air (see Fig. 5a).
Fig. 10. SNMS depth profiles of the elements Fe, Al and Cr in the oxide scale on Aluchrom YHf after oxidation at 900 °C in humid air for a), 10 h, b), 100 h and c), 1000 h.
of RE [35]. Also inclusions of REs such as Hf and especially Zr, which prevail e.g. in Aluchrom YHf, tend to induce microporosity in the alumina scale [36]. This might be an additional rapid pathway for Cr from the alloy to the oxide surface. Apart from the thinner, denser oxide scale the slightly lower Cr vaporization rates of PM 2000 may be affected by the smoother morphology of the oxide surface compared to
Fig. 11. Chromium vaporization rates as a function of time at 800 °C in humid air for non-preoxidized and aluminized samples of Nicrofer 45 TM and Ferrotherm 4828.
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Fig. 12. SEM cross sections and plan views of the oxide scale of Nicrofer 45 TM, a), b), after aluminizing, c), d), after aluminizing and annealing at 800 °C for 100 h in air, and, e), f), of Ferrotherm 4828 after aluminizing and annealing at 800 °C for 100 h in air.
The vaporization rates of these two alloys were very similar after these oxidation times (see Fig. 4). After 100 h of oxidation at 800 °C in air the alumina scale on Aluchrom YHf reached a thickness of 0.6–0.9 μm and the Cr vaporization rates were more than two orders of magnitude lower than that of Nicrofer 6025 HT after 500 h at 800 °C. The Cr vaporization rates of Aluchrom YHf were more than 3 orders of magnitude lower than those of Ducrolloy. Similar vaporization rates were measured for the preoxidized samples (see Fig. 1). This demonstrates the high Cr retention capability of alumina. The types of alumina phases formed after oxidation at 800 °C were identified by XRD as α-alumina, in the case of Nicrofer 6025 HT, and as a mixture of α- and γ-alumina, in the case of Aluchrom YHf. In the latter case, the alumina scale consisted of a layered structure consisting of an outer layer of metastable γ-alumina, which formed characteristic platelets, and an inner layer of stable α-alumina (see Figs. 5a and 8a). Similar results were reported by other authors [10,37,41]. The formation of the layered structure can be explained by the growth of metastable aluminas by outward cation diffusion [16,42] and the growth of α-alumina by inward oxygen diffusion [43].
After 150 h of oxidation at 800 °C in air the metastable θ-alumina phase was detected on Aluchrom YHf by XRD (see Fig. 7), next to γand α-alumina. In the literature several papers are available dealing with the metastable alumina phases formed at lower temperatures. These range from θ-alumina [12,13,15,32,44] to γ-alumina [16,45] and combinations of the metastable phases of γ-, δ- and θ-alumina [19]. In the present work, it was found that the θ-, γ- and α-alumina phases co-exist after oxidation at 800 °C for 1000 h in air and that the amount of metastable phases increased with time (see Fig. 7). It is generally assumed that metastable aluminas are formed at the beginning of the oxidation process and are replaced by stable α-alumina at times ranging from minutes to hours depending on the temperature [12,23]. The sequence of the phase transformation ranges is given in the literature beginning with metastable γ-alumina to δ- and θ-alumina to finally stable α-alumina [23,46]. The formation and morphology of the metastable aluminas are influenced by several factors such as small deviations in the chemical composition [15], pre-existing oxides [10], presence of water vapor [47], surface defects such as grain boundaries [16] or previous cold deformation [32].
