Diffusion and growth mechanism of Al2O3 scales on ferritic Fe-Cr-Al alloys

Materials Science and Engineering A247 (1998) 248 – 262

Diffusion and growth mechanism of Al2O3 scales on ferritic Fe-Cr-Al alloys K. Messaoudi, A.M. Huntz *, B. Lesage Laboratoire de Me´tallurgie Structurale d’Orsay, UA CNRS 1107, Uni6ersite´ Paris XI, 91405, Orsay, France Received 22 April 1997; received in revised form 12 September 1997

Abstract In order to determine the growth mechanism of a alumina scales, cationic and anionic diffusion studies have been performed on scales developed on three alumina former Fe-Cr-Al alloys. Two of them were classically elaborated and differed by the presence or not of implanted yttrium. The third one is an ODS alloy, the MA 956. First, the oxidation behaviour in oxygen was characterised in a temperature range 1000–1200°C, with morphological studies and analyses of the scales. Then, anionic self diffusion coefficients were determined by isotopic exchange 16O– 18O, cationic self diffusion was simulated by Cr diffusion and the penetration profiles were determined by SIMS. These profiles were analysed considering two domains: the first one relative to effective diffusion and the second to grain boundary diffusion. The f values, fractions of sites associated with the grain boundaries were modified by considering the roughness of the scales. The anionic and cationic self-diffusion coefficients are compared with those extrapolated from massive aluminas and the experimental parabolic oxidation constants are compared with those calculated with diffusionnal data. The results obtained are discussed by taking into account the differences in the scale morphology which looks like an important parameter. Then, the effect of active elements on the morphology and on the transport and mechanical properties is commented. © 1998 Elsevier Science S.A. All rights reserved. Keywords: Fe-Cr-Al alloys; Al2O3 scales; Massive Al2O3; Self-diffusion; Morphology effect

1. Introduction At high temperatures in aggressive environment, alloys are hoped to form a protective scale. The scale is protective if it is continuous, adherent and characterised by a slow growth rate. Al2O3 scales are often used but in order to improve the protective qualities of such scales, it is necessary to understand the growth process. Many attempts were investigated, using different techniques, to determine the growth mechanism of the Al2O3 scale on Fe-Cr-Al alloys but they lead to scattered results [1–3]. Some authors suggest that Al2O3 scale grows predominantly by outward Al diffusion [4 –6], others have reported that scale growth occurs by inward oxygen diffusion [7 – 9] and some results indicate that the growth mechanism occurred by both cation and anion transport in the Al2O3 scales on Fe-Cr-Al [10]. The most recent works are those of Clemens et al. * Corresponding author. Tel.: + 33 1 69416318; fax: + 33 1 69417833. 0921-5093/98/$19.00 © 1998 Elsevier Science S.A. All rights reserved. PII S0921-5093(97)00711-9

[11] concerning oxygen diffusion in an alumina scale developed on ODS MA 956 alloy and they assumed that the alumina scale grows exclusively by oxygen diffusion along grain boundaries. Diffusion results were also obtained in alumina developed on b-NiAl by cationic and anionic self diffusion tests and also by an electrochemical method [12]. Prescott et al., for alumina scales formed on b-NiAl and FeAl showed that the scales grew through a combination of Al and oxygen diffusion, but they did not determine diffusion coefficients from their penetration profiles [13]. Another possibility consists in considering the diffusion coefficients determined in massive aluminas undoped and doped with yttrium [14–17]. When extrapolated to the range of oxidation temperatures, from recent results of Prot and Le Gall [14–17], the diffusion coefficients are lower than those obtained in alumina scales. For instance, the order of magnitude of the bulk self diffusion coefficients in alumina scale is about 10 − 17 cm2 s − 1 at 1100°C [11,12], while the value extrapolated from Prot and Le Gall works is rather about 10 − 22 cm2 s − 1.

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Moreover, the oxidation constants calculated from diffusion coefficients determined on massive aluminas are smaller than the experimental ones. In order to clarify these differences, diffusion experiments were performed on alumina scales grown on two Fe-Cr-Al alloys of different origins. As will be seen further on, they both develop an alumina scale but of various morphologies. The different alloys were oxidised in 1 atm oxygen and the parabolic oxidation constants were experimentally determined so that it will be possible to compare them with the calculated ones from the diffusion data. The oxygen diffusion coefficients in the alumina scales are determined by isotopic exchange 16O – 18O. Due to the difficulties in studying aluminium diffusion and since the chromium diffusion coefficient in alumina is of the same order of magnitude as that for aluminium diffusion [18], the cationic self-diffusion is simulated by chromium diffusion. Moreover, chromium ions are isovalent with aluminium ions and chromia and alumina presents a complete mutual solubility. Both anionic and cationic penetration profiles were established by SIMS (CNRS Bellevue) and analysed considering two domains as suggested by Tsaı¨ et al. [19,20] and shown in (Fig. 1). The first one is related to effective diffusion Deff and the second one to grain boundary diffusion Dgb. Deff is given by Hart [21]: Deff = (1− f )Db + fDgb

(1)

where Db is the bulk diffusion coefficient, Dgb is the diffusion coefficient along grain boundaries and f the fraction of sites associated with grain boundaries. Generally f is given by f =3d/f, where f is the average grain size and d is the grain boundary width. This is correct for smooth surfaces as in the case of massive aluminas, but must be modified if rough surfaces are considered as it is the case for scales. According to [19,20], fr, the fraction of sites associated with grain boundaries for rough surfaces is given by:

fr =

3d (f+ 4Ra)

249

(2)

where Ra is the surface roughness. In the case of oxygen diffusion, the first part of the penetration profiles allows to determine Deff according to Fick’s solution given for constant superficial concentration [22]:





x (C(x) − Cs) =erf (3) (C0 − Cs) 2 Defft where C0 is the natural tracer concentration in the sample, equal to 0.2% and Cs is the constant superficial tracer concentration. From the second part of the profiles corresponding to grain boundary diffusion, Dgb can be determined by the equation proposed by Whipple–Le Claire [22]: Dgbd= 0.661(− Pgb) − 5/3(4Db/t)1/2

(4)

where Pgb is the slope of the curve ln C =f(x 6/5) plotted in the second part of the experimental profiles, and t is the diffusion time. Taking d= 10 − 7 cm and combining Eqs. (4) and (1), a second-order equation with an unknown Db is obtained (1−fr)Db + (1.332× 107)fr(− Pgb) − 5/3(t) − 1/2 −Deff =0 (5) Db can be calculated with Eq. (5), and then Dgb from Eq. (4). In the case of cationic diffusion, 54Cr was deposited as a thin layer on the oxide scale by vacuum evaporation and Deff is obtained from the Fick’s solution for thin films: C(x) = M(pDefft) − 1/2 exp(− x 2/4Defft)

(6)

which leads to Deff =

1 4pt

(7)

where p is the slope of the curve ln(54Cr)= f(x 2) plotted in the first part of the diffusion profile. The bulk diffusion coefficient and the grain boundary diffusion coefficient are then determined as previously.

2. Materials and experiments

Fig. 1. Diffusion profile analysed considering two domains.

The chemical composition of the Fe-Cr-Al alloys is given in Table 1. It can be observed that Imphy alloy is doped with zirconium and that MA956 is doped with Y2O3. Imphy alloy was produced by classical metallurgy while MA956, supplied by INCO Alloy International, Hereford UK, was produced by mechanical alloying and is an oxide dispersion strengthened (ODS) alloy. Thus, it has a high creep resistance. Specimen of area : 1 cm2 were cut with a diamond saw and polished with SiC paper up to 1200 grade. Then, they were ultrasonically cleaned in volasil and washed with iso-

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Table 1 Chemical composition of the alloys Elements (wt %)

(ODS alloy) (S.A. Imphy)

Fe

Cr

Al

Si

Mn

Ni

Zr

S (ppm)

C

Ca

bal bal

19.7 22.6

4.62 4.36

0.09 0.85

0.12 0.38

0.07 0.35

0.13

20 60

0.01 0.05

0.01

propanol. Incorporation of yttrium in the Imphy alloy has been performed by implantation (CSNSM, Orsay) at 150 keV with a dose of 1016 atm cm − 2. Thus, in such a case, this alloy is Y and Zr doped. The oxidation kinetic tests were carried out in a SETARAM thermobalance in 1 atm of oxygen during 48 h at temperatures ranging from 1000 to 1200°C. For each case, it was verified that the scale is made of a-alumina. On the ODS MA 956 alloy, the scale is compact with very well defined grains as shown in (Fig. 2), while the scale formed on the Imphy alloy is convoluted and contains many porosities, (Fig. 3), even in the case of the Y-doped alloy (Fig. 4). As will be seen later see (see Table 3), the scale thickness is about 0.5 mm after oxidation at 1000°C and between 3 and 6 mm after oxidation at 1200°C. After the oxidation, a thin layer of the tracer 54Cr was deposited by evaporation. The profile of chromium, before oxidation, is given in Fig. 5 and indicates that the chromium film thickness is :60 nm. Then, the diffusion treatment was performed in a temperature range 1000 – 1200°C in 1 atm oxygen pressure. Thus, the alumina scales were cooled at 300°C h − 1 between oxidation and diffusion. According to Tsaı¨ et al. works [14,15], it can be said that such a procedure does not induce any damages in the alumina scales. For oxygen diffusion, after the oxidation treatment in 1 atm of 16O during 48 h, without any intermediate cooling, diffusion tests with 18O were established in 0.1 atm during 1 and 2.5 h for MA 956 and Imphy alloy respectively, in the temperature range 1000 –1200°C. The penetration profiles of 18O and 54Cr were established by SIMS (Secondary ion mass spectroscopy, CNRS Bellevue) using a 10 keV Cs + ion source. The scanned area was 250 × 250 mm, and the analysed zone was 60 mm in diameter. The sputtering rate was determined by measuring the crater depth with a profilometer. 3. Oxidation kinetics and morphology of the scales Representative kinetics curves for the oxidation of the two types of alloys at various temperatures are shown in Fig. 6. The determination of the parabolic rate constants kp (g2 cm − 4 s − 1) was carried out by using a complete law analysis:

t= M0 + M1

  m m + M2 S S

Y2O3

Ti

0.5

0.41

P

0.002

2

with kp =

1 M2

(8)

m the weight gain and S the sample surface, which leads to a better fit of the experimental values than the classical parabolic analysis (m/S)2 = kpt, as shown by Fig. 7 which gives an example of comparison between complete and parabolic law analysis. From these determinations, all kc (cm2 s − 1) values, defined by (xox)2 = kct with xox the alumina scale thickness, were calculated with the relation: kc =





