High-temperature oxidation of FeCrAl alloys: the effect ...

MK_mk17722 – 2.4.03/druckhaus köthen

A. Dimyati et al.: High-temperature oxidation of FeCrAl alloys

A. Dimyatia,c, H. J. Penkallab, P. Untoroc, D. Naumenkob, W. J. Quadakkersb, J. Mayera a

Gemeinschaftslabor für Elektronenmikroskopie, RWTH, Aachen, Germany Forschungszentrum Jülich, Jülich, Germany c National Nuclear Energy Agency, Kawasan PUSPIPTEK, Serpong, Indonesia

© 2003 Carl Hanser Verlag, Munich, Germany

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Not for use in internet or intranet sites. Not for electronic distribution.

b

High-temperature oxidation of FeCrAl alloys: the effect of Mg incorporation into the alumina scale Dedicated to Professor Dr. Dr. h. c. Manfred Rühle on the occasion of his 65th birthday

Apart from reactive element additions, commercial FeCrAl alloys contain a number of poorly controlled or unavoidable impurities, which might significantly affect the material's oxidation behaviour. In the present study it was found that magnesium, if present in a concentration of only 80 ppm, can substantially affect composition and morphology of the alumina surface scales. This was found for scales formed after short (a few minutes) as well as very long (up to 1000 h) oxidation treatments. Magnesium appeared to diffuse outward through the alumina scale and it became enriched in the outer part of the oxide scale, forming a layer of MgAl2O4 spinel. An important effect associated with the Mg enrichment and spinel formation was the development of porosity within the scale which likely affects the scale growth rate and adherence, however, did not fundamentally change the sub-parabolic character of the scale thickening rate, which has previously frequently been reported for pure -alumina scales. Keywords: High-temperature oxidation; FeCrAl alloys; Transmission electron microscopy; Focused ion beam technique

1. Introduction FeCrAl alloys are used as functional or constructional materials in a large number of high-temperature applications. In most technically relevant atmospheres and at temperatures as high as 1400 °C, these materials with typical composition Fe-20Cr-5Al (mass%) form very protective alumina scales that prevent rapid corrosive degradation of the metallic components. However, in spite of the excellent stability of the alumina surface scales, the lifetime of the FeCrAl-based alloy components can be limited by oxidation, because the scale-forming element Al is present only in a limited amount within the alloy matrix. After long exposure times, growth and eventually spallation of the protective alumina scales can lead to a substantial depletion of this Al reservoir in the component. Below a critical Al level, the protective alumina scale can no longer be maintained. As the oxidation products of the major alloying elements, i. e., Fe and Cr, are not protective at high temperatures, a rapid “breakaway” corrosion process occurs, eventually resulting in a failure of the component within only a few hours [1]. 180

In attempts to extend the oxidation-limited lifetime of the currently available FeCrAl-based components, it has been found that both the growth rate and the scale adherence can substantially be improved by minor additions of reactive elements (REs) [2 – 4]. Typical REs that are added in the range of 0.05 mass% to the main composition of commercial wrought alloys are Y, La, Ce, Zr, and Hf. The effects of REs on the high-temperature oxidation resistance of alumina-forming alloys include a decrease in the oxidation rate, a change in the scale structure and morphology as well as an improved scale adherence. The main physical effects responsible for the above observations have been claimed to be the suppression of the Al diffusion in the alumina scale and the prevention of the deleterious sulphur segregation to the scale/metal interface [5 – 8]. In the fundamental studies of the RE effect, mostly pure metals or model alloys have been used, and the effect of many common impurities, which are unavoidably present in commercial materials, has hardly been investigated. One of the impurities which is frequently observed in commercial alloys is Mg. Although its concentration is commonly quite low, i. e., 20 – 100 ppm, recent studies [9, 10] indicated that it can substantially alter the morphology and thus the growth rates of alumina scales on FeCrAl alloys. In the present study, the oxidation behaviour of a commercial FeCrAl alloy has been investigated, with the main emphasis on the mechanisms of Mg incorporation in the alumina scales. By studying the time dependence of the evolution of the oxide film on a time scale which ranges from a few minutes to 1000 h, detailed insights into the scale growth mechanisms were obtained. Transmission electron microscopy (TEM) was used as a main tool for the microstructural characterisation. Owing to the layered scale morphology found already after very short exposure times, specimen preparation was very critical and the focused ion beam (FIB) technique proved to be indispensable for obtaining thin sections with uniform thickness through the complete oxide scales. The results of the present studies give fundamental insight into the scale growth mechanisms which extend beyond the effect of the Mg incorporation and will be discussed in detail.

