High temperature oxidation of Fe–Al and Fe–Cr–Al alloys: The role of Cr as a chemically active element

Corrosion Science 52 (2010) 3394–3404

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High temperature oxidation of Fe–Al and Fe–Cr–Al alloys: The role of Cr as a chemically active element E. Airiskallio a,b, E. Nurmi a,b,c,*, M.H. Heinonen a,b, I.J. Väyrynen a,b, K. Kokko a,b, M. Ropo d,e, M.P.J. Punkkinen a,b,f, H. Pitkänen g, M. Alatalo g, J. Kollár h, B. Johansson f,i, L. Vitos f,h,i a

Department of Physics and Astronomy, University of Turku, FI-20014 Turku, Finland Turku University Centre for Materials and Surfaces (MatSurf), Turku, Finland Graduate School of Materials Research, Turku, Finland d Department of Information Technology, Åbo Akademi, FI-20500 Turku, Finland e Fritz-Haber-Institut der Max-Planck-Gesellschaft, D-14195 Berlin, Germany f Applied Materials Physics, Department of Materials Science and Engineering, Royal Institute of Technology, SE-10044 Stockholm, Sweden g Department of Mathematics and Physics, Lappeenranta University of Technology, P.O. Box 20, FI-53851 Lappeenranta, Finland h Research Institute for Solid State Physics and Optics, P.O. Box 49, Budapest H-1525, Hungary i Department of Physics and Materials Science, Uppsala University, SE-75121 Uppsala, Sweden b c

a r t i c l e

i n f o

Article history: Received 20 April 2010 Accepted 18 June 2010 Available online 30 June 2010 Keywords: A. Alloy A. Steel B. AES C. Oxidation C. Passive films

a b s t r a c t Good high-temperature corrosion resistance of Fe–Al alloys in oxidizing environments is due to the aAl2O3 film which is formed on the surface provided temperature is above 900 °C and the Al-content of the alloy exceeds the critical value. Ab initio calculations combined with experiments on Fe–13Al, Fe– 18Al, Fe–23Al and Fe–10Cr–10Al alloys show that the beneficial effect of Cr on the oxidation resistance is significantly related to bulk effects. The comparison of experimental and calculated results indicates a clear correlation between the Fe–Cr chemical potential difference and the formation of the protective oxide scales. Ó 2010 Elsevier Ltd. All rights reserved.

1. Introduction Fe–Cr–Al alloys are often used as high-temperature corrosion resistant materials [1–3] due to their ability to form a highly stable and protective oxide scale on the open surface when exposed to oxidizing environment. The physical properties and corrosion resistance of Fe–Cr–Al as a function of the chemical composition of the alloy has been studied quite extensively, see e.g. [4–20]. Although the overall effects and phenomena of the oxidation and corrosion of metals has been well characterized [21,22], the atomic processes at successive time steps of the oxidation or even the chemical composition of the surface prior to the oxidation are less well studied. The purpose of our investigation is to extend the understanding of the Fe–Cr–Al surfaces towards the atomic scale. The knowledge obtained can be used to shed more light on the state of the surface prior to oxidation. This will help us build a bot-

* Corresponding author at: Department of Physics and Astronomy, University of Turku, FI-20014 Turku, Finland. Tel.: +358 2 33 5661; fax: +358 2 333 5070. E-mail address: eknurm@utu.fi (E. Nurmi). 0010-938X/$ - see front matter Ó 2010 Elsevier Ltd. All rights reserved. doi:10.1016/j.corsci.2010.06.019

tom up picture of the physical phenomena responsible for the excellent high temperature oxidation resistance of Fe–Cr–Al alloys. The corrosion resistance of Fe–Cr–Al in oxidizing atmosphere is due to the formation of a highly protective chromium and aluminum oxide layer on the surface, which effectively separates the oxidizing atmosphere from the pure alloy. Because a-Al2O3 is very stable at high temperatures it would be beneficial to maximize its content at the surface. Unfortunately, for most of the Fe-alloy applications the most straightforward procedure of increasing the Al-content in bulk is not an acceptable solution. This is due to the fact that increasing Al content makes Fe–Al alloys brittle, which poses a natural upper bound for the Al-content in these alloys regarding to most of the applications. [23] Fortunately, there exist additional alloying elements that can boost the formation of the Al-oxide scale on the surface up to such a level that the Al-content in bulk can be kept within the acceptable limits regarding to the mechanical properties of the alloy. One of the most common additional elements for this purpose is Cr. This phenomenon, called the third element effect (TEE), is still considered a phenomenon without generally accepted explanation [20,24–28]. In the present investigation we concentrate on the surface phenomena closely

E. Airiskallio et al. / Corrosion Science 52 (2010) 3394–3404

related to TEE. To give an extensive account on the energetics of the pure Fe–Cr–Al surfaces at the atomic resolution we have calculated the differences of the atomic chemical potentials of the components in the bulk and at the (0 0 1) surface of Fe–Cr–Al. The excellent corrosion resistance of Fe–Cr–Al in oxidizing environments is based on the rapid formation of the oxide surface scale of the right type. Iron oxide on the surface is quite vulnerable to corrosion in most cases. Chromium oxide provides very good corrosion protection, but only at low and intermediate temperatures. At low temperatures Cr improves the oxidation resistance of Fe–Cr–Al due to the fast formation of Cr2O3 compared to that of a-Al2O3 [29,30]. Since at high temperature Cr2O3 transforms to volatile compounds, e.g. CrO3 [31], the corrosion protection at elevated temperatures requires the more stable aluminum oxide scales on the surface. Depending on the concentrations of the alloys, details of the oxidizing environment, as well as criteria used for the oxidation resistance, the literature references show marked variance in the upper limits of the temperature allowing the protective oxide scales of Cr2O3 and Al2O3 to form and sustain. The most often reported values for the upper limits of the feasible temperature range are 900– 1100 °C for Cr2O3 and 1200–1400 °C for Al2O3 [3,7,25,32,33,31,34–36]. Depending on the composition of the Fe–Cr–Al ternary alloy different types of oxide scales are formed on the open surface when exposed to the oxidizing environments. Because the right composition of the alloy is one of the key prerequisites for a good corrosion protection, we have calculated the chemical potential differences of Fe–Cr–Al within a wide range of concentrations, Cr: 0–25 at% and Al: 0–20 at%. Detailed understanding of the equilibrium surface composition requires knowledge of various properties of a material, e.g. vacancy formation, migration, diffusion, lattice relaxