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Anomalous temperature dependence of oxidation kinetics during steam oxidation of ferritic steels in the temperature range 550–650 °C
Corrosion Science 46 (2004) 2301–2317 www.elsevier.com/locate/corsci
Anomalous temperature dependence of oxidation kinetics during steam oxidation of ferritic steels in the temperature range 550–650 C a,* _ J. Zurek , E. Wessel a, L. Niewolak a, F. Schmitz b, T.-U. Kern b, L. Singheiser a, W.J. Quadakkers a a
Forschungszentrum J€ulich GmbH, IWV-2, 52425 J€ulich, FRG b Siemens Power Generation, M€ulheim, FRG Received 6 August 2003; accepted 16 January 2004 Available online 12 March 2004
Abstract The oxidation behavior of a number of selected ferritic steels in a simulated steam environment at temperatures between 550 and 650 C was studied. In the prevailing test gas, some of the studied 9–12% Cr steels tended to exhibit an anomalous temperature dependence of the oxidation behavior. This means, that the oxidation rates do not steadily increase with increasing temperature. At higher temperatures, some of the studied steels tend to form a very thin and protective oxide scale whereas at lower temperature rapidly growing, less-protective oxides are being developed. The anomalous temperature dependence is related to differences in chromium distribution in the inner part of the oxide scale. The effect is observed for steels with intermediate-Cr contents (10–12%) whereas steels with either lower or higher Cr contents exhibit an increasing oxidation rate with increasing temperature. 2004 Elsevier Ltd. All rights reserved. Keywords: A. Steel; B. SIMS; C. High temperature corrosion; Selective oxidation
*
Corresponding author. Tel.: +48-2461-61-55-22; fax: +48-2461-61-36-99. _ E-mail address: [email protected] (J. Zurek).
0010-938X/$ - see front matter 2004 Elsevier Ltd. All rights reserved. doi:10.1016/j.corsci.2004.01.010
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1. Introduction The need for reduction of CO2 emission and fuel consumption is a leading factor for the requirement of the utilities to increase the thermal efficiency of fossil fired power generation plants [1]. An increase in the steam temperature from common 535 to 650 C and the steam pressure from 185 to 300 bar would enable a reduction in fuel consumption and a reduction of CO2 emission by more than 25% [2]. A number of high strength 9–12% Cr steels have been developed for application as construction materials in such advanced power plants [3]. Whereas all these steels possess excellent oxidation resistance during service in oxygen or air [4], it is well known that in water vapour containing environments, as prevailing under power plant service conditions, the oxidation rates of a number of these steels are significantly enhanced [5–7]. Recently it was shown that during short term isothermal exposure in water vapour-rich environments, some of these steels tended to exhibit a bell-shape temperature dependence of the oxidation rate in the temperature range 600–800 C [8]. Whereas poorly protective oxides were formed at 600 and 650 C, very thin chromium-rich oxides were found at 700 C and especially 800 C, resulting in very low oxidation rates. In the present study it will be shown, that such an anomalous temperature dependence can also occur during exposure in steam if only lower, i.e. more practically relevant temperatures for current power plant application (550–650 C), are being considered. It will also be illustrated, that the effect does not only occur in laboratory tests with short exposure times of up to typically 100 h [8], but it is also of great significance for very long time service in practical application.
2. Experimental The compositions of the studied materials, delivered in form of forged bars, are listed in Table 1. The steels varied in Cr content between 1.3% and 12.4%. One of the steels contained a substantial addition of Co. Specimens with a size of 20 · 10 · 2 mm were machined. For the oxidation studies, the specimens were ground to 800 grit surface finish and exposed to an Ar + 50%H2 O atmosphere. The exposures were carried out for up to 1000 h in the temperature range 550–650 C, whereby the specimens were cooled to room temperature every 250 h for weight measurements. After exposure the oxidation products were characterised by optical metallography, scanning electron microscopy (SEM) with energy dispersive X-ray analysis (EDX) Table 1 Composition of the investigated steels (mass %) Steeldesignation
C
Cr
Mo
Si
Mn
Co
V
W
Ni
Cu
A B C X20
0.17 0.10 0.07 0.20
1.31 10.4 10.8 12.4
0.95 1.05 0.86 1.00
0.51 0.09 0.3 0.26
0.59 0.45 0.62 1.5
– 0.012 6.52 0.03
0.21 0.18 0.27 0.054
– 0.98 – 0.06
0.11 0.73 0.72 0.51
0.07 0.028 0.006 0.1
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and sputtered neutrals mass spectrometry (SNMS). Before mounting for metallographic cross-section analyses, the specimens were sputtered by a thin gold layer and subsequently electroplated by nickel. This coating provided protection of the surface oxide layer during grinding and polishing and it ensured a better optical contrast between oxide and mounting material. 3. Results and discussion 3.1. Different types of temperature dependence Fig. 1(a)–(d) shows the weight change data of the four studied steels during exposure in Ar + 50%H2 O up to 1000 h at different temperatures. In the following the weight change data are given as relative data. The 1% Cr steel is commonly used in practice at temperatures up to 550 C. Therefore all the measured weight change data are normalised to the 1000 h value of the 1% Cr steel (steel A) at 550 C. The low-Cr steel A and the Co-free 10% Cr steel B clearly exhibit increasing oxidation rates with increasing temperature. However, the two other steels show a much more complex temperature dependence of the oxidation rates. After 1000 h of exposure the highest weight change in case of X20 is found for the temperature 550 C. For this steel the weight change drops upon temperature change to 600 C and again increases if the temperature is raised to 625 or 650 C. The Co-containing steel C again 7 625˚C
(a) 8 6
600˚C
4 2
550˚C
0
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Relative weight change
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(b)
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1.6
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(c)
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1 0.8 0.6 0.4
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0 0
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Time (h)
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625˚C
0.8 650˚C
0.6 0.4 0.2 0
600˚C
0
200
400
600
800
1000
Time (h)
Fig. 1. Weight change of various steels during exposure at temperatures between 550 and 650 C in Ar–50%H2 O; (a) 1% Cr steel A, (b) 10% Cr steel B, (c) Co-containing 10% Cr steel C and (d) 12% Cr steel X20.
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2304 8
A
6
B
4
2
C X20
0 540 550 560 570 580 590 600 610 620 630 640 650 660 Temperature (°C)
Fig. 2. Temperature dependence of weight change of the various steels after oxidation in Ar–50%H2 O for 250 h, relative to steel A at 550 C.
exhibits a different type of temperature dependence. After 1000 h exposure at 625 and 650 C the weight changes are clearly lower than those observed during 550 C exposure, whereby the 600 C-value is only slightly smaller than that at 550 C. It is important to note that the high value at 600 C is already reached after 250 h, in other words, the oxidation rate between 250 and 1000 h is similarly small as that observed at 625 and 650 C. Fig. 2 summarizes the weight change data for the four steels after 250 h exposure at the various temperatures. 3.2. Scale morphology in case of normal temperature dependence Fig. 3 shows typical microstructures of oxide scales at the 1% Cr steel A after 1000 h exposure at 550 and 625 C. The scales nearly exclusively consist of magnetite (Fe3 O4 ). In agreement with the gravimetric data, the scale formed at 625 C is substantially thicker than that formed at 550 C. After the 625 C exposure, a thin
Fig. 3. Metallographic cross-sections of 1% Cr steel A showing oxide scales after 1000 h oxidation in Ar–50%H2 O: (a) 550 C and (b) 625 C.
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zone of wustite (FeO) with minor Cr2 O3 precipitates is found near the oxide/alloy interface. The specimen exposed at 550 C exhibits small amounts of hematite (Fe2 O3 ) near the scale/gas interface. In both scales substantial void formation is observed. After the 625 C exposure the voids are occurring in the outer magnetite layer and near the scale/alloy interface. During exposure at 550 C, they concentrate at the interface between inner and outer scale [9], resulting in a tendency to spalling of the outer part of the oxide layer. Fig. 4 shows the microstructures of the scales on the Co-free 10% Cr steel B after 1000 h exposure at 550 and 650 C. Very similar morphologies were found after exposure at 600 and 625 C. The scales typically consist of an outer Fe3 O4 layer and an inner layer containing (Fe,Cr)3 O4 stringers in an Fe3 O4 ‘‘matrix’’ [9]. The two layers are separated by a near-ideally straight line representing the original metal surface. Internal oxide stringers embedded in an FeO-matrix are observed near the scale/alloy interface. After exposure at 550 C the outer scale exhibits substantial void formation near the interface with the inner scale. Additional voids, however of smaller size, are formed near the interface between inner scale and internal oxidation zone. After 650 C exposure this void formation in the inner zone is more pronounced. The few number of voids in the outer scale are present throughout the Fe3 O4 layer. The SEM-picture in Fig. 5 shows, that in addition to the voids which
Fig. 4. Metallographic cross-sections of 10% Cr steel B showing oxide scales after 1000 h oxidation in Ar–50%H2 O: (a) 550 C, (b) 650 C and (c) 550 C––high magnification.
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Fig. 5. High magnification of area near the scale/alloy interface on 10% Cr steel B after 1000 h exposure at 650 C in Ar–50%H2 O (compare Fig. 4b).
can be easily distinguished by light microscopy (Fig. 4), the inner scale contains an infinite by large number of microvoids, next to the precipitates of (Fe,Cr)3 O4 . The transverse cracks visible in the cross-section in Fig. 4b are likely to be related to stresses originating from differences in thermal expansion coefficient between alloy and oxide at high temperatures, the thermal expansion coefficient of magnetite being larger than that of the steel [10], resulting in tensile stresses in the oxide during cooling. In some of the cases, hematite is formed in the scale near the interface with the gas (Fig. 3a). Comparison of the equilibrium oxygen partial pressure in the test gas (approximately 10 8 bar at 650 C) with the Fe2 O3 /Fe3 O4 -equilibrium oxygen partial pressure (approximately 10 13 bar at 650 C) clearly reveals that hematite is the iron oxide phase, which is in equilibrium with the gas atmosphere. In most cases, however, no hematite was found in the outer scale. Apparently no equilibrium between gas and oxide surface has been established during the time of exposure. The ratedetermining step of the scale growth process is in that case not the cation and/or anion transport in the scale, but probably the kinetics of the gas/solid reactions at the scale/gas interface. One might speculate about possible reasons for the formation of the large and tiny voids observed in the scales (Fig. 5). An explanation might be, that they originate from condensation of cation vacancies, if one assumes the growth in magnetite and wustite to primarily occur via Fe cations [11]. In the inner part of the scale, the interfaces between the Fe3 O4 ‘‘matrix’’ and the finely distributed precipitates of (Fe,Cr,Mn)O4 , which were shown to partly originate from the oxidation of carbides [9], could supply numerous vacancy condensation sites resulting in an extremely large number of tiny voids (Fig. 5). As the outer part of the scale consists of virtually pure magnetite, the number of vacancy condensation sites in this part of the scale is extremely small, resulting in fewer but larger voids (Fig. 4).
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Summarizing it can be said that the composition and morphology of the scales formed on steels A and B does not fundamentally differ when exposed at different temperatures. As the rates of anion/cation transport in the scales as well as the rates for the surface reaction are expected to increase with increasing temperature, it is obvious, that the oxidation rates should increase with increasing temperature, as shown in Fig. 1a and b. 3.3. Steels showing anomalous temperature dependence Fig. 6 shows the oxide scales formed on the Co-containing 10% Cr steel C after 1000 h exposure at different temperatures. In agreement with the weight change data in Fig. 1c, the thicknesses of the scales formed at 550 and 600 C are quite similar. However, the scale morphologies show substantial differences. After exposure at 550 C the inner part of the scale exhibits a wide zone of internal oxide precipitates in an FeO-matrix. The outer scale consist of Fe3 O4 plus Fe2 O3 , whereby the latter is not only present at the scale/gas interface but also locally, more inward, i.e. near the large gap in the outer scale. This observation allows the following conclusions:
Fig. 6. Metallographic cross-sections of Co-containing 10% Cr steel C after 1000 h oxidation in Ar– 50%H2 O: (a) 550 C, (b) 600 C and (c) 625 C.
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• The outer surface of the scale is in (near-) equilibrium with the gas (in contrary to the case described in the previous section). • The outer scale allows molecular gas transport. Also after 600 C exposure, the outer scale contains minor amounts of hematite. The gap in the scale is located near the boundary between inner and outer scale i.e. the original metal surface. The most striking difference with the 550 C specimen is the virtual absence of the internal oxidation zone consisting of FeO + Cr2 O3 . In agreement with the weight change data, an increase in temperature to 625 C again leads to a dramatic change in scale morphology (Fig. 6c). A very protective oxide is formed which, according to SNMS-analyses, mainly consists of an inner Cr2 O3 -layer and minor amounts of Fe2 O3 in the outer part of the scale (Fig. 7). No indication was found for presence of Co in the surface scale. A very similar scale morphology was found after the 650 C exposure. Based on these observations, in combination with findings reported in literature [8], the temperature dependence observed for steel C can be explained in the following way: At 550 C the scale initially grows via cation diffusion in the magnetite and wustite lattice. Chromium diffusion in the alloy and in the oxide is so slow that it can be considered as being virtually immobile. It becomes incorporated in the inner part of the scale due to oxidation of the Cr-rich carbides prevailing in the alloy, i.e. the carbides become embedded by the inwardly moving oxidation front [9]. The Cr-rich stringers in the inner scale thus exhibit a sort of ‘‘fingerprint’’ of the typical carbide
100
90
80
Concentration (at%)
Fe 70
O 60
50
40
Cr
30
20
10
Co
Mn 0 0
500
1000
1500
2000
2500
3000
3500
4000
4500
5000
Sputtering Time (s)
Fig. 7. SNMS depth profile of oxide scale on Co-containing 10% Cr steel C after 1000 h oxidation at 625 C in Ar–50%H2 O (compare Fig. 6c).
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Fig. 8. Secondary electron image (a) and X-ray maps showing element distribution in scale formed on Cocontaining 10% Cr steel C after 1000 h exposure at 550 C: (b) O-distribution, (c) Cr-distribution, (d) Fedistribution and (e) Co-distribution (compare Fig. 6a).
morphology in this type of steels. This is confirmed by the element distribution in Fig. 8, showing that in the scale the Cr is hardly enriched compared to the alloy Cr concentration. At 625 and 650 C, Cr diffusion in the alloy in direction of the scale/alloy interface is sufficiently fast to stabilize a chromia based scale, with remnants of transient Feoxides (Fig. 6c). Probably, the partial dissolution of the Cr-rich carbides in the alloy matrix at the higher temperatures contributes to the faster Cr-transport in the alloy.
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The scale formation at 600 C is ‘‘intermediate’’ between the 550 and 625 C case. Based on the initial, very rapid weight change, (Fig. 1c) the scale formation mechanism at 600 C is more similar to that at 550 C than that at 625 or 650 C. However, after longer times the oxidation rate resembles that observed at the higher temperatures. The latter would mean, that during the 600 C exposure somewhere in the scale a protective Cr-rich oxide becomes being stabilized. This is confirmed by
Fig. 9. Secondary electron image (a) and X-ray maps showing element distribution in scale formed on Cocontaining 10% Cr steel C after 1000 h exposure at 600 C: (b) O-distribution, (c) Cr-distribution, (d) Fedistribution and (e) Co-distribution (compare Fig. 6b).
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the element mapping in Fig. 9. Whereas the overall Cr enrichment in the inner part of the scale is not very pronounced if compared with the Cr alloy concentration, a strong enrichment is visible in a narrow band near the scale/alloy interface. Beneath this, a Cr depletion is clearly visible in the alloy as a result of the selective oxidation of Cr. With the analysis methods used in the present study, it could not unequivocally be derived, whether the Cr-rich band consists of pure chromia or e.g. a chromium-rich (Fe,Cr)3 O4 . Qualitatively, these features in respect to the temperature dependence of the oxidation behavior of C related to differences in diffusion rates of Cr in the alloy can also be observed in case of the 12% Cr steel X20. After exposure at 550 C the inner scale shows only very limited Cr enrichment compared to the alloy matrix (Fig. 10). The Cr-enriched regions virtually exclusively consist of oxidized carbides, thus expressing the original carbide morphology in the steel. After exposure at 625 C the scale is still not ideally protective, but its thickness is smaller than that formed at 550 C (Fig. 11a and b). This is correlated with the clear differences in Cr-distribution in the inner scale. The 625 C specimen exhibits Cr-rich stringers in the inner oxide layer, indicating more rapid Cr diffusion from the bulk alloy in direction of the scale (Fig. 12). The enhanced Cr diffusion allows the formation of a Cr-rich (Fe,Cr)3 O4 layer at the scale alloy interface. This preferential oxidation of Cr results in a Cr
Fig. 10. Secondary electron image (a) and X-ray maps showing element distribution in scale formed on 12% Cr steel X20 after 1000 h exposure at 550 C: (b) O-distribution, (c) Cr-distribution and (d) Fedistribution (compare Fig. 11a).
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Fig. 11. Metallographic cross-sections of 12% Cr steel X20 after 1000 h oxidation in Ar–50%H2 O: (a) 550 C and (b) 625 C.
Fig. 12. Secondary electron image (a) and X-ray maps showing element distribution in scale formed on 12% Cr steel X20 after 1000 h exposure at 625 C: (b) O-distribution, (c) Cr-distribution and (d) Fedistribution (compare Fig. 11b).
depletion beneath the Cr-rich spinel layer in the alloy. After continued Cr depletion, eventually the Cr-rich spinel can no longer remain stable. As a result, Fe-rich oxides are subsequently being formed. Repetition of this process leads to the microstructure in the inner part of the scale shown in Fig. 12.
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3.4. Mechanisms of scale formation for steels with different Cr content Ehlers et al. [8], Ehlers and Quadakkers [15] and Thiele et al. [16] recently showed that commercial ferritic steels with Cr contents between 9% and 12% Cr may exhibit a ‘‘bell-shape temperature dependence’’ of the oxidation rate when exposed in water vapour rich environments at temperatures in the range 600–800 C. It was claimed that this effect is related to the temperature dependence of the diffusion coefficient for scale forming alloying elements, primarily Cr (and possibly Mn, Si, V), in the alloy and the solubility of Cr-rich phases (carbides) in the steel matrix [12]. In the present investigation it was found that during exposure in Ar–50%H2 O such an inverse temperature dependence can also occur if only the temperature range of 550–650 C, i.e. the envisaged application temperature range of the mentioned steels, is considered. Besides, it was found that for some of the steels the temperature dependence can be more complex than recently described for the temperature range 600–800 C. For rationalizing the present results, we first consider the scale formation of ‘‘lowCr’’ steels (Cr contents of, approximately 9% or less). At 550 C the diffusion of Cr (and the other scale forming elements) in the alloy is so slow [13], that incorporation in the scale mainly occurs by embedding of Cr-rich carbides in the inner part of the scale due to the inwardly moving scale/alloy interface. In other words, hardly any selective oxidation of Cr occurs. In the inner scale, the oxidized carbides prevail as particles of (Fe,Cr)3 O4 spinel and the distribution of these Fe/Cr-spinel precipitates exhibits practically a ‘‘fingerprint’’ of the carbide distribution in the steel. Reduction of the scale growth rate by the presence of Cr compared to that of a Cr-free steel is at that temperature practically only due to a reduction of the actually available area for cation diffusion by the embedded Cr-rich spinel particles in the inner part of the scale. Variation in Cr content in these ‘‘low-Cr’’ steels (example 1 in Fig. 13) will thus
650°C
log Kp
600°C
550°C
5
650°C 550°C
6
7
8
9
1
2
10
11
3
600°C
4
12
13
14
15
16
Cr content (Mass.-%)
Fig. 13. Schematic illustration of oxidation rate as function of Cr content in the temperature range 550– 650 C. Numbers 1–5 illustrate alloy compositions, which exhibit different types of temperature dependence of the oxidation rates.
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Fig. 14. Area near the scale/alloy interface on 10% Cr steel B after 1000 h exposure at 650 C in Ar– 50%H2 O.
have only little effect on the oxidation rate. The part of the scale, which is in direct contact with the steel, frequently consists of an inhomogeneous mixture of wustite [15,17] and Cr-rich precipitates of chromia and Cr-rich Fe/Cr-spinel (Fig. 4c). The rate-determining step of the scaling process is the diffusion in the scale or the kinetics of a surface reaction. An increase in temperature from 550 to 600 or 650 C will for these ‘‘low-Cr’’ steels thus lead to an increase in the oxidation rate. However, as diffusion in the alloy becomes enhanced, the tendency for Cr to selectively oxidize becomes more pronounced than at 550 C. This results in the formation of Cr-rich stringers of Fe/Crspinel in the inner part of the scale. The formation and morphology of these stringers (Fig. 14) can be explained in the following way: The more selective oxidation of Cr tends to result in formation of a protective Crrich spinel layer at the scale/alloy interface [15]. This leads, however, to Cr depletion in the underlying alloy matrix. After the Cr content in this region has decreased beneath a critical level, the spinel formation can no longer be sustained. Then further oxidation proceeds via magnetite (and/or wustite) formation. As a result, the Cr depletion gradually vanishes until eventually formation of the Cr-rich spinel becomes possible again. This repetitive process leads to the typical morphology of the inner scale showing Cr-rich spinel stringers embedded in an Fe3 O4 ‘‘matrix’’ with inner scale as e.g. illustrated in Fig. 14. The number and/or size of the Cr-rich stringers in the inner scale will increase with increasing Cr content. This will result in a smaller volume fraction of magnetite in which Fe cations can rapidly diffuse and thus the oxidation rate will be more dependent on Cr content than at 550 C. However, the Cr concentration in the alloy is still not sufficiently high to allow the formation of a protective Cr-rich layer, e.g. of FeCr2 O4 , at the scale/alloy interface which is stable during long term exposure. The rate-determining step for the scale growth process will be either the transport of Fe cations through the magnetite or the kinetics of a surface reaction. At the higher test
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temperatures (600–650 C) the latter is more likely to be the case as illustrated by the fact that during quite long exposure times, the outer scale does not contain hematite (Fig. 4b), i.e. the oxide phase which would exist in thermodynamic equilibrium with the gas atmosphere [15]. The oxidation rate will thus increase with increasing temperature. In the case of ‘‘high-Cr’’ steels (example 5 in Fig. 13), the Cr content is sufficiently high to stabilize a protective, Cr-rich oxide scale. The oxidation rate of the very slowly growing scale will be determined by transport processes in the scale and thus also for this type of steels the oxidation rate will increase with increasing temperature. It has been shown previously [15], that the critical Cr concentration at which the change from protective to non-protective oxidation occurs, depends on temperature. The critical Cr content above which a ‘‘drop in Kp ’’ occurs, decreases with increasing temperature [15]. This was claimed to be mainly caused by the temperature dependence of Cr diffusion and the carbide solubility in the steel [15]. Possibly, a further aspect is the enhanced reaction rate between the inner wustite layer with the Cr-rich internal precipitates to form Fe/Cr-spinel. Considering these facts in combination with the scale formation mechanisms described above for ‘‘low-Cr’’ and ‘‘high-Cr’’ steels, the oxidation rate as function of protective-scale forming alloying elements at different temperatures must exhibit a shape as shown in Fig. 13. This schematic figure at least qualitatively explains, that steels with a Cr content in the range of approximately 10–12% can exhibit various types of temperature dependence of the oxidation rate, as actually experimentally observed (Fig. 2). The scaling rate of the ‘‘intermediate-Cr’’ steels designated by the numbers 2–4 in Fig. 13 is determined by the repetitive process of stringers of Cr-rich spinel and subsequent magnetite formation, described before. A temperature increase in the range between 550 and 650 C can lead to an increase as well as a decrease of the oxidation rate, depending on the exact Cr content and the actual temperature step. It is important to mention, that in Fig. 13 the only ‘‘protective-scale forming element’’ considered is Cr. Of course, the commercial steels in reality commonly contain a large number of additional alloying elements such as Si, Mn, V, C, N, Mo, W and/or Co. The four first mentioned elements are strong oxide formers and can thus substantially shift the location at which the drop in oxidation rate in Fig. 13 occurs. From previous studies [14,15] it may be derived that Si shifts the ‘‘Kp -drop’’ to lower Cr contents, whereas Mn has the opposite effect. Co additions also shift the ‘‘Kp -drop’’ to lower Cr contents whereby this Co effect seems to be stronger at 650 C than at lower temperatures [15]. Co is not incorporated into the protective surface scale and thus its positive effect possibly relates to an increase in the Cr activity, the diffusivity of Cr in the alloy and/or the carbide solubility. A complicating factor in the interpretation of the oxidation behavior of the discussed type of materials in steam is that the oxidation mechanism can change with time. Steels with a composition near to the boundary between protective and nonprotective behavior can initially behave like a ‘‘low-Cr’’ steel, and subsequently exhibit formation of a protective layer at or near the scale/steel interface as e.g. observed in the present study for steel C during exposure at 600 C (Figs. 1c and 9).