2. Experimental For the experimental investigations, a wrought commercial FeCrAl alloy, produced by conventional casting followed  Carl Hanser Verlag, München

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Table 1. Chemical composition of the studied FeCrAl alloy in mass %a or mass ppmb. Fea

Cra

Ala

Mgb

Nib

Mnb

Mob

Sib

Yb

Tib

Zrb

Hfb

Sb

Nb

Cb

Base

20.10

5.68

80

1550

2040

60

900

600

35

330

340

20

21

230

by hot and cold rolling, was used. The composition of the studied FeCrAl alloy is shown in Table 1. A number of alloying additions have been added to improve the high-temperature mechanical properties and also to increase the oxidation resistance. In addition, Mg was found as the main impurity with a concentration of 80 ppm. From the starting material, specimens 20 mm × 10 mm × 1.5 mm in size were cut, ground using SiC paper, mechanically polished with 1 lm diamond paste and finally degreased ultrasonically in a detergent prior to oxidation. The isothermal oxidation exposures were performed in synthetic air at 1200 °C in a resistance-heated furnace. The exposure times were 2 and 10 min, 2, 20, 100, 600 and 1000 h. Oxidation times were counted, starting from the time at which the specimen was moved into the hot furnace. For a detailed analysis of the scale growth kinetics, isothermal oxidation tests up to 100 h were carried out in a SETARAM microbalance. After oxidation, the specimens were examined by optical metallography, scanning electron microscopy (SEM) and TEM, coupled with energy-dispersive X-ray analysis (EDX) and electron energy loss spectroscopy (EELS). TEM cross-sections were prepared by conventional techniques and on an FIB workstation. For the conventional preparation, thin slices of the material including the alumina scale were embedded in special alumina cross-section holders, sliced, dimpled and finally ionthinned to electron transparency [11]. For the FIB specimens, first a tungsten cap layer was deposited on the surface in the form of a thin stripe. The stripe is located along the top edge of the thin lamella which is then prepared by cutting trenches perpendicular to the surface on both sides of this lamella. In the final step, the lamellae are extracted using an ex situ lift-out process and placed on a thin carbon support film. The TEM characterisation was performed on a Jeol 2000 FX and a Philips CM 20. Phase analysis in the TEM crosssections was carried out by quantitative electron diffraction, EDX and EELS analysis.

growing oxide with important consequences for the scale morphology. The dominating precipitates are chromium carbides (Cr23C6) which decorate part of the grain boundaries in the as-received material. Fig. 1c shows the typical arrangement of the chromium carbides as extracted on a carbon replica. Besides equiaxed particles which decorate the grain boundary itself, a typical butterfly-shaped arrangement of thin lamellae growing from the grain boundary into the surrounding matrix can be observed. In addition to these precipitates, HfC and (Hf/Zr)C particles can also frequently be observed. During the oxidation heat treat-

3. Experimental results 3.1. Effect of heat treatment on alloy microstructure In the as-received condition, the alloy microstructure consists of severely deformed grains of approximately 100 lm in length (diameter) and 10 lm in thickness (Fig. 1a). During exposure at high temperature, the material tends to recrystallize and after short oxidation time has a stable microstructure consisting of equiaxed grains of approx. 500 lm size, as is shown in Fig. 1b for the material isothermally oxidised for 100 h at 1200 °C. Detailed studies were also performed on the carbide precipitates present in the alloy and their evolution during the heat treatment. As will be shown in the subsequent sections, they are not only active as possible grain growth inhibitors in the alloy matrix, but can also be incorporated in the Z. Metallkd. 94 (2003) 3

Fig. 1. (a) Optical micrograph of a cross-section of the alloy in the asreceived condition. The microstructure shows grains with a high aspect ratio elongated in the rolling direction. (b) Microstructure of the alloy matrix after 100 h of isothermal oxidation in air at 1200 °C. (c) TEM micrograph of an extraction replica of the alloy in the as-received condition. The grain boundaries are decorated with chromium carbides, and thin lamellae extend into the bulk of the alloy.