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Fig. 6 shows the influence of temperature on the Cr vaporization rate of non-preoxidized Aluchrom YHf. The oxide scales formed at 900 °C and 1000 °C consisted exclusively of α-alumina. Even after relatively short oxidation times of 10 h at 900 °C and 1000 °C no metastable alumina phases could be detected by XRD. This is in contrast to the oxidation test at 800 °C where the oxide scale always consisted of γ-, θ-, and α-alumina. Accordingly, there must be a transition temperature between 800 °C and 900 °C above in which only α-alumina is formed. In the literature, this transition temperature is given as 900 °C [13], 950 °C [48] or 1050 °C [49]. This transition temperature will, however, depend on exposure time. The Cr vaporization rates at 1000 °C are more than one order of magnitude higher than at 800 °C and a factor of 3–4 higher than at 900 °C. The time dependence of the Cr vaporization rates at 1000 °C and 900 °C are comparable but differ from that at 800 °C. At 800 °C the Cr vaporization rates show a stronger decrease with increasing time. During the initial stage of oxidation the Cr vaporization strongly depends on the growth of the alumina scale. The time dependence of the Cr vaporization rates was fitted to a power law time dependence of the form y = a ∙ tb. The exponents of the best fits to the curves at 900 °C and 1000 °C were b = −0.67 and b = −0.69, respectively. For the curve of the Cr vaporization rates at 800 °C an exponent of b = −1.01 was obtained for the best fit (see Fig. 6). These results indicate that the kinetics of the Cr vaporization at 900 °C and 1000 °C are similar but differ from those at 800 °C. This is in agreement with the observations of other authors who measured a fundamental change in oxidation kinetics with the temperature increasing from 800 °C to 900 °C [12,50]. In the literature, this effect was attributed to the formation of metastable aluminas at lower temperatures [16], sometimes accompanied by crack formation caused by the transformation of metastable to stable alumina [41]. After exposures of 1000 h at 1000 °C, 900 °C and 800 °C in air the SEM images of the cross sections of the alumina scales in Fig. 8 showed alumina scales with thicknesses of 3.5–4.2 μm, 1.4–2.6 μm and 0.9– 2.6 μm, respectively. It is noteworthy that the alumina scales formed during oxidation at 800 °C and 900 °C within 1000 h have similar thicknesses. This can be explained by the formation of different types of alumina phases. It is well known from the literature that metastable aluminas grow faster than α-alumina [12,13,22]. Metastable aluminas are cation-deficient and are therefore considered to possess a lower Cr retention capability than α-alumina [27,28]. This cannot be confirmed by transpiration experiments performed in the work. This might be explained by the formation of a dense α-alumina layer under the metastable γ-/θ-alumina scale, as already mentioned. In Fig. 6, the Cr vaporization rates measured at 1000 °C after the same exposure time are shown to be a factor 4–5 higher than those at 900 °C. The alumina scale formed on Aluchrom YHf after an oxidation of 1000 h at 1000 °C is about twice as thick as that which was formed at 900 °C (see Fig. 8). Vaporization tests [51] with preoxidized samples of Aluchrom YHf and PM 2000 at 800 °C, 900 °C and 1000 °C showed that the Cr vaporization rates follow an Arrhenius-type temperature dependence. According to the Hertz–Langmuir-equation the vaporization rates are proportional to the partial pressures. In our experiments we did not directly measure the partial pressures but rather the quantities of vaporized elements per time and area. As the most abundant volatile Crspecies are of the one-atomic form CrO2(OH)2, CrO2(OH) and CrO3 the measured vaporization rates are directly proportional to the partial pressures. Therefore one can follow in this context: 1 pðCrÞ e exp − T
ð1Þ
This type of temperature dependence is valid for the vaporization rates as well as for the diffusion. The question remains of how the Cr content in the alumina scale is influenced by oxide growth. It is known that Cr2O3 and Fe2O3 are
formed in the initial stage of oxidation on the surface of ferritic alumina-forming alloys, before the dense Al2O3 layer is formed [10,11,14,15]. After a few minutes or hours (depending on temperature) the transient oxides are replaced by Al2O3. It is proposed that Cr2O3 crystals act as nucleation sites for α-Al2O3 [10]. The initially high Cr concentration in the oxide scale thus drops significantly within the first few hours of oxidation as was shown by Quadakkers et al. [11,52]. The SNMS measurements after oxidation at 900 °C in air for 10 h and 100 h and 1000 h in Fig. 10a–c shows that a gradient of Cr exists within the alumina scale. With increasing oxidation time the gradient flattens and after 1000 h of oxidation at 900 °C the concentration of Cr at the surface is larger by about a factor of 2 than after 10 h or 100 h. This will lead to Cr vaporization rates that are higher by a factor of 2 after long oxidation times. The decrease of the Cr gradient in the alumina scale after 1000 h of oxidation can be explained by the decrease of the scale thickening rate of the alumina scale with increasing time and thickness of the scale. The concentration gradient of Cr is then flattened by the