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The formation of this protective scale at the oxide/alloy interface results in hampered transport of Fe cations into the scale and thus in void and gap formation due to condensation of inwardly moving vacancies. As a result, the initially formed outer magnetite layer transforms into hematite and the scale becomes very prone to spallation. This explains the observation frequently made in practice, that in the temperature range 550–650 C, surface oxides formed on 10–12% Cr steels can in some cases be more prone to spallation than e.g. those formed on a 1% Cr steel. The anomalous temperature dependence of the oxidation behavior of the 10–12% Cr steels has far-reaching implications for the evaluation of test results and their extrapolation to longer times and other temperatures. Oxidation rates determined in accelerated tests at high temperature (for instance: 650 C) cannot simply be used to predict materials behavior at lower temperatures. Also the opposite is true: relatively high oxidation rates at low temperatures (550 C) do not necessarily mean that the steel will exhibit high oxidation rates at 650 C.
4. Conclusions In steam environments some of the newly developed ferritic 9–12% Cr steels tend to exhibit an anomalous temperature dependence of the oxidation behavior in the temperature range 550–650 C. That means, that under these conditions the oxidation rates of a number of the new steels do not show the common increase of the scaling rate with increasing temperature. Depending on steel type, the maximum oxidation rates may be observed at lower temperatures (e.g. 550 or 600 C). The reasons for this ‘‘kinetics inversion’’ behavior are related to two concurring processes occurring upon temperature increase: • enhanced in-scale diffusion and surface reaction kinetics, resulting in an increase of the oxidation rate, • enhanced incorporating of Cr in the scale, resulting in a decrease of the oxidation rate. The temperature dependence can substantially be affected by commonly present alloying elements, which may modify the scale composition, the Cr diffusion and/or the carbide solubility in the alloy. The anomalous oxidation behavior of the 10–12% Cr steels has far-reaching implications for the evaluation of test results and their extrapolation to longer times and higher or lower temperatures. References [1] K. Weizierl, Kohlekraftwerke der Zukunft, VGB Kraftwerkstechnik 74 (2) (1994). [2] R. Blum, J. Hald, W. Bendick, A. Rosselet, J.C. Vaillant, Neuentwicklungen hochwarmfester ferritisch-martensitischer St€ahle aus den USA, Japan und Europa, VGB Kraftwerkstechnik 74 (8) (1994).
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[3] J.B. Mariott, K. Pithan, A review of the properties of 9–12% Cr steels for use as HP/IP rotors in advanced steam turbines, Office for Official Publications of the European Communities, Catalog number: CD-NA-1187-EN-C, 1989. [4] S. Jianian, Z. Longjiang, L. Tiefan, Oxidation of Metals 48 (3/4) (1997). [5] P. Mayer, A.V. Manolescu, The role of structural and compositional factors in corrosion of ferritic steels in steam, high temperature corrosion, National Association of Corrosion Engineers (1983) 368. [6] A. Rahmel, J. Tobolski, Einfluss von Wasserdampf und Kohlendioxid auf die Oxidation von Eisen in Sauerstoff bei Hohen Temperaturen, Corrosion Science 5 (1965) 333. [7] Y. Ikeda, K. Nii, Mechanism of accelerated oxidation of Fe–Cr alloys in water vapour containing atmosphere, Transaction of the Japanese Institute of Metals 26 (1984) 52. [8] R.J. Ehlers, E.J. Smaardijk, H.J. Penkala, A.K. Tyagi, L.Singheiser, W.J. Quadakkers, Effect of steel composition on the bell-shape temperature dependence of oxidation in water vapour containing environments, in: Proceedings of the International Corrosion Congress, Cape Town, 27.09.1999– 01.10.1999, Proceedings, Paper 336. [9] R.J. Ehlers, P.J. Ennis, L. Singheiser, W.J. Quadakkers, T. Link, Significance of scale spalling for the lifetime of ferritic 9–10% Cr steels during oxidation in water vapour at temperatures between 550 and 650 C, in: Proceedings of the European Federation of Corrosion Workshop 2001, No. 34. [10] I. Armitt, D.P. Holmes, The spalling of steam grown oxide from superheater and reheater tube steels, Technical Planning Study 76-655, Final Report, Feb. 1978 EPRI, 3412 Hillview Ave., P.O. Box 10412, Palo Alto, CA 94303, USA, 1978. [11] P. Kofstad, High Temperature Corrosion, Elsevier Applied Science Publishers Ltd, 1988, p. 257. [12] R.N. Durham, B. Gleeson, D.J. Young, Chromium carbide dissolution during oxidation of chromia forming alloys, in: Proceedings of the International Corrosion Congress, Cape Town, 27.09.1999– 01.10.1999, Proceedings, Paper 178. [13] C. Piehl, Z. T^ okei, H.J. Grabke, The role of fast diffusion paths in the selective oxidation of chromium steels, Defect and Diffusion Forum 194–199 (2001) 1689–1694, Oxidation of Metals 37 (1/2) (1992). [14] H.E. Townsend, The effects of alloying elements on the corrosion of steel in industrial atmospheres, in: Proceedings of the International Corrosion Congress, Cape Town, 27.09.1999–01.10.1999, Proceedings, Paper 243. [15] R.J. Ehlers, W.J. Quadakkers, Oxidation von ferritischen 9–12% Cr-St€ahlen in wasserdampfhaltigen Atmosph€ aren bei 550–650 C, Report Forschungszentrum J€ ulich, J€ ul-3883, ISSN 0944-2952, June 2001. [16] M. Thiele, W.J. Quadakkers, F. Schubert, H. Nickel, Hochtemperaturkorrosion von ferritischen und austenitischen St€ahlen in simulierten Rauchgasen kohlbefeuerter Kraftwerke, Report Forschungszentrum J€ ulich, J€ ul.-3712, ISSN 0944-2952, J€ ulich, FRG, 1999, p. 63. [17] M. Thiele, H. Teichmann, W. Schwarz, W.J. Quadakkers, Korrosionsverhalten von ferritischen und austenitischen St€ahlen in simulierten rauchgasen von stein und braunkohlebefeuerten Kraftwerken, VGB Kraftwerkstechnik 77 (2) (1997) 129–134.