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ment, a grain coarsening and clustering of the precipitates can be observed and the number of chromium carbide particles decreases as the existing low amount of carbon in the alloy becomes tied up in the thermodynamically more stable carbides of Hf and/or Zr.

In order to get a detailed information on the initial stages of oxidation, the respective heat treatments in air were stopped after a sequence of time intervals by removing the specimen rapidly from the hot zone in the furnace. The time intervals discussed in this section are 2 and 10 min, and 2 and 20 h, respectively. A detailed characterisation of the resulting scale thickness, morphology and chemical composition was carried out on FIB prepared TEM cross-sections. After a total time of 2 min at a temperature of 1200 °C (including the time for heating up the specimen by moving it rapidly into the hot furnace), a dense Al2O3 scale with a thickness between 200 and 300 nm has formed. A TEM micrograph of an FIB cross-section of the sample is shown in Fig. 2a. In the oxide scale, grains with two distinctly different sizes appear. By analytical investigations, it could be shown that both types of grains consist of a phase with a nominal composition of Al2O3. By careful analysis of diffraction patterns, it could also be verified, that both types of grains are formed by the alpha phase, i. e., a bimodal size distribution of alpha alumina grains exists after the first 2 min of heat treatment. Occasionally, as is shown in Fig. 2b, small Cr2O3 particles can be observed which are located mainly near the al-


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3.2. Initial stages of oxidation

Fig. 2. (a) TEM micrograph of the FIB section through the scale formed during a total time of 2 min at 1200 °C (including the time to bring the specimen to temperature). (b) In the oxide scale, Cr2O3 particles can be observed at the boundaries between individual grains in specific areas.

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Fig. 3. A complex layer system has formed in the oxide scale after 10 min of exposure in air at 1200 °C. Besides the small and large α-Al2O3 grains, an outer layer with characteristic pores and frequently observed MgAl2O4 spinel grains has grown.

loy matrix at oxide grain boundaries and at the boundary between the small and the large alumina grains. From their characteristic size distribution and their relative arrangement, it can be suspected that these Cr2O3 particles are formed from chromium carbide (Cr23C6) predecessors at grain boundaries near the original matrix surface (Fig. 1c). It could, however, not be clarified whether the Cr2O3 particles are formed by direct oxidation of the chromium carbides incorporated in the scale, or whether dissolution of chromium in the oxide formed during heating with subsequent precipitation upon cooling plays the dominant role. After 10 min of exposure, the oxide scale has reached a thickness of about 500 nm and, compared to the result of the shorter oxidation time, distinct changes in the scale morphology can be observed (Fig. 3). In the inner part of the oxide scale, the same characteristic features can still be observed, i. e., a layer of small Al2O3 grains with some Cr2O3 grain boundary precipitates terminates the oxide scale towards the alloy matrix, and some larger Al2O3 grains form a centre layer of the scale. However, on top of these grains, a third layer has formed towards the outer surface. This third layer is characterised by the existence of MgAl2O4 spinel grains at the top near the surface and exhibits small pores with typical dimensions of 50 – 100 nm underneath. The spinel grains have dimensions of 100 – 200 nm and do not form a dense layer, but can be observed frequently in the topmost layer. The chemical nature of the phase distribution in the oxide scale was investigated by two-dimensional mapping with EDX. Figure 4 shows the resulting maps for three different chemical elements, Al, O and Mg, in a characteristic specimen area. Whereas the oxygen signal, with the exception of occasional pores, is distributed evenly throughout the scale, a distinct Al depletion is visible in the outer layer. The areas in which a lower Al signal is observed match nicely with the occurrence of Mg, indicating clearly that spinel grains were formed. A quantitative analysis of the signal resulted in the correct atomic ratios of Al and Mg expected for spinel. Z. Metallkd. 94 (2003) 3



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Fig. 4. Bright-field image (top left) and EDX distribution maps for O, Al, and Mg of a characteristic specimen area. A reduction in the Al signal is observed in the areas in which Mg occurs, indicating that stoichiometric spinel has formed.