MAl2O3 2 k 3MOrAl2O3 p

(9)

after having verified that, in all cases, the scale is made up of a-alumina. kc and xox values are reported in Tables 2 and 3 respectively, and plotted in an Arrhenius diagram in Fig. 8. These kc values are of the same order of magnitude as given in the literature [23]. These results show firstly, that the growth of the scale is predominantly controlled by a diffusionnal process, though a chemical process at one of the interfaces with a linear rate constant kl = 1/M1 (see Fig. 7b), is also controlling the growth rate, at least during a first oxidation period. Secondly, it can be remarked that the weight gain is greater in the case of the Y-implanted Imphy alloy than in the case of the unimplanted Imphy alloy (but Zr-doped) or the ODS alloy, and the difference increases with the temperature, see (Figs. 6 and 8). This indicates that the incorporation of yttrium in a metallic form instead of an oxide dispersion has not a beneficial effect on the oxidation rate at high temperatures, at least when zirconium is present. It was already observed that the simultaneous presence of zirconium and yttrium, or a too high active element level are not beneficial for the high temperature resistance of alloys [10,24]. The good effect of the Y2O3 oxide dispersion was already shown [9,11] and it appeared that the benificial effect depends on the oxide dispersion amount which must be higher than a critical value given as equal to 0.02% in order to form a diffusion barrier [25]. In the case of the MA956 alloy, this critical value is reached. The beneficial effect of yttrium as an oxide dispersion especially consists in improving the scale adherence. This is also the case here. The SEM observations and analyses by EDX show that the scale developed on the Imphy alloy either

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Fig. 2. SEM micrographs of the outer surface (a) and cross section (b) of the MA956 alloy.

implanted or unimplanted is convoluted and contains a lot of porosity (Figs. 3 and 4). In the case of the ODS alloy, the scale is very compact with well defined grains (see Fig. 2). Fig. 2a, Fig. 3a and Fig. 4a are relative to the oxide outer surface morphology while Fig. 2b, Fig. 3b and Fig. 4b show cross sections of the scale. It appears that the average grain size is somewhat different between the two types of alloys: : 1 mm for the alumina scale on the Imphy alloys, and : 3 mm for the alumina scale on the ODS alloy. Moreover, the alumina scales formed on the ODS alloy consist of two distinctive layers, an inner one with columnar grains and a thin outer layer with small equiaxed grains. In the case of the alumina scale grown on the ODS alloy, yttrium is detected in the scale on both parts, i.e. near the outer surface and at the inner interface, with a decrease of the yttrium content in the intermediate part of the scale. Titanium is localised at the inner interface of the alumina scale.

4. Diffusion results

4.1. Anionic self-diffusion Firstly, anionic diffusion was studied in the alumina scale developed on the Imphy unimplanted and Y-implanted alloys. The results of the oxygen diffusion coefficients in the alumina scale developed on the unimplanted and Y-implanted Imphy alloys are summarised in Tables 4 and 5 and Fig. 9a and b, respectively. The ratio Dgb/Db only varies between 101 and 103 and, in the case of the scale formed on the Y-implanted alloy, the activation energy of the grain boundary diffusion is greater than that

obtained for bulk diffusion. This was already observed for diffusion in massive oxides [26,27]. If the order of magnitude of the activation energy is estimated, it gives : 410 and 290 kJ mol − 1 for bulk diffusion and :500 and 570 kJ mol − 1 for grain boundary diffusion in unimplanted and implanted cases respectively. It is observed that the activation energy of grain boundary diffusion is greater than the activation energy for bulk diffusion, especially in the case of the implanted alloy. Such curious comparison of the activation energy of bulk and grain boundary diffusion was already obtained in ceramics [17,26]. In the case of the Y-implanted alloy, the bulk diffusion is slightly greater than in the unimplanted alloy but the grain boundary diffusion is slightly smaller. The differences are very small. By comparing the present results with those extrapolated from Prot and Le Gall works [15–17] in undoped and Y-doped massive aluminas, respectively (values given in Table 6 and in Fig. 10), it can be remarked that both bulk and grain boundary diffusion coefficients in scales are greater by about five orders of magnitude than those in massive aluminas. Nevertheless, it can be seen that, due to the high activation energy of grain boundary diffusion found by Prot, Le Gall et al. at higher temperatures than our temperature range, the extrapolation leads to curious results as the grain boundary diffusion coefficients are of the same order of magnitude as the bulk diffusion coefficients at low temperatures (1000°C). So, comparisons with such extrapolated coefficients must be taken with caution. If our bulk diffusion coefficients are now compared with other literature data concerning oxygen bulk diffusion in massive aluminas [8,12,15,16,28–31], (Fig. 11), a good agreement is obtained only with the extrapolated

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Fig. 3. SEM micrographs of the outer surface (a) and cross section (b) of the unimplanted Imphy alloy.