slower diffusion transport. The transpiration experiments with aluminized samples in Fig. 11 show that surface modification by aluminizing offers a similar protection against Cr vaporization of chromia-forming alloys as found for the alumina-forming steels. The measured Cr vaporization rates of non-preoxidized samples of the aluminized Ni-base alloy Nicrofer 45 TM and the aluminized austenitic steel Ferrotherm 4828 are comparable to that of Aluchrom YHf. After annealing at 800 °C for 100 h in air the alumina scales grown on the aluminized samples have a thickness of 1.2–1.8 μm, which is about twice the thickness of the alumina scale on Aluchrom YHf after the same oxidation time (see Fig. 12). Grabke [24] and Prasanna et al. [13] reported that minor additions of Y suppress the formation of θ-alumina and encourage the formation of α-alumina. In contrast to Aluchrom YHf, the aluminized samples contain no Y. The faster growth of the alumina scale on the aluminized alloys can therefore be explained by the slower phase transformation from rapid-growing metastable γ- and θ-aluminas to slowly-growing stable α-alumina [12,13,22]. Another explanation might be, that the rate of nucleation of α-Al2O3 depends on the underlying metallic matrix [53]. For aluminized Ferrotherm 4828 it was observed that the morphology of the alumina scale formed on the alloy grains differed from that near the alloy grain boundaries (see Fig. 12). This can be attributed to a slower diffusion of aluminium in the austenitic grains compared to the diffusion in the ferritic grains and along the grain boundaries. The more rapid aluminium diffusion in the ferritic grains is also responsible for the deeper ingress of aluminium in Ferrotherm 4828 than in Nicrofer 45 TM. The narrow Al-diffusion zone of 30–200 μm of the aluminized alloys is expected to not substantially affect the mechanical properties of the bulk of the alloy. Due to the presence of the aluminium reservoir in the surface-near regions the alumina scales on aluminized alloys grow and are self-healing and are therefore more durable and reliable than e.g. ceramic coatings. Aluminizing thus offers an effective and attractive measure to obtain protection against Cr vaporization by a thermally grown and self-healing oxide scale. 5. Conclusions The measurement of Cr vaporization from alumina scales on alumina-forming and aluminized alloys by the transpiration method shows that Cr vaporization rates from alumina scales are about three orders of magnitude lower than those from pure chromia scales. It can be assumed that the kinetics of the Cr release is governed by the scale growth. At 1000 °C and 900 °C the kinetics of the scale growth are similar. At 800 °C a different scale growth kinetics could be identified from the results of Cr transpiration experiments which can be explained by the formation of metastable γ- and θ-aluminas. The chromium retention of metastable aluminas was comparable to that of stable α-alumina.
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The Cr retention of aluminized surfaces of Ni- and Fe-base alloys was proven to be as efficient as that of alumina-forming steels. Thus, aluminizing offers the possibility to reduce chromium vaporization from high-temperature alloys by about two to three orders of magnitude compared to chromia-forming materials. This makes aluminizing a very interesting method for the protection of metal components in SOFC systems which do not have a current-conducting function. Acknowledgements The authors are grateful to their colleagues at Research Centre Jülich for their assistance with the work. Special thanks are due to Mr. H. J. Penkalla and Ms. D. Esser (IWV-2) for TEM/FIB studies, Mr. U. Breuer (ZCH) for SNMS measurements and Mr. P. Lersch (IWV-2) for XRD measurements. ThyssenKrupp VDM, Werdohl (Germany), Plansee AG, Reutte (Austria) and Kanthal AB, Hallstahammar (Sweden) are acknowledged for providing the materials. A special acknowledgment is given to Klaus Hilpert who suddenly and unexpectedly died in retirement. We acknowledge him for his valuable contributions to this paper. References [1] W.J. Quadakkers, J. Piron-Abellàn, V. Shemet, L. Singheiser, Mater. High Temp. 20 (2003) 115. [2] W.A. Meulenberg, A. Gil, E. Wessel, H.P. Buchkremer, D. Stöver, Oxid. Metals 57 (2002) 1. [3] D. Das, M. Miller, H. Nickel, K. Hilpert, in: U. Bossel (Ed.), Proceedings of the 1st European Solid Oxide Fuel Cell Forum, Lucerne, Switzerland, October 3–7, (1994) 703 (EPFL). [4] S. Taniguchi, M. Kadowaki, H. Kawamura, T. Yasuo, Y. Akiyama, Y. Miyake, T. Saitoh, J. Power Sources 55 (1995) 73. [5] K. Hilpert, D. Das, M. Miller, D.H. Peck, R. Weiß, J. Electrochem. Soc. 143 (1996) 3642. [6] Y. Matsuzaki, I. Yasuda, Solid State Ionics 132 (2000) 271. [7] M. Stanislowski, E. Wessel, K. Hilpert, T. Markus, L. Singheiser, J. Electrochem. Soc. 154 (2007) A295. [8] I. Itoh, M. Fukuya, R. Hisatomi, H. Morimoto, K. Ohmura, H. Tanaka, F. Fudanki, M. Arakawa, Nippon Steel Tech. Rep. 64 (1995) 69. [9] G.C. Wood, Oxid. Metals 2 (1970) 11. [10] M.W. Brumm, H.J. Grabke, Corr. Sci. 33 (1992) 1677. [11] W.J. Quadakkers, A. Elschner, W. Speier, H. Nickel, Appl. Surf. Sci. 52 (1991) 271. [12] W.J. Quadakkers, K. Schmidt, H. Grübmeier, E. Wallura, Mater. High Temp. 10 (1992) 23. [13] K.M.N. Prasanna, A.S. Khanna, R. Chandra, W.J. Quadakkers, Oxid. Metals 46 (1996) 465.