In a next step after the very short exposure times, specimens were oxidised for 2 and 20 h, respectively, and subsequently analysed. Besides the increase in thickness, no major differences in the chemistry and morphology of the scale were found over the whole time interval from 2 to 20 h. Thus, the results will be summarised using the 20 h specimen as example. Figure 5a shows a TEM micrograph of an FIB cross-section for the scale obtained after 20 h of oxidation. The scale appears to consist of columnar grains near the scale/alloy interface and equiaxed grains near the oxide surface. The columnar grains form a dense layer on top of the substrate, which extends to at least half of the total scale

Fig. 5. (a) TEM micrograph of the FIB cross-section of the oxide scale after 20 h, and elemental distribution images for (b) Al and (c) Mg obtained by mapping with EDX for a characteristic specimen area.

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thickness in all of the specimens investigated in this time interval. As the scale thickness increases with time, the columnar grains not only grow in the growth direction of the scale, but a grain coarsening also occurs in the lateral direction, in such a way that the aspect ratio of the columnar grains remains nearly constant. The same also holds for the microstructure in the upper part of the oxide scale, in which both the equiaxed grains and the existing pores show a coarsening during scale growth in this time interval. In terms of the pores, no clear evidence for either a decrease or an increase of their volume fraction with increasing oxidation time could be found. Analytical TEM investigations revealed that an almost dense layer of spinel grains (MgAl2O4) covers the top of the oxide layer after 20 h of oxidation. The results of EDX analysis correspond to an MgAl2O4 phase, with minor amounts of Mn, which substitutes for the Mg. The distribution of the spinel grains can nicely be seen in elemental distribution images obtained by mapping with EDX, as is demonstrated for the material heat-treated for 20 h in Figs. 5b and c. All the spinel grains, which have been observed, reside at or near the surface of the oxide scale. 3.3. Microstructure after oxidation for 100 to 1000 h After 100 h of oxidation time, only gradual changes in the morphology of the oxide scale can be observed, and the resulting microstructure is very similar to the one obtained after 20 h of oxidation. Fig. 6 shows an entire FIB lamellae extracted from the scale obtained after 100 h of oxidation. The dense columnar central part of the oxide layer, a porous layer, and the top layer with equiaxed grains can clearly be identified. EDX analysis revealed that an almost continuous spinel layer persists at the surface. A fracture surface of the specimen heat-treated for 1000 h is shown in Fig. 7a. It is evident that the dense layer with columnar grains has increased in thickness, but also that near the outer surface of the scale a porous layer with equiaxed grains still exists. The SEM micrograph depicted in Fig. 7a was acquired with back-scattered electrons (BSE) and shows the phases containing large amounts of heavy elements in bright contrast. Elongated precipitates and clusters of precipitates appearing in bright contrast can be seen along grain boundaries in the inner part of the scale. TEM analysis revealed that the precipitates in the inner part of the scale consist of Hf/Zr carbide (Fig. 7b). This is in agreement with previous studies [12] which showed that the alloy carbide precipitates, if becoming embedded in

Fig. 6. TEM micrograph of the entire FIB lamellae prepared as crosssection through the scale on the material oxidised for 100 h at 1200 °C.