Fig. 4. SEM micrographs of the outer surface (a) and cross section (b) of the Y-implanted Imphy alloy.

values from Oishi and Kingery works [28] who determined diffusion coefficients on monocrystalline alumina. In fact, it was shown that the so-determined diffusion coefficient corresponds to an apparent diffusion coefficient taking into account both bulk diffusion and diffusion in dislocations which were introduced in great quantities during production. All other values are significantly smaller than ours. The present results can also be compared with those obtained on a-alumina scales. But, for those obtained by Clemens et al. on MA956 [11], the scale morphology notably differs from the morphology in Figs. 3 and 4 and is rather close to the morphology shown in Fig. 2.

So, the comparison does not seem to be adequate. Secondly, it is possible to compare our results with those obtained by Balmain et al. [12] in the case of alumina grown on b-NiAl. In this case, the scale morphology is close to that shown in (Figs. 3 and 4). Consequently, our results are relatively close to those obtained by Balmain (Fig. 11) who also found bulk and grain boundary diffusion coefficients much greater than those determined in massive aluminas.

4.1.1. MA 956 alloy The anionic self-diffusion coefficients in the alumina scale developed on MA956 are reported in Table 7 and

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plotted in Fig. 12. In this case, the ratio Dgb/Db is about 104 –105 and the activation energy of bulk and grain boundary diffusion have fairly the same order of magnitude, i.e. : 370 and 390 kJ mol − 1 for bulk and grain boundary diffusion respectively. Even if the remark made previously about the extrapolation of diffusion results obtained at higher temperatures is always valid, the comparison of the present diffusion results with those extrapolated [15 – 17] offers the advantage to indicate (Fig. 13 and compare results of Tables 6 and 7) that, in this case, there is agreement between bulk diffusion coefficients in Y-doped massive alumina and the scale formed on MA956. It also means that the bulk diffusion coefficients in the scale formed on MA956 are not far from most of the results given in the literature for oxygen bulk diffusion (see Fig. 11). Our grain boundary diffusion coefficients are still greater than those extrapolated from results obtained on massive aluminas. As shown by (Fig. 14a), the bulk diffusion coefficients in alumina formed on MA956 are smaller than

Fig. 7. Comparison between parabolic law (a) and complete law (b) analysis. Unimplanted Imphy alloy oxidised at 1200°C.

Fig. 5. Profile of chromium after oxidation and before diffusion test.

Fig. 6. Kinetics curves for all alloys at 1000, 1100 and 1200°C.

those determined in alumina formed on Imphy alloys. The difference is not so marked if the grain boundary diffusion coefficients are compared (Fig. 14b). Now, if our results are compared with those obtained recently by Clemens et al. on an alumina scale formed on the same alloy than in the present work [11], and whose values are reported in Table 8 and (Fig. 15), it appears that their bulk diffusion coefficient deduced from the first domain of the experimental penetration profile corresponds to our effective diffusion coefficient. So, it is suggested that they did not determine bulk diffusion coefficients but effective diffusion coefficients. Due to this reason, their grain boundary diffusion coefficient are greater than ours. Moreover, the difference between their grain boundary diffusion coefficients and ours is also due, at least in part, to the fact that they considered the scale growth exclusively controlled by oxygen grain boundary transport and that they did not take into account the surface roughness.

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Table 2 Experimental kc (cm2 s−1) determined on the alloys at different temperatures in 1 atm of oxygen T (°C)

Unimplanted-Imphy Y-implanted-Imphy ODS alloy

1000

1100

1200

1.8×10−14 1.8×10−14 1.6×10−14

2.4×10−13 4.7×10−13 2.3×10−13

1.1×10−12 2.1×10−12 5.0×10−13

Table 3 Calculated thickness values (mm) of the oxide scale developed on each alloy at 1000, 1100 and 1200°C T (°C)

Unimplanted-Imphy Y-implanted Imphy ODS alloy

1000

1100

1200

0.5 0.5 0.5

2.0 2.8 2.0

4.3 6.0 3.0

4.2. Cationic self-diffusion Taking into account the rather surprising results obtained for anionic diffusion in the scale formed on Imphy alloy, while diffusion results in the scale developed on MA956 seems satisfactory, the cationic self diffusion coefficients were only determined in the scale grown on the ODS alloy. Similarly to anionic diffusion, the penetration profiles were analysed considering the first domain related to an effective diffusion. The results are reported in Table 9 and (Fig. 16). The ratio Dgb/Db is about 103 – 104, thus slightly smaller than for anionic diffusion in the same scale (see Fig. 12) and the activation energy of grain boundary diffusion ( :430 kJ mol − 1) is greater than

that for bulk diffusion (: 250 kJ mol − 1), which was not the case for anionic diffusion. Always comparing cationic and anionic self diffusion, (Fig. 17a), it clearly appears that the cationic self diffusion is more important than the anionic one, especially for the bulk diffusion. Clemens et al. [11], in their analysis of the penetration profiles, neglected the cationic diffusion and attributed the 18O concentration in the outer part of the scale to outward transport of Ti and Y and not to outer aluminium diffusion. But, the present results indicate that cationic diffusion cannot be neglected and this is confirmed by the results of both Moya et al. for Cr diffusion [18] and Le Gall et al. [14] for aluminium diffusion, in undoped alumina single crystals (Fig. 17b and Tables 10 and 11). Moya et al. found that chromium diffuses slightly faster than aluminium and Le Gall found that aluminium diffusion is slightly faster than oxygen diffusion. The comparison of our results with those obtained by Balmain et al. [12], (Fig. 17b), does not show important differences. This is curious as the alumina scale formed on NiAl (Balmain works) has not a morphology similar to that obtained on MA956.