Anomalous temperature dependence of oxidation kinetics during steam oxidation of ferritic steels in the temperature range 550–650 C a,* _ J. Zurek , E. Wessel a, L. Niewolak a, F. Schmitz b, T.-U. Kern b, L. Singheiser a, W.J. Quadakkers a a
Forschungszentrum J€ulich GmbH, IWV-2, 52425 J€ulich, FRG b Siemens Power Generation, M€ulheim, FRG Received 6 August 2003; accepted 16 January 2004 Available online 12 March 2004
Abstract The oxidation behavior of a number of selected ferritic steels in a simulated steam environment at temperatures between 550 and 650 C was studied. In the prevailing test gas, some of the studied 9–12% Cr steels tended to exhibit an anomalous temperature dependence of the oxidation behavior. This means, that the oxidation rates do not steadily increase with increasing temperature. At higher temperatures, some of the studied steels tend to form a very thin and protective oxide scale whereas at lower temperature rapidly growing, less-protective oxides are being developed. The anomalous temperature dependence is related to differences in chromium distribution in the inner part of the oxide scale. The effect is observed for steels with intermediate-Cr contents (10–12%) whereas steels with either lower or higher Cr contents exhibit an increasing oxidation rate with increasing temperature. 2004 Elsevier Ltd. All rights reserved. Keywords: A. Steel; B. SIMS; C. High temperature corrosion; Selective oxidation
*
Corresponding author. Tel.: +48-2461-61-55-22; fax: +48-2461-61-36-99. _ E-mail address: [email protected] (J. Zurek).
0010-938X/$ - see front matter 2004 Elsevier Ltd. All rights reserved. doi:10.1016/j.corsci.2004.01.010
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1. Introduction The need for reduction of CO2 emission and fuel consumption is a leading factor for the requirement of the utilities to increase the thermal efficiency of fossil fired power generation plants [1]. An increase in the steam temperature from common 535 to 650 C and the steam pressure from 185 to 300 bar would enable a reduction in fuel consumption and a reduction of CO2 emission by more than 25% [2]. A number of high strength 9–12% Cr steels have been developed for application as construction materials in such advanced power plants [3]. Whereas all these steels possess excellent oxidation resistance during service in oxygen or air [4], it is well known that in water vapour containing environments, as prevailing under power plant service conditions, the oxidation rates of a number of these steels are significantly enhanced [5–7]. Recently it was shown that during short term isothermal exposure in water vapour-rich environments, some of these steels tended to exhibit a bell-shape temperature dependence of the oxidation rate in the temperature range 600–800 C [8]. Whereas poorly protective oxides were formed at 600 and 650 C, very thin chromium-rich oxides were found at 700 C and especially 800 C, resulting in very low oxidation rates. In the present study it will be shown, that such an anomalous temperature dependence can also occur during exposure in steam if only lower, i.e. more practically relevant temperatures for current power plant application (550–650 C), are being considered. It will also be illustrated, that the effect does not only occur in laboratory tests with short exposure times of up to typically 100 h [8], but it is also of great significance for very long time service in practical application.
2. Experimental The compositions of the studied materials, delivered in form of forged bars, are listed in Table 1. The steels varied in Cr content between 1.3% and 12.4%. One of the steels contained a substantial addition of Co. Specimens with a size of 20 · 10 · 2 mm were machined. For the oxidation studies, the specimens were ground to 800 grit surface finish and exposed to an Ar + 50%H2 O atmosphere. The exposures were carried out for up to 1000 h in the temperature range 550–650 C, whereby the specimens were cooled to room temperature every 250 h for weight measurements. After exposure the oxidation products were characterised by optical metallography, scanning electron microscopy (SEM) with energy dispersive X-ray analysis (EDX) Table 1 Composition of the investigated steels (mass %) Steeldesignation
C
Cr
Mo
Si
Mn
Co
V
W
Ni
Cu
A B C X20
0.17 0.10 0.07 0.20
1.31 10.4 10.8 12.4
0.95 1.05 0.86 1.00
0.51 0.09 0.3 0.26
0.59 0.45 0.62 1.5
– 0.012 6.52 0.03
0.21 0.18 0.27 0.054
– 0.98 – 0.06
0.11 0.73 0.72 0.51
0.07 0.028 0.006 0.1
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and sputtered neutrals mass spectrometry (SNMS). Before mounting for metallographic cross-section analyses, the specimens were sputtered by a thin gold layer and subsequently electroplated by nickel. This coating provided protection of the surface oxide layer during grinding and polishing and it ensured a better optical contrast between oxide and mounting material. 3. Results and discussion 3.1. Different types of temperature dependence Fig. 1(a)–(d) shows the weight change data of the four studied steels during exposure in Ar + 50%H2 O up to 1000 h at different temperatures. In the following the weight change data are given as relative data. The 1% Cr steel is commonly used in practice at temperatures up to 550 C. Therefore all the measured weight change data are normalised to the 1000 h value of the 1% Cr steel (steel A) at 550 C. The low-Cr steel A and the Co-free 10% Cr steel B clearly exhibit increasing oxidation rates with increasing temperature. However, the two other steels show a much more complex temperature dependence of the oxidation rates. After 1000 h of exposure the highest weight change in case of X20 is found for the temperature 550 C. For this steel the weight change drops upon temperature change to 600 C and again increases if the temperature is raised to 625 or 650 C. The Co-containing steel C again 7 625˚C
(a) 8 6
600˚C
4 2
550˚C
0
Relative weight change
Relative weight change
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6
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550˚C
2 1 0
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1000
0
200
400
1.6
800
1000
1.6 550˚C
(c)
1.4 1.2
600˚C
1 0.8 0.6 0.4
625˚C
0.2
650˚C
0 0
200
400
600
Time (h)
800
1000
Relative weight change
Relative weight change
600
Time (h)
Time (h)
550˚C
(d)
1.4 1.2 1
625˚C
0.8 650˚C
0.6 0.4 0.2 0
600˚C
0
200
400
600
800
1000
Time (h)
Fig. 1. Weight change of various steels during exposure at temperatures between 550 and 650 C in Ar–50%H2 O; (a) 1% Cr steel A, (b) 10% Cr steel B, (c) Co-containing 10% Cr steel C and (d) 12% Cr steel X20.
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2304 8
A
6
B
4
2
C X20
0 540 550 560 570 580 590 600 610 620 630 640 650 660 Temperature (°C)
Fig. 2. Temperature dependence of weight change of the various steels after oxidation in Ar–50%H2 O for 250 h, relative to steel A at 550 C.
exhibits a different type of temperature dependence. After 1000 h exposure at 625 and 650 C the weight changes are clearly lower than those observed during 550 C exposure, whereby the 600 C-value is only slightly smaller than that at 550 C. It is important to note that the high value at 600 C is already reached after 250 h, in other words, the oxidation rate between 250 and 1000 h is similarly small as that observed at 625 and 650 C. Fig. 2 summarizes the weight change data for the four steels after 250 h exposure at the various temperatures. 3.2. Scale morphology in case of normal temperature dependence Fig. 3 shows typical microstructures of oxide scales at the 1% Cr steel A after 1000 h exposure at 550 and 625 C. The scales nearly exclusively consist of magnetite (Fe3 O4 ). In agreement with the gravimetric data, the scale formed at 625 C is substantially thicker than that formed at 550 C. After the 625 C exposure, a thin
Fig. 3. Metallographic cross-sections of 1% Cr steel A showing oxide scales after 1000 h oxidation in Ar–50%H2 O: (a) 550 C and (b) 625 C.