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Fig. 7. (a) Fracture surface of the oxide scale after 1000 h of oxidation, SEM micrograph showing the BSE signal. (b) TEM micrograph of Hf/Zr carbide and Y-rich precipitates in the oxide scale near the alloy interface.

the inwardly growing alumina scale, initially remain stable as long as they prevail in the inner part of the scale where the oxygen partial pressure is extremely small. Upon further scale thickening by inward growth, the carbides become completely embedded in the scale and are thus subjected to a higher oxygen partial pressure. This leads to transformation of the carbides into oxides [12]. The size of such Zr/Hf-rich oxide phases obviously depends on the size of the original carbide precipitates in the alloy matrix. Depending on the alloy manufacturing method, the Hf/Zr-rich precipitates in the alloy sometimes exist in form of relatively large clusters, frequently in combination with other reactive elements such as Y. If they become embedded in the inwardly growing alumina scale, they can occasionally be found as globular precipitates in the oxide scale. An SEM image of a corresponding area is shown in Fig. 8, along with the EDX elemental maps for the three reactive elements. In addition, the EDX maps for Mg and Al reveal that a more or less continuous layer of spinel is still present on top of the oxide scale. 3.4. Time dependence of scale thickness The FIB-cut TEM specimens are ideally suited for a thickness measurement in the TEM. Since the original surface has been covered with a W cap layer, a surface marker exists which helps to obtain a quantitative measure of the

oxide scale thickness. The thickness measurements were performed for all the oxidation times extending from 2 min to 100 h in the furnace. Only in the last specimen (100 h, Fig. 6), the W cap is not visible all along the surface, i. e., the thickness measurement in this case tends to underestimate the actual thickness. For the specimens heat treated for 600 and 1000 h, the scale thickness was measured from SEM images. Since the measurements cover a large range of oxidation times, i. e., four orders of magnitude on a time scale, it is very difficult to plot the results on a linear time scale. However, for a constant mass density of the oxide scales, the measured thickness increase should exhibit the same time dependence as the weight gain. Therefore, the measured thickness values were plotted as a function of t1/3, which reflects the near-to-cubic time dependence of the experimental isothermal weight change data during isothermal oxidation up to a total of 100 h, which was previously also found for FeCrAl-based oxide dispersion-strengthened alloys [13]. The results are shown in Fig. 9, and a linear relationship between scale thickness and t1/3 should exist. It can be seen that in particular the first four data points can very nicely be fitted with a straight line. Only the last data points show some scatter, which may be attributed to fact that the topmost part of the scale after 100 h of oxidation may have been sputtered away during FIB preparation. For longer oxidation times (between 600 and 1000 h), a substantial deviation from the near-cubic time dependence occurs. Possible reasons for this effect will be discussed in the following section.

4. Discussion In the present investigations, a detailed TEM characterisation of the oxide scale development on FeCrAl alloys has been performed. It has been shown that Mg, although being a minor impurity, can have a major influence on the oxide scale morphology and the phases present in the scale. The detailed analysis of a time sequence of specimens, for which the oxidation process was interrupted after total times ranging from 2 min to 1000 h, has given insight into the growth and microstructural features of the oxide scale as function of time. After short oxidation times, it was found that a pronounced bimodal size distribution prevails, with large grains of alpha alumina residing on top of smaller grains be-

Fig. 8. SEM micrograph (BSE signal) and EDX maps of a cross-section after 1000 h of oxidation at 1200 °C. The scale surface is on the left side of the image, the matrix alloy on the right side. A dense spinel layer can be observed on the surface of the scale, and clusters of Hf-, Zr- and Y-rich precipitates are visible in the scale.

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Fig. 9. Time dependence of the increase of the scale thickness measured from the FIB sections on TEM micrographs and from SEM images for the specimens oxidised for 600 and 1000 h. The thickness data are plotted as a function of a t1/3 time scale. A linear relationship is observed for the first 600 h of oxidation, indicating that the cubic time dependence is nicely obeyed. Only for the sample oxidised for 1000 h, a larger thickness than expected is obtained. This may be attributed to the formation of microcracks which lead to an enhanced oxidation rate.