5. Comparison between the experimental oxidation constants and the constants calculated with diffusion data The oxidation constant was calculated according to Wagner’s theory, i.e. using the following equation: kc(cal)=

&

pext O

2

p int O2

[Danion + (b/a)Dcation]fln pO2

int for an oxide Ma Ob, with p ext O2 and p O2 the oxygen pressure at the outer and inner interfaces of the scale, respectively, taken as the equilibrium pressures. If it is considered that the concentration of the defects which ensure the scale growth by diffusion does not depend on the oxygen pressure, this equation leads to: int kc = [Danion + (b/a)Dcation]ln(p ext O2 /p O2) ext O2

Fig. 8. Experimental kc (cm2 s − 1) values for each alloy at 1000, 1100 and 1200°C.

(10)

(11)

In our case, a= 2 and b= 3, p = 1 atm (oxygen pressure of the oxidation experiments) and p int O2 is calculated considering that the thermodynamic equilibrium is reached at the inner interface. The value of the ratio int ln(p ext O2 /p O2) for each temperature is given in Table 12. This assumption is based on the work of Ramanarayanan et al. [32] and on the fact that it is wellknown that alumina behaves as an extrinsic material: the defect concentration is imposed by the impurities. The calculated values of the oxidation constants, considering several mechanisms for the scale growth (bulk, grain boundary or both, i.e. taking Deff), are given in Table 13 for the scale on MA 956 alloy.

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Table 4 Anionic Deff, Db and Dgb in Al2O3 scale developed on the unimplanted Imphy alloy T (°C)

f

f

fr

Deff (cm2 s−1)

(−Pgb)

Db(cm2 s−1)

Dgb(cm2 s−1)

1000 1100 1200

0.3 0.6 0.8

0.01 0.005 0.00375

1.30×10−3 1.15×10−3 1.07×10−3

7.5×10−18 1.1×10−16 1.8×10−15

1.0×107 1.3×106 1.1×106

6.5×10−18 4.3×10−17 1.4×10−15

7.4×10−16 6.1×10−14 4.6×10−13

Table 5 Anionic Deff, Db and Dgb in Al2O3 scale developed on the Y-implanted Imphy alloy T(°C)

f

f

fr

Deff (cm2 s−1)

(−Pgb)

Db(cm2 s−1)

Dgb (cm2 s−1)

1000 1100 1200

0.25 0.5 0.75

0.11 0.006 0.004

1.33×10−3 1.20×10−3 1.09×10−3

3.8×10−17 1.3×10−16 2.2×10−15

3.1×107 2.6×106 1.3×106

3.8×10−17 1.0×10−16 1.7×10−15

2.8×10−16 2.8×10−14 3.9×10−13

In the case of Imphy alloys, for which only anionic self-diffusion coefficients were determined, it was assumed that the cationic self-diffusion is of the same

order of magnitude as the anionic diffusion. Thus, the values of the oxidation constants, calculated by taking into account the effective diffusion coefficient, are given in Tables 14 and 15 for unimplanted and Y-implanted alloy respectively. For the scale developed on MA 956, it appears that the experimental values are in agreement with the calculated values if both bulk and grain boundary diffusion mechanisms are considered, Fig. 18a. In the case of Imphy alloys, the experimental parabolic oxidation constants are higher than the calculated values (Fig. 18b) as it was also for the results obtained by Balmain et al. [12].

6. Discussion

6.1. Good beha6iour of the ODS alloy The results obtained with the ODS alloys are satisfactory:

Table 6 Extrapolated anionic self-diffusion coefficients in the undoped and Y-doped massive aluminas [15 – 17] T (°C)

Undoped massive alumina

Fig. 9. (a) Db and Dgb in the scale developed on the unimplanted Imphy alloy. (b) Db and Dgb in the scale developed on the Y-implanted Imphy alloy.

Y-doped massive alumina

1000

1100

1200

Db (cm2 s−1)

1.6×10−24

1.3×10−22

5.6×10−21

Dgb (cm2 s−1)

2.5×10−22

1.4×10−19

3.4×10−17

Db (cm2 s−1)

1.6×10−21

3.4×10−20

4.6×10−19

Dgb (cm2 s−1)

1.0×10−22

2.5×10−20

2.9×10−18

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scale formed on the ODS alloy (see Fig. 2), the grains are equiaxed, the scale is flat and compact, the inner interface is also flat. There is some limited internal oxidation. The alumina scale formed on the Imphy alloys (Figs. 3 and 4) is very convoluted, buckled. The scale is not compact and the inner interface shows important waves suggesting the presence, during the scale growth, of mechanical instabilities. The consequence of the differences in the morphology is confirmed by ionic images of 18-oxygen (Fig. 19), made on cross sections of samples oxidised first in 16O2 (48 h), then in 18O2 (: 1–2 h). In (Fig. 19a), relative to a treatment of the ODS alloy, the 18O-isotope is concentrated in the scale and somewhat in the internal oxide particles, indicating that oxygen has diffused by both bulk and grain boundary diffusion. On the contrary, in the case of the Imphy alloy (here doped with yttrium, but results are similar without yttrium as a doping element), the 18O has very quickly gone through the scale and diffused in the underlying substrate to a large distance. It means that the scale does not act as a protective barrier and that probably gaseous oxygen has reached the inner interface via porosities and then has diffused into the substrate. In such a case, an analysis of the penetration profiles using diffusion equations in a solid medium is not rigorous and comparisons with other diffusion results, or oxidation constant calculations are not adequate. Wagner’s theory cannot be used in such cases. Now, the question is why such various morphologies are obtained for a-alumina scales? This will be treated in Section 6.3, after the diffusion results have been commented.