2415
[14] J. Jedliński, A. Glazkov, M. Konopka, G. Borchardt, E. Tscherkasova, M. Bronfin, M. Nocuń, Appl. Surf. Sci. 103 (1996) 205. [15] D. Naumenko, W.J. Quadakkers, A. Galerie, Y. Wouters, S. Jourdain, Mater. High Temp. 20 (2003) 287. [16] H. El Kadiri, R. Molins, Y. Bienvenu, M.F. Horstemeyer, Oxid. Metals 64 (2005) 63. [17] R. Molins, A.M. Huntz, Mater. Sci. Forum 461–464 (2004) 29. [18] H. El Kadiri, R. Molins, Y. Bienvenu, Mater. Sci. Forum 461–464 (2004) 1107. [19] G. Barthomé, E. N'Dah, Y. Wouters, A. Galerie, Mater. Corr. 56 (2005) 389. [20] R. Streiff, J. Phys. IV, C9 3 (1993) 17. [21] H. Simon, M. Thoma, Angewandte Oberflächentechnik für metallische Werkstoffe, Hanser Verlag, München, 1985. [22] G.C. Rybicki, J.L. Smialek, Oxid. Metals 31 (1989) 275. [23] J. Doychak, J.L. Smialek, T.E. Mitchell, Met. Trans. A 20A (1989) 499. [24] H.J. Grabke, Mater. Sci. Forum 251–254 (1997) 149. [25] K. Hilpert, M. Miller, Z. Metallk. 83 (1992) 739. [26] E.J. Opila, Mater. Sci. Forum 461–464 (2004) 765. [27] W.J. Quadakkers, F. Baumanns and H. Nickel, German Patent No. DE 44 10 711 C1 (1994). [28] W.J. Quadakkers, German Patent No. DE 195 47 699 A1 (1995). [29] H. Pillhöfer, M. Thoma, H. Walter and P. Adam, German Patent No. DE 41 19 967 C1 (1991). [30] H.K. Hofmeister, R.v. Haeseler, O. Glemser, Z. Elektrochem. 64 (1960) 513. [31] M. Stanislowski, U. Seeling, D.-H. Peck, S.-K. Woo, L. Singheiser, K. Hilpert, Solid State Ionics 176 (2005) 2523. [32] W.J. Quadakkers, Werkst. Korr. 41 (1990) 659. [33] W.J. Quadakkers, L. Singheiser, Mater. Sci. Forum 369–372 (2001) 77. [34] W.J. Quadakkers, H. Holzbrecher, K.G. Briefs, H. Beske, Oxid. Met. 32 (1989) 67. [35] B. Pint, J. Am. Ceram. Soc. 86 (2003) 686. [36] E. Wessel, V. Kochubey, D. Naumenko, L. Niewolak, L. Singheiser, W.J. Quadakkers, Scripta Mat. 51 (2004) 987. [37] P.T. Moseley, K.R. Hyde, B.A. Bellamy, G. Tappin, Corr. Sci. 24 (1984) 547. [38] M.J. Bennett, R. Newton, J.R. Nicholls, Mater. High Temp. 20 (2003) 347. [39] A. Dimyati, H.J. Penkalla, P. Untoro, D. Naumenko, W.J. Quadakkers, J. Mayer, Z. Metallkd. 94 (2003) 180. [40] P. Kofstad, J. Phys. Chem. Solids 44 (1983) 129. [41] W.J. Quadakkers, D. Naumenko, E. Wessel, V. Kochubey, L. Singheiser, Oxid. Metals 61 (2004) 17. [42] B.A. Pint, J.R. Martin, L.W. Hobbs, Solid State Ionics 78 (1995) 99. [43] J. Jedlinski, R. Borchardt, Oxid. Metals 36 (1991) 318. [44] W.C. Hagel, Corrosion 21 (1965) 316. [45] L. Marechal, B. Lesage, A.-M. Huntz, R. Molins, Mater. High Temp. 20 (2003) 295. [46] K. Wefers, G.M. Bell, Techn. Paper, vol. 19, Alcoa Res. Lab., 1972. [47] J.R. Regina, J.N. DuPont, A.R. Marder, Oxid. Metals 61 (2004) 69. [48] J.A. Nychka, D.R. Clarke, Oxid. Metals 63 (2005) 325. [49] H. Al-Badairy, D. Naumenko, J. Le Coze, G.J. Tatlock, W.J. Quadakkers, Mater. High Temp. 20 (2003) 405. [50] M.J. Bennett, R. Newton, J.R. Nicholls, H. Al-Badairy, G.J. Tatlock, Mater. Sci. Forum 461–464 (2004) 463. [51] M. Stanislowski, PhD thesis, RWTH Aachen University, Germany (2006). [52] W.J. Quadakkers, A. Elschner, H. Holzbrecher, K. Schmidt, W. Speier, H. Nickel, Mikrochim. Acta 107 (1992) 197. [53] D. Clemens, V. Vosberg, L.W. Hobbs, U. Breuer, W.J. Quadakkers, H. Nickel, J. Anal. Chem. 355 (1996) 703.