longing to the same phase. The two different populations of the phase can only be understood if two different nucleation mechanisms have been active during these initial oxidation steps. A possible explanation is that, during heating up the specimens in air, metastable phases of alumina [14] grow as large grains on the surface of the specimen and, before reaching the final temperature, already undergo a phase transformation to form large grains of the α phase. Further oxidation at a temperature of 1200 °C then proceeds via nucleation and growth of much smaller α phase grains. The nucleation of these smaller grains at the scale/alloy interface and the presence of Cr2O3 particles at the interfaces between the small and the large oxide grains suggest that the main mechanism for the formation of the small α phase grains is based on the inward diffusion of oxygen. The reason for the presence of the small chromia particles in the scale could not unequivocally be identified, however their formation must be correlated with a reaction occurring at the bare metal surface in the very early stages of oxidation. A possible source for the chromia might be the Cr23C6 grain boundary precipitates prevailing at the alloy surface. Upon start of the oxidation, they will be oxidised and subsequently become embedded in the growing alumina scale. A second possible mechanism was proposed in Ref. [15]. SNMS analyses revealed the transient alumina-based scale on FeCrAl, based oxide dispersion-strengthened (ODS) alloys to contain substantial amounts of chromium oxide. Depending on the actually prevailing scale growth mechanism, the transient chromia was retained at the oxide surface or became embedded in the alumina-based scale upon further oxidation treatment. After longer oxidation times, the chromia particles dissolved in the alumina, an effect which was also observed in the present study. From the present results, it is not possible to unequivocally derive, which of the mentioned mechanisms is responsible for the formation of the chromia particles in the transient surface oxide scale. The first indications for spinel formation, induced by incorporation of Mg and minor amounts of Mn, were obZ. Metallkd. 94 (2003) 3

served after 10 min of oxidation, both elements diffusing from the bulk alloy through the inner α-Al2O3 layer to the scale surface, probably via scale grain boundaries, as described for Y in Ref. [13]. According to the results of various analyses, evidence was found that the Mg enrichment in the outer scale occurs by diffusion of Mg from the bulk alloy into the oxide scale. It is interesting to note that the spinel layer persists up to the highest oxidation times and seems to obey the same growth kinetics as the total scale thickness. Even though being a minor impurity, Mg thus seems to play an important role in the overall diffusion processes and may contribute to the excellent oxidation behaviour of the present alloy by slowing down the inward diffusion of oxygen. The formation of the large voids in the outer scale also seems to be related to the spinel formation. The spinel nucleation is likely to result in a volume change, induced by the change in lattice structure (hexagonal to cubic), by the differences in atomic volume of alumina and spinel, and/or by the stresses induced by the phase transformation. This is corroborated by the experimental finding that the first voids in the scale form in a systematic way along with the first spinel grains found after 10 min of oxidation time. A scientifically as well as technologically important observation of the present studies is the sub-parabolic, nearcubic time dependence of the alumina scale thickening rate (Fig. 9). If scale growth would proceed, as proposed in classical Wagner’s oxidation theory, by lattice diffusion of aluminium and/or oxygen ions, the scaling kinetics would obey a parabolic time dependence, i. e.: x2 ¼ kt

ð1Þ

in which x is the scale thickness, t the time and k the oxidation rate constant. Eq. (1) is conveniently described in terms of an area-specific mass change (Dm) ðDmÞ2 ¼ Kp t

ð2Þ

where Kp is the parabolic rate constant. In practice, Kp is commonly determined by determining the slope of a plot of (Dm)2 versus time, in other words d ðDmÞ2 ¼ Kp dt

ð3Þ

Alumina scales on reactive element doped FeCrAl alloys have frequently been found to grow to a large extend by grain boundary diffusion, especially of oxygen [7, 16]. It is easy to show that also for this growth process the scaling kinetics obey a parabolic time dependence as long as the grain boundary density is independent of scale thickness and oxidation time. The k value in Eq. (1) can then be described by an expression of the type [17, 18] k/

D0 d Dl r RT

ð4Þ

in which D0 is the oxygen grain boundary diffusion coefficient, d the grain boundary width, r the grain size, Dl the oxygen chemical potential gradient across the scale, R the gas constant and T the oxidation temperature. It is obvious from Eq. (4) that parabolic oxidation does not exist if the oxide grain size is time-dependent. Several authors (see, e. g., Ref. [19]) have used a similar approach to explain 185



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non-parabolic oxidation of various types of oxides. In most cases, it was observed or assumed that during the high-temperature exposure growth of the oxide grains occurred. For the case of alumina scales on FeCrAl-based oxide dispersion-strengthened alloys, it was found [13, 15] that at not too high temperatures (e. g., 1100 °C) the non-constant value of r was not primarily related to oxide grain growth but to the fact that the grains being formed at time t2 are larger than those being formed at time t1 (with t2 > t1). As the scales appeared to grow nearly exclusively by oxygen grain boundary diffusion, the grain size increased in the main scale growth direction, i. e., from the scale/gas towards the scale/alloy interface. The oxidation kinetics could in many cases quite accurately be described by a power law-type time dependence [13] Dm ¼ ks tn

ð5Þ

whereby the value of n was found to be close to 0.33, i. e., a near-cubic time dependence prevailed. This observation might be explained by assuming a grain growth as function of time, e. g., of the type r ∼ r0 + tm, as was e. g., done in Refs. [13, 19]. However, this is probably not a completely adequate description in the case of alumina scales. The main reason for the sub-parabolic oxidation is the increase of the grain size from the scale/gas towards the scale/ alloy interface, which was also found in the present investigations. Additionally, indications were found that at the high oxidation temperature of 1200 °C the grain diameter perpendicular to the growth direction also increases with time. From Eq. (4), it is then obvious that the scale growth kinetics should obey a sub-parabolic time dependence, as found experimentally by the scale thickness measurements (Fig. 9). This is in full agreement with the weight change data measured by thermogravimetry during isothermal oxidation up to 100 h (Fig. 10a). The thermogravimetric data until 100 h oxidation (Fig. 10a) can quite accurately be described by a power law time dependence [Eq. (5)] whereby the growth exponent n is quite close to 1/3, i. e., a near-cubic time dependence prevails, as has frequently been observed previously for alumina formers [13, 15]. Although a sub-parabolic oxidation can easily be explained as discussed above, it is presently not yet clarified why the exponent n is frequently very close to 1/3. Our TEM studies revealed that in the present case the alumina scale growth does not exclusively proceed via inward grain boundary oxygen diffusion. The time dependence of the detailed scale composition obviously leads to the conclusion that outward transport of Mg and, to a lesser extend, Mn contributes to the scale growth process. Considering the limited solubility of both elements in -alumina, it seems to be justified to assume that the outward transport of these minor elements also proceeds via grain boundaries rather than via the alumina lattice. For the reasons described above, the time dependence of Mg and Mn incorporation in the scale should then also obey a sub-parabolic time dependence and would therefore not substantially alter the ideal, sub-parabolic time dependence expected for a pure -alumina scale, as long as no fundamental changes in growth mechanisms occur. If an ideal, sub-parabolic time dependence [Eq. (5)] would prevail over the total time period of the thermogravimetry analysis (i. e., 100 h), a conventional plot of (Dm)2 as function of time does not reveal a line with a con186

stant slope which equals Kp, because ðDmÞ2 ¼ ks2 t2n

ð6Þ

and consequently d ðDmÞ2 ¼ 2nks2 t2n1 dt or in logarithmic form

ð7Þ

d ðDmÞ2 ¼ C þ ð2n 1Þ log t ð8Þ dt A double-logarithmic plot of dðDmÞ2 /dt as function of time should reveal a straight line with a slope of 2n – 1. If n < 0.5, i. e., for sub-parabolic oxidation kinetics, the slope should be negative and for the case of n = 1/3, the slope of the curve should be – 1/3. Figure 10b indeed shows a negative slope, however, the curve cannot be represented by a single straight line. Initially, it exhibits a very steep slope. This might relate to the early stages of oxidation in which log

Fig. 10. (a) Mass change during isothermal oxidation at 1200 °C in synthetic air. The mass change data have been fitted using a power law time dependence. The derived growth rate exponent n is near 0.33, i. e., a sub-parabolic oxidation kinetics prevails. (b) Double-logarithmic plot of dðDmÞ2 /dt as function of time derived from the mass change data in (a). (c) Mass change during discontinuous oxidation at 1200 °C in air. The mass change data have been fitted using a power law time dependence, and a somewhat higher growth rate exponent compared to (a) is obtained.

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the external as well as the internal equi-axed grains are being developed. The second branch of the curve corresponds to the growth of the columnar grains, accompanied by a grain boundary density decreasing with time. In this time period the slope appears indeed to be close to – 1/3, i. e., the value which would prevail in case of ideal cubic kinetics. It is difficult to explain, why, after approximately 20 h, the slope of the curve again tends to decrease. A possible explanation might be that the Mg reservoir in the alloy becomes exhausted due to its extremely small concentration. After that, the scale growth would continue by α-alumina formation only, and the additional scale thickening, i. e., weight change, imparted by Mg would vanish. A simple mass balance consideration, assuming that the 80 ppm Mg in the bulk alloy would completely be tied-up in the scale as MgAl2O4, reveals that the maximum thickness of a completely dense spinel layer would be approximately 2 lm. Considering the fact that the spinel layer is quite porous and that it contains minor amounts of Mn, the spinel thickness analyses from the TEM studies illustrate, that it is certainly not unlikely to assume, that after around 20 h of oxidation, practically the total amount of Mg initially present in the alloy is exhausted, which will undoubtedly affect the scaling kinetics. Beyond oxidation times of hundred hours (Fig. 10c), the growth exponent n [Eq. (5)], tends to become larger than the value of approximately 0.33. This effect might be related to formation of microcracks in the scale, as the longterm data in Fig. 10c do not relate to real isothermal exposures, but to a test with intermediate cooling to room temperature every 100 h. Another reason for the increasing n value is that after longer oxidation times additional factors start to affect the scaling kinetics. Commercial FeCrAl alloys, such as the one used in the present study, commonly contain precipitates of carbides, nitrides and/or oxides originating from the added reactive elements (Y, Hf, Zr). These precipitates sometimes prevail as inhomogeneously distributed clusters in the alloy matrix. Upon prolonged oxidation, more and more of these clusters become embedded in the inwardly growing oxide scale [12, 20], which leads to locally enhanced scale growth (Fig. 8), thereby disturbing the idealised growth process described above for oxidation times of less than approximately 100 h.

5. Conclusions • From the very early stages to about 600 h of oxidation, the scale growth follows a near-t1/3 time dependence which can be explained by assuming the scale growth to predominantly occur along oxide grain boundaries, whereby the grain boundary density increases in the dominating scale growth direction. • Magnesium is easily incorporated in alumina scales on FeCrAl alloys, although it is present in the alloy only in very low concentrations (≈ 80 ppm). Magnesium, together with Mn, diffuses through the inner -Al2O3 layer in direction of the scale/gas interface. Reaction of the Mg with alumina leads to spinel formation accompanied by initiation of voids in the outer part of the scale. • The amount of Mg impurity should, therefore, be carefully controlled in FeCrAl alloy manufacturing because it can substantially affect the properties of the protective alumina layer. Z. Metallkd. 94 (2003) 3

• Cr-, Hf-, and Zr-carbides are not only important for the properties of the alloy matrix, but when incorporated into the growing scale may influence the scale growth properties. For Hf and Zr, the beneficial reactive element effect dominates for shorter oxidation times, whereas embedding of HfC and ZrC particulates in the scale after longer oxidation times leads to enhanced diffusion and accelerated scale growth. Future research will include a more detailed investigation at shorter time intervals of the initial stages of oxidation, and an analysis of alloys for which the surface activity of Mg has been increased by Mg implantation in an ion accelerator. The authors thank Dr. A. Schertel (FEI Germany) for his expert help with FIB specimen preparation, Dr. Syahril and Dr. T. E. Weirich for help with diffraction pattern analysis, and gratefully acknowledge support by the German ministry of education and research (BMBF) and the Indonesian national nuclear energy agency (BATAN) in the framework of the bilateral co-operation project IDN 99/004.

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(Received November 21, 2002) Correspondence address Prof. Dr. J. Mayer Gemeinschaftslabor für Elektronenmikroskopie Ahornstr. 55, D-52074 Aachen, Germany Tel.: +49 241 802 4345 Fax: +49 241 802 2313 E-mail: [email protected]

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