6.2. Diffusion in protecti6e alumina scales Fig. 10. (a) Diffusion coefficients in the scale developed on the Imphy alloy (unimplanted with yttrium) and the extrapolated diffusion coefficients in undoped massive alumina [15]. (b) Diffusion coefficients determined in the scale developed on the Y-implanted Imphy alloy and the extrapolated diffusion coefficients in Y-doped massive alumina [16].

1. The ratio Dgb/Db has a satisfying order of magnitude, 2. There is a good agreement between the bulk diffusion coefficients determined in massive alumina and in alumina scales, 3. The oxidation constants calculated on the basis of the diffusion results agree well with the experimental ones. It is not the case with the diffusion results obtained in the case of the Imphy alloys. The differences in the behaviour are directly related to the scale morphology. In the case of the alumina

As it was shown previously that the diffusion results obtained in the alumina scale formed on the ODS alloy are available, some conclusions can be drawn. 1. The first important point is that the cationic diffusion is more important than the anionic diffusion, contrary to assumptions made in the literature. 2. This result agrees with those obtained in massive aluminas: Le Gall [14] observed, for bulk diffusion, that aluminium diffuses faster than oxygen. 3. It means that, in the calculation of the oxidation constant (for compact scales), the cationic self diffusion cannot be neglected. 4. Always in the calculation of the oxidation constant, it is clear that the relevant diffusion coefficient is an effective one. 5. The growth of compact a-alumina scales occurs by counter-current diffusion of oxygen and aluminium via both the bulk and the grain boundaries which act as short-circuits.

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Fig. 11. Comparison between our results and those given by the literature for oxygen bulk diffusion in alumina scales and in massive aluminas.

6. The agreement between the calculated and the experimental oxidation constants, in the case of a compact scale, confirms the analysis of the penetration profiles. In other words, it confirms the fact that the first part of the penetration profiles is related to effective diffusion and that the surface roughness has to be considered. 7. It is not possible to discuss about differences between grain boundary diffusion in alumina scales and grain boundary diffusion in massive aluminas as these values have been extrapolated from results obtained at higher temperatures, with a probably different grain boundary chemistry. 8. Evidently, the presence of active elements, such as yttrium or titanium which enter the composition of the ODS alloy, must have an effect on the diffusion. This will be discussed just below.

6.3. Why the differences in the morphology of the alumina scales Several ideas can be developed to discuss the morphology differences of alumina scales. First, some differences can be related to the possible formation of the transition aluminas. Also the influence of the doping elements cannot be neglected, not only on the transport properties, but also and probably especially on the scale morphology and on its mechanical properties.

Concerning the possible formation of transition aluminas, it looks like if this problem especially concerns materials such as intermetallics, particularly b-NiAl. Indeed a lot of works on the oxidation of b-NiAl mentions the formation of transition aluminas [2,10,12,33–36], while it is not the case of literature data on oxidation of Fe-Cr-Al alloys. For instance, Le Coze found that after 10 min oxidation of Fe-Cr-Al alloys at 1050°C, a-alumina was present, which is not the case for Fe3Al treated in the same conditions (J. Le Coze, personal communication). Prunier, (V. Prunier, personal communication) similarly found on Kanthal A1 a gap in kp values around 1000°C probably related to a change in oxidation mechanism. It means that the growth of transition aluminas on Fe-Cr-Al alloys can only be evidenced at low temperatures (B1000°C). Thus, the present work is at the limit of the possibilities of observation of transition aluminas. It is suggested that chromium accelerates the formation of a-alumina, as it was shown when chromium is added to b-NiAl [33]. One could think that the difference in the oxidation weight gain between the undoped and Y-doped Imphy alloy could be due to the effect of yttrium on the transition aluminas. Indeed, it was shown that metallic yttrium delays the transformation of transition aluminas to a-alumina [35], which leads to a greater weight gain due to the fact that the growth rate of transition aluminas is greater than the growth rate of a-alumina

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258

Table 7 Deff, Db and Dgb anionic self-diffusion coefficients in Al2O3 scale developed on the MA 956 alloy T (°C)

f

Deff (cm2 s−1)

(−Pgb)

Db (cm2 s−1)

Dgb (cm2 s−1)

1000 1100 1200

0.03 6.60×10−3 6.42×10−3

4.1×10−18 3.8×10−17 1.3×10−16

4.2×106 5.2×105 9.0×105

4.7×10−21 7.9×10−21 5.9×10−19

1.4×10−16 5.8×10−15 2.0×10−14

[31]. But, such a phenomenon leads to a greater weight gain only during a given oxidation time. For instance, J. Balmain observed for b-NiAl that, at 1100°C, the first oxidation stage related to the growth of transition aluminas extends on 3 – 4 h maximum. In the present work, the differences observed in the oxidation kinetic curves of undoped and Y-doped Imphy alloys concern all the oxidation treatment (see Fig. 6). In other words, in the present work, the Y-doped Imphy alloy shows only one oxidation stage whose parabolic oxidation

Fig. 12. Bulk and grain boundary diffusion coefficients in the scale developed on the MA 956 alloy.

Fig. 13. Comparison between diffusion coefficients in massive Ydoped alumina [16,17] and alumina scale developed on the MA 956 alloy.

constant is greater than that obtained with the undoped alloy (Table 2). Thus, our feeling is that phase transformation does not occur (or is not an important parameter) in the oxidation conditions used in the present work. Now, it is clear that the active elements must have an influence on the transport properties in the alumina scales. An extensive literature concerns this problem, as well on massive aluminas (conductivity, diffusion, creep, microstructure, EXAFS studies...) as on alumina scales (isotopic exchange tests, conductivity measurements, diffusion...). Nevertheles, several considerations suggest that these effects are not so important for the growth rate of alumina scales: 1. The order of magnitude of the growth rate of alumina scales is always the same (at a given temperature), which suggests that the effect of impurities or doping element is of minor importance, 2. This is confirmed by the small differences in the diffusion coefficients in undoped and Y-doped aluminas [15–17], 3. This is also confirmed by the similar behaviour of the ODS alloy, doped with Y2O3, and the Imphy alloy (doped with Zr): the kinetic curves are similar to each other although the two alloys are very different (see Fig. 6), 4. When differences are discussed by taking into account the active element effect, all possibilities can be found [1]: authors describe the active element effect as decreasing either the oxygen diffusion or the aluminium diffusion.... Also, it can be noted that active elements are sometimes localised at the inner interface, sometimes near the outer interface, etc... Thus, Pint’s model [37] as well as Pieraggi’s model (poisoned interface) [38] are perhaps operative in some cases, but not always and it seems that, even if active elements can have an effect on the growth rate, this is of minor importance. Nevertheless, it is perhaps interesting to mention some ideas deduced from observations in the present work: 1. Synergetic phenomena occur. Thus, the effect of Y+ Zr is rather bad while each of these elements are potentially considered as able to improve the oxidation resistance;

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259

Fig. 15. Comparison between our results and those obtained by Clemens et al. [11].

forms yttrium oxide immediately at the beginning of the oxidation. The Y2O3 oxide dispersion in the alloy leads to a decrease of the oxidation rate probably related to its stability. Clearly, the present work indicates that active elements have an important influence on the morphology of the scale and on the mechanical properties. Table 9 Deff, Db and Dgb cationic diffusion coefficients in Al2O3 scale developed on the MA 956 alloy D(cm2 s−1) Fig. 14. (a) Comparison between bulk diffusion coefficients in the scale developed on unimplanted and Y-implanted Imphy alloys and MA 956 alloy. (b) Comparison between grain boundary diffusion coefficients in the scale developed on the unimplanted and implanted Imphy alloys and MA 956 alloy.

T (°C)

Deff

(−Pgb)

Db

Dgb

1000 1100 1200

4.4×10−17 4.0×10−16 2.2×10−15

7.4×106 1.9×106 7.0×105

3.0×10−18 4.7×10−17 6.8×10−17

1.4×10−15 5.4×10−14 3.3×10−13

Table 8 Anionic self-diffusion, Db and Dgb determined by Clemens et al. [11] in an Al2O3 scale developed on the MA 956 alloy T (°C)

Db (cm2 s−1) Dgb (cm2 s−1)

1000

1100

1200

1.3×10−13

1.0×10−17 1.0×10−12

5.7×10−12

2. Yttrium, when added in a metallic form, has not the same influence as yttrium oxide additions. Indeed, the scale developed on the ODS MA 956 alloy is more compact and protective than the scale developed on the Imphy alloy. The metallic yttrium, in the present work, increases the oxidation rate, perhaps because it is not thermodynamically stable and

Fig. 16. Cationic bulk and grain boundary diffusion coefficients in the scale developed on the MA 956 alloy.

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scale formed on the Imphy alloy (Fig. 3b and Fig. 4b) are representative of mechanical instabilities. Moreover, the active elements or impurities incorporated in the scales can have an effect on the intrinsic mechanical properties of the scale, as it was shown for instance for yttrium doping in massive alumina [40,41], but it is not easy to discuss this point as it would be necessary to take into account the amount and the chemical nature of each element in all parts of the scale (influence of the oxygen pressure). Finally, one could suggest the formation of an intermediate film at the inner interface which would have mechanical properties intermediate between those of the substrate and of the scale, but the present work did not allow to observe it.

7. Conclusion Diffusion characteristics in alumina scales are of great interest in determining the growth mechanism of such scales at high temperatures. Thus, anionic and cationic diffusion tests were performed on alumina scales developed either on a traditional Fe-Cr-Al alloy undoped or Y-doped by implantation or on an ODS alloy. At first, the oxidation behaviour of these alloys was characterised and the parabolic oxidation constants Table 10 Extrapolated cationic self-diffusion coefficients in undoped massive alumina [14] T (°C)

Db (cm2 s−1) Fig. 17. (a) Comparison between D(O) and D(Cr) in the scale developed on the MA 956 alloy. (b) Comparison between our results concerning cationic diffusion in the scale developed on the MA 956 alloy and those given by the literature for cationic diffusion.

Several works have already indicated that active elements can modify the grain size of the alumina scale. In many cases, as for massive aluminas [39], it is mentioned that yttrium induces a decrease of the grain size of the alumina scales and such a phenomenon will increase the creep rate and promote the stress relaxation. But, this point cannot be discussed here for the ODS alloy since there is no similar alloy without yttrium. When stress relaxation is not possible, i.e. when deformation of the substrate is not operative, then, the stresses in the scale can be relaxed by buckling, as observed for the Imphy alloy and the scale is not compact and contains cracks and pores. Besides, the waves observed at the inner interface of the alumina

1000°C

1100°C

1200°C

1.8×10−22

6.2×10−21

1.3×10−19

Table 11 Db for Cr determined by Moya et al. [18] in the undoped alumina single crystals T (°C)

Db (cm2 s−1)

1000°C

1100°C

1200°C

1.6×10−18

7.7×10−18

3.6×10−17

Table 12 int Ratio ln(p ext O2 /p O2) for each temperature in the alumina scale grown in 1 atm O2 T (°C)

p int O2 int ln(p ext O2 /p O2)

1000

1100

1200

10−35 78.28

10−32 71.38

10−28 62.17

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Table 13 Comparison between the calculated parabolic constants assuming various mechanisms of scale growth and the experimental ones, for the MA 956 alloy T (°C)

1000 1100 1200

Diffusion mechanism Lattice diffusion Db (cm2 s−1)

Grain boundary Dgb (cm2 s−1)

Effective diffusion Deff (cm2 s−1)

Experimental kc (cm2 s−1)

3.5×10−16 5.0×10−15 6.3×10−15

1.7×10−13 6.1×10−12 3.2×10−11

5.4×10−15 5.0×10−14 2.1×10−13

1.6×10−14 2.3×10−13 5.0×10−13

were determined in the temperature range 1000– 1200°C. The scale morphology was observed by SEM and analyses were performed to determine the localisation of some doping elements. It appears that the diffusion coefficients are closely connected to the morphology of the scale. The alumina scale formed on the classical alloy without yttrium is largely buckled, convoluted with an interface metal/oxide very disturbed, with regular waves. On the Y-implanted alloy, the scale is not at all compact. In such cases, the present work clearly indicates that the penetration profiles cannot be analysed using classical diffusion equations. The scale developed on the ODS alloy, is made of compact alumina grains. Then, the penetration profiles can be analysed in terms of classical diffusion. Considering that the first part of the profile is related to an effective diffusion coefficient, and taking into account the oxide surface roughness, the anionic and cationic diffusion coefficients are in agreement with those in massive aluminas, and lead to agreement between the experimental and the calculated parabolic oxidation constants. This confirms the validity of the analysis.

The cationic diffusion coefficients are greater both in the bulk and in the grain boundary than the anionic ones. This indicates that, contrary to some literature data, cationic diffusion cannot be neglected in the alumina growth mechanism.

Table 14 Comparison between the calculated parabolic constants and the experimental ones for the Imphy alloy unimplanted with yttrium Temperature (°C) 1000 Calculated kc (cm2 s−1) Experimental kc (cm2 s−1)

1100

1200

1.0×10−16 1.6×10−15 2.6×10−14 1.8×10−14 2.4×10−13 1.1×10−12

Table 15 Comparison between the calculated parabolic constants and the experimental ones for the Y-implanted Imphy alloy Temperature (°C) 1000 Calculated kc (cm2 s−1) Experimental kc (cm2 s−1)

1100

1200

1.5×10−15 2.0×10−14 2.9×10−13 1.8×10−14 2.4×10−13 1.1×10−12

Fig. 18. (a) Comparison between kc (exp) and kc (cal) assuming various mechanisms for the alumina scale growth on the MA 956 alloy. (b) Calculated and experimental values of kc in the case of alumina scales developed on unimplanted and Y-implanted Imphy alloys.

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Fig. 19. (a) SIMS 18O image on the cross section of a MA 956 sample oxidized 48 h in 16O then 1 h in 18O. (b) SIMS 18O image on the cross section of a Imphy sample oxidized 48 h in 16O then 1 h in 18O.

The differences in the behaviour of the various alloys and in the morphology of the scales are discussed taking into account various parameters. It is thought that perturbations with transitional aluminas are not operative for these alloys at the test temperatures. Again, the effect of active elements is not easy to discuss and would need to have information on their amount, their chemical nature and their localisation.... Anyway, it is clear that a beneficial effect of yttrium in the oxidation resistance is obtained especially when this element is added as an oxide dispersion Y2O3. It is suggested also that these active elements rather act on the mechanical properties of both the scale and the substrate than on the transport properties.[39 –41]

Acknowledgements C. Dolin is gratefully acknowledged for his investigation to carry out the diffusion profiles by SIMS at CNRS Bellevue.

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