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zone of wustite (FeO) with minor Cr2 O3 precipitates is found near the oxide/alloy interface. The specimen exposed at 550 C exhibits small amounts of hematite (Fe2 O3 ) near the scale/gas interface. In both scales substantial void formation is observed. After the 625 C exposure the voids are occurring in the outer magnetite layer and near the scale/alloy interface. During exposure at 550 C, they concentrate at the interface between inner and outer scale [9], resulting in a tendency to spalling of the outer part of the oxide layer. Fig. 4 shows the microstructures of the scales on the Co-free 10% Cr steel B after 1000 h exposure at 550 and 650 C. Very similar morphologies were found after exposure at 600 and 625 C. The scales typically consist of an outer Fe3 O4 layer and an inner layer containing (Fe,Cr)3 O4 stringers in an Fe3 O4 ‘‘matrix’’ [9]. The two layers are separated by a near-ideally straight line representing the original metal surface. Internal oxide stringers embedded in an FeO-matrix are observed near the scale/alloy interface. After exposure at 550 C the outer scale exhibits substantial void formation near the interface with the inner scale. Additional voids, however of smaller size, are formed near the interface between inner scale and internal oxidation zone. After 650 C exposure this void formation in the inner zone is more pronounced. The few number of voids in the outer scale are present throughout the Fe3 O4 layer. The SEM-picture in Fig. 5 shows, that in addition to the voids which
Fig. 4. Metallographic cross-sections of 10% Cr steel B showing oxide scales after 1000 h oxidation in Ar–50%H2 O: (a) 550 C, (b) 650 C and (c) 550 C––high magnification.
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Fig. 5. High magnification of area near the scale/alloy interface on 10% Cr steel B after 1000 h exposure at 650 C in Ar–50%H2 O (compare Fig. 4b).
can be easily distinguished by light microscopy (Fig. 4), the inner scale contains an infinite by large number of microvoids, next to the precipitates of (Fe,Cr)3 O4 . The transverse cracks visible in the cross-section in Fig. 4b are likely to be related to stresses originating from differences in thermal expansion coefficient between alloy and oxide at high temperatures, the thermal expansion coefficient of magnetite being larger than that of the steel [10], resulting in tensile stresses in the oxide during cooling. In some of the cases, hematite is formed in the scale near the interface with the gas (Fig. 3a). Comparison of the equilibrium oxygen partial pressure in the test gas (approximately 10 8 bar at 650 C) with the Fe2 O3 /Fe3 O4 -equilibrium oxygen partial pressure (approximately 10 13 bar at 650 C) clearly reveals that hematite is the iron oxide phase, which is in equilibrium with the gas atmosphere. In most cases, however, no hematite was found in the outer scale. Apparently no equilibrium between gas and oxide surface has been established during the time of exposure. The ratedetermining step of the scale growth process is in that case not the cation and/or anion transport in the scale, but probably the kinetics of the gas/solid reactions at the scale/gas interface. One might speculate about possible reasons for the formation of the large and tiny voids observed in the scales (Fig. 5). An explanation might be, that they originate from condensation of cation vacancies, if one assumes the growth in magnetite and wustite to primarily occur via Fe cations [11]. In the inner part of the scale, the interfaces between the Fe3 O4 ‘‘matrix’’ and the finely distributed precipitates of (Fe,Cr,Mn)O4 , which were shown to partly originate from the oxidation of carbides [9], could supply numerous vacancy condensation sites resulting in an extremely large number of tiny voids (Fig. 5). As the outer part of the scale consists of virtually pure magnetite, the number of vacancy condensation sites in this part of the scale is extremely small, resulting in fewer but larger voids (Fig. 4).
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Summarizing it can be said that the composition and morphology of the scales formed on steels A and B does not fundamentally differ when exposed at different temperatures. As the rates of anion/cation transport in the scales as well as the rates for the surface reaction are expected to increase with increasing temperature, it is obvious, that the oxidation rates should increase with increasing temperature, as shown in Fig. 1a and b. 3.3. Steels showing anomalous temperature dependence Fig. 6 shows the oxide scales formed on the Co-containing 10% Cr steel C after 1000 h exposure at different temperatures. In agreement with the weight change data in Fig. 1c, the thicknesses of the scales formed at 550 and 600 C are quite similar. However, the scale morphologies show substantial differences. After exposure at 550 C the inner part of the scale exhibits a wide zone of internal oxide precipitates in an FeO-matrix. The outer scale consist of Fe3 O4 plus Fe2 O3 , whereby the latter is not only present at the scale/gas interface but also locally, more inward, i.e. near the large gap in the outer scale. This observation allows the following conclusions:
Fig. 6. Metallographic cross-sections of Co-containing 10% Cr steel C after 1000 h oxidation in Ar– 50%H2 O: (a) 550 C, (b) 600 C and (c) 625 C.
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• The outer surface of the scale is in (near-) equilibrium with the gas (in contrary to the case described in the previous section). • The outer scale allows molecular gas transport. Also after 600 C exposure, the outer scale contains minor amounts of hematite. The gap in the scale is located near the boundary between inner and outer scale i.e. the original metal surface. The most striking difference with the 550 C specimen is the virtual absence of the internal oxidation zone consisting of FeO + Cr2 O3 . In agreement with the weight change data, an increase in temperature to 625 C again leads to a dramatic change in scale morphology (Fig. 6c). A very protective oxide is formed which, according to SNMS-analyses, mainly consists of an inner Cr2 O3 -layer and minor amounts of Fe2 O3 in the outer part of the scale (Fig. 7). No indication was found for presence of Co in the surface scale. A very similar scale morphology was found after the 650 C exposure. Based on these observations, in combination with findings reported in literature [8], the temperature dependence observed for steel C can be explained in the following way: At 550 C the scale initially grows via cation diffusion in the magnetite and wustite lattice. Chromium diffusion in the alloy and in the oxide is so slow that it can be considered as being virtually immobile. It becomes incorporated in the inner part of the scale due to oxidation of the Cr-rich carbides prevailing in the alloy, i.e. the carbides become embedded by the inwardly moving oxidation front [9]. The Cr-rich stringers in the inner scale thus exhibit a sort of ‘‘fingerprint’’ of the typical carbide
100
90
80
Concentration (at%)
Fe 70
O 60
50
40
Cr
30
20
10
Co
Mn 0 0
500
1000
1500
2000
2500
3000
3500
4000
4500
5000
Sputtering Time (s)
Fig. 7. SNMS depth profile of oxide scale on Co-containing 10% Cr steel C after 1000 h oxidation at 625 C in Ar–50%H2 O (compare Fig. 6c).
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Fig. 8. Secondary electron image (a) and X-ray maps showing element distribution in scale formed on Cocontaining 10% Cr steel C after 1000 h exposure at 550 C: (b) O-distribution, (c) Cr-distribution, (d) Fedistribution and (e) Co-distribution (compare Fig. 6a).
morphology in this type of steels. This is confirmed by the element distribution in Fig. 8, showing that in the scale the Cr is hardly enriched compared to the alloy Cr concentration. At 625 and 650 C, Cr diffusion in the alloy in direction of the scale/alloy interface is sufficiently fast to stabilize a chromia based scale, with remnants of transient Feoxides (Fig. 6c). Probably, the partial dissolution of the Cr-rich carbides in the alloy matrix at the higher temperatures contributes to the faster Cr-transport in the alloy.
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The scale formation at 600 C is ‘‘intermediate’’ between the 550 and 625 C case. Based on the initial, very rapid weight change, (Fig. 1c) the scale formation mechanism at 600 C is more similar to that at 550 C than that at 625 or 650 C. However, after longer times the oxidation rate resembles that observed at the higher temperatures. The latter would mean, that during the 600 C exposure somewhere in the scale a protective Cr-rich oxide becomes being stabilized. This is confirmed by
Fig. 9. Secondary electron image (a) and X-ray maps showing element distribution in scale formed on Cocontaining 10% Cr steel C after 1000 h exposure at 600 C: (b) O-distribution, (c) Cr-distribution, (d) Fedistribution and (e) Co-distribution (compare Fig. 6b).
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the element mapping in Fig. 9. Whereas the overall Cr enrichment in the inner part of the scale is not very pronounced if compared with the Cr alloy concentration, a strong enrichment is visible in a narrow band near the scale/alloy interface. Beneath this, a Cr depletion is clearly visible in the alloy as a result of the selective oxidation of Cr. With the analysis methods used in the present study, it could not unequivocally be derived, whether the Cr-rich band consists of pure chromia or e.g. a chromium-rich (Fe,Cr)3 O4 . Qualitatively, these features in respect to the temperature dependence of the oxidation behavior of C related to differences in diffusion rates of Cr in the alloy can also be observed in case of the 12% Cr steel X20. After exposure at 550 C the inner scale shows only very limited Cr enrichment compared to the alloy matrix (Fig. 10). The Cr-enriched regions virtually exclusively consist of oxidized carbides, thus expressing the original carbide morphology in the steel. After exposure at 625 C the scale is still not ideally protective, but its thickness is smaller than that formed at 550 C (Fig. 11a and b). This is correlated with the clear differences in Cr-distribution in the inner scale. The 625 C specimen exhibits Cr-rich stringers in the inner oxide layer, indicating more rapid Cr diffusion from the bulk alloy in direction of the scale (Fig. 12). The enhanced Cr diffusion allows the formation of a Cr-rich (Fe,Cr)3 O4 layer at the scale alloy interface. This preferential oxidation of Cr results in a Cr
Fig. 10. Secondary electron image (a) and X-ray maps showing element distribution in scale formed on 12% Cr steel X20 after 1000 h exposure at 550 C: (b) O-distribution, (c) Cr-distribution and (d) Fedistribution (compare Fig. 11a).
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Fig. 11. Metallographic cross-sections of 12% Cr steel X20 after 1000 h oxidation in Ar–50%H2 O: (a) 550 C and (b) 625 C.
Fig. 12. Secondary electron image (a) and X-ray maps showing element distribution in scale formed on 12% Cr steel X20 after 1000 h exposure at 625 C: (b) O-distribution, (c) Cr-distribution and (d) Fedistribution (compare Fig. 11b).
depletion beneath the Cr-rich spinel layer in the alloy. After continued Cr depletion, eventually the Cr-rich spinel can no longer remain stable. As a result, Fe-rich oxides are subsequently being formed. Repetition of this process leads to the microstructure in the inner part of the scale shown in Fig. 12.
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3.4. Mechanisms of scale formation for steels with different Cr content Ehlers et al. [8], Ehlers and Quadakkers [15] and Thiele et al. [16] recently showed that commercial ferritic steels with Cr contents between 9% and 12% Cr may exhibit a ‘‘bell-shape temperature dependence’’ of the oxidation rate when exposed in water vapour rich environments at temperatures in the range 600–800 C. It was claimed that this effect is related to the temperature dependence of the diffusion coefficient for scale forming alloying elements, primarily Cr (and possibly Mn, Si, V), in the alloy and the solubility of Cr-rich phases (carbides) in the steel matrix [12]. In the present investigation it was found that during exposure in Ar–50%H2 O such an inverse temperature dependence can also occur if only the temperature range of 550–650 C, i.e. the envisaged application temperature range of the mentioned steels, is considered. Besides, it was found that for some of the steels the temperature dependence can be more complex than recently described for the temperature range 600–800 C. For rationalizing the present results, we first consider the scale formation of ‘‘lowCr’’ steels (Cr contents of, approximately 9% or less). At 550 C the diffusion of Cr (and the other scale forming elements) in the alloy is so slow [13], that incorporation in the scale mainly occurs by embedding of Cr-rich carbides in the inner part of the scale due to the inwardly moving scale/alloy interface. In other words, hardly any selective oxidation of Cr occurs. In the inner scale, the oxidized carbides prevail as particles of (Fe,Cr)3 O4 spinel and the distribution of these Fe/Cr-spinel precipitates exhibits practically a ‘‘fingerprint’’ of the carbide distribution in the steel. Reduction of the scale growth rate by the presence of Cr compared to that of a Cr-free steel is at that temperature practically only due to a reduction of the actually available area for cation diffusion by the embedded Cr-rich spinel particles in the inner part of the scale. Variation in Cr content in these ‘‘low-Cr’’ steels (example 1 in Fig. 13) will thus
650°C
log Kp
600°C
550°C
5
650°C 550°C
6
7
8
9
1
2
10
11
3
600°C
4
12
13
14
15
16
Cr content (Mass.-%)
Fig. 13. Schematic illustration of oxidation rate as function of Cr content in the temperature range 550– 650 C. Numbers 1–5 illustrate alloy compositions, which exhibit different types of temperature dependence of the oxidation rates.
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Fig. 14. Area near the scale/alloy interface on 10% Cr steel B after 1000 h exposure at 650 C in Ar– 50%H2 O.
have only little effect on the oxidation rate. The part of the scale, which is in direct contact with the steel, frequently consists of an inhomogeneous mixture of wustite [15,17] and Cr-rich precipitates of chromia and Cr-rich Fe/Cr-spinel (Fig. 4c). The rate-determining step of the scaling process is the diffusion in the scale or the kinetics of a surface reaction. An increase in temperature from 550 to 600 or 650 C will for these ‘‘low-Cr’’ steels thus lead to an increase in the oxidation rate. However, as diffusion in the alloy becomes enhanced, the tendency for Cr to selectively oxidize becomes more pronounced than at 550 C. This results in the formation of Cr-rich stringers of Fe/Crspinel in the inner part of the scale. The formation and morphology of these stringers (Fig. 14) can be explained in the following way: The more selective oxidation of Cr tends to result in formation of a protective Crrich spinel layer at the scale/alloy interface [15]. This leads, however, to Cr depletion in the underlying alloy matrix. After the Cr content in this region has decreased beneath a critical level, the spinel formation can no longer be sustained. Then further oxidation proceeds via magnetite (and/or wustite) formation. As a result, the Cr depletion gradually vanishes until eventually formation of the Cr-rich spinel becomes possible again. This repetitive process leads to the typical morphology of the inner scale showing Cr-rich spinel stringers embedded in an Fe3 O4 ‘‘matrix’’ with inner scale as e.g. illustrated in Fig. 14. The number and/or size of the Cr-rich stringers in the inner scale will increase with increasing Cr content. This will result in a smaller volume fraction of magnetite in which Fe cations can rapidly diffuse and thus the oxidation rate will be more dependent on Cr content than at 550 C. However, the Cr concentration in the alloy is still not sufficiently high to allow the formation of a protective Cr-rich layer, e.g. of FeCr2 O4 , at the scale/alloy interface which is stable during long term exposure. The rate-determining step for the scale growth process will be either the transport of Fe cations through the magnetite or the kinetics of a surface reaction. At the higher test
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temperatures (600–650 C) the latter is more likely to be the case as illustrated by the fact that during quite long exposure times, the outer scale does not contain hematite (Fig. 4b), i.e. the oxide phase which would exist in thermodynamic equilibrium with the gas atmosphere [15]. The oxidation rate will thus increase with increasing temperature. In the case of ‘‘high-Cr’’ steels (example 5 in Fig. 13), the Cr content is sufficiently high to stabilize a protective, Cr-rich oxide scale. The oxidation rate of the very slowly growing scale will be determined by transport processes in the scale and thus also for this type of steels the oxidation rate will increase with increasing temperature. It has been shown previously [15], that the critical Cr concentration at which the change from protective to non-protective oxidation occurs, depends on temperature. The critical Cr content above which a ‘‘drop in Kp ’’ occurs, decreases with increasing temperature [15]. This was claimed to be mainly caused by the temperature dependence of Cr diffusion and the carbide solubility in the steel [15]. Possibly, a further aspect is the enhanced reaction rate between the inner wustite layer with the Cr-rich internal precipitates to form Fe/Cr-spinel. Considering these facts in combination with the scale formation mechanisms described above for ‘‘low-Cr’’ and ‘‘high-Cr’’ steels, the oxidation rate as function of protective-scale forming alloying elements at different temperatures must exhibit a shape as shown in Fig. 13. This schematic figure at least qualitatively explains, that steels with a Cr content in the range of approximately 10–12% can exhibit various types of temperature dependence of the oxidation rate, as actually experimentally observed (Fig. 2). The scaling rate of the ‘‘intermediate-Cr’’ steels designated by the numbers 2–4 in Fig. 13 is determined by the repetitive process of stringers of Cr-rich spinel and subsequent magnetite formation, described before. A temperature increase in the range between 550 and 650 C can lead to an increase as well as a decrease of the oxidation rate, depending on the exact Cr content and the actual temperature step. It is important to mention, that in Fig. 13 the only ‘‘protective-scale forming element’’ considered is Cr. Of course, the commercial steels in reality commonly contain a large number of additional alloying elements such as Si, Mn, V, C, N, Mo, W and/or Co. The four first mentioned elements are strong oxide formers and can thus substantially shift the location at which the drop in oxidation rate in Fig. 13 occurs. From previous studies [14,15] it may be derived that Si shifts the ‘‘Kp -drop’’ to lower Cr contents, whereas Mn has the opposite effect. Co additions also shift the ‘‘Kp -drop’’ to lower Cr contents whereby this Co effect seems to be stronger at 650 C than at lower temperatures [15]. Co is not incorporated into the protective surface scale and thus its positive effect possibly relates to an increase in the Cr activity, the diffusivity of Cr in the alloy and/or the carbide solubility. A complicating factor in the interpretation of the oxidation behavior of the discussed type of materials in steam is that the oxidation mechanism can change with time. Steels with a composition near to the boundary between protective and nonprotective behavior can initially behave like a ‘‘low-Cr’’ steel, and subsequently exhibit formation of a protective layer at or near the scale/steel interface as e.g. observed in the present study for steel C during exposure at 600 C (Figs. 1c and 9).
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The formation of this protective scale at the oxide/alloy interface results in hampered transport of Fe cations into the scale and thus in void and gap formation due to condensation of inwardly moving vacancies. As a result, the initially formed outer magnetite layer transforms into hematite and the scale becomes very prone to spallation. This explains the observation frequently made in practice, that in the temperature range 550–650 C, surface oxides formed on 10–12% Cr steels can in some cases be more prone to spallation than e.g. those formed on a 1% Cr steel. The anomalous temperature dependence of the oxidation behavior of the 10–12% Cr steels has far-reaching implications for the evaluation of test results and their extrapolation to longer times and other temperatures. Oxidation rates determined in accelerated tests at high temperature (for instance: 650 C) cannot simply be used to predict materials behavior at lower temperatures. Also the opposite is true: relatively high oxidation rates at low temperatures (550 C) do not necessarily mean that the steel will exhibit high oxidation rates at 650 C.
4. Conclusions In steam environments some of the newly developed ferritic 9–12% Cr steels tend to exhibit an anomalous temperature dependence of the oxidation behavior in the temperature range 550–650 C. That means, that under these conditions the oxidation rates of a number of the new steels do not show the common increase of the scaling rate with increasing temperature. Depending on steel type, the maximum oxidation rates may be observed at lower temperatures (e.g. 550 or 600 C). The reasons for this ‘‘kinetics inversion’’ behavior are related to two concurring processes occurring upon temperature increase: • enhanced in-scale diffusion and surface reaction kinetics, resulting in an increase of the oxidation rate, • enhanced incorporating of Cr in the scale, resulting in a decrease of the oxidation rate. The temperature dependence can substantially be affected by commonly present alloying elements, which may modify the scale composition, the Cr diffusion and/or the carbide solubility in the alloy. The anomalous oxidation behavior of the 10–12% Cr steels has far-reaching implications for the evaluation of test results and their extrapolation to longer times and higher or lower temperatures. References [1] K. Weizierl, Kohlekraftwerke der Zukunft, VGB Kraftwerkstechnik 74 (2) (1994). [2] R. Blum, J. Hald, W. Bendick, A. Rosselet, J.C. Vaillant, Neuentwicklungen hochwarmfester ferritisch-martensitischer St€ahle aus den USA, Japan und Europa, VGB Kraftwerkstechnik 74 (8) (1994).
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[3] J.B. Mariott, K. Pithan, A review of the properties of 9–12% Cr steels for use as HP/IP rotors in advanced steam turbines, Office for Official Publications of the European Communities, Catalog number: CD-NA-1187-EN-C, 1989. [4] S. Jianian, Z. Longjiang, L. Tiefan, Oxidation of Metals 48 (3/4) (1997). [5] P. Mayer, A.V. Manolescu, The role of structural and compositional factors in corrosion of ferritic steels in steam, high temperature corrosion, National Association of Corrosion Engineers (1983) 368. [6] A. Rahmel, J. Tobolski, Einfluss von Wasserdampf und Kohlendioxid auf die Oxidation von Eisen in Sauerstoff bei Hohen Temperaturen, Corrosion Science 5 (1965) 333. [7] Y. Ikeda, K. Nii, Mechanism of accelerated oxidation of Fe–Cr alloys in water vapour containing atmosphere, Transaction of the Japanese Institute of Metals 26 (1984) 52. [8] R.J. Ehlers, E.J. Smaardijk, H.J. Penkala, A.K. Tyagi, L.Singheiser, W.J. Quadakkers, Effect of steel composition on the bell-shape temperature dependence of oxidation in water vapour containing environments, in: Proceedings of the International Corrosion Congress, Cape Town, 27.09.1999– 01.10.1999, Proceedings, Paper 336. [9] R.J. Ehlers, P.J. Ennis, L. Singheiser, W.J. Quadakkers, T. Link, Significance of scale spalling for the lifetime of ferritic 9–10% Cr steels during oxidation in water vapour at temperatures between 550 and 650 C, in: Proceedings of the European Federation of Corrosion Workshop 2001, No. 34. [10] I. Armitt, D.P. Holmes, The spalling of steam grown oxide from superheater and reheater tube steels, Technical Planning Study 76-655, Final Report, Feb. 1978 EPRI, 3412 Hillview Ave., P.O. Box 10412, Palo Alto, CA 94303, USA, 1978. [11] P. Kofstad, High Temperature Corrosion, Elsevier Applied Science Publishers Ltd, 1988, p. 257. [12] R.N. Durham, B. Gleeson, D.J. Young, Chromium carbide dissolution during oxidation of chromia forming alloys, in: Proceedings of the International Corrosion Congress, Cape Town, 27.09.1999– 01.10.1999, Proceedings, Paper 178. [13] C. Piehl, Z. T^ okei, H.J. Grabke, The role of fast diffusion paths in the selective oxidation of chromium steels, Defect and Diffusion Forum 194–199 (2001) 1689–1694, Oxidation of Metals 37 (1/2) (1992). [14] H.E. Townsend, The effects of alloying elements on the corrosion of steel in industrial atmospheres, in: Proceedings of the International Corrosion Congress, Cape Town, 27.09.1999–01.10.1999, Proceedings, Paper 243. [15] R.J. Ehlers, W.J. Quadakkers, Oxidation von ferritischen 9–12% Cr-St€ahlen in wasserdampfhaltigen Atmosph€ aren bei 550–650 C, Report Forschungszentrum J€ ulich, J€ ul-3883, ISSN 0944-2952, June 2001. [16] M. Thiele, W.J. Quadakkers, F. Schubert, H. Nickel, Hochtemperaturkorrosion von ferritischen und austenitischen St€ahlen in simulierten Rauchgasen kohlbefeuerter Kraftwerke, Report Forschungszentrum J€ ulich, J€ ul.-3712, ISSN 0944-2952, J€ ulich, FRG, 1999, p. 63. [17] M. Thiele, H. Teichmann, W. Schwarz, W.J. Quadakkers, Korrosionsverhalten von ferritischen und austenitischen St€ahlen in simulierten rauchgasen von stein und braunkohlebefeuerten Kraftwerken, VGB Kraftwerkstechnik 77 (2) (1997) 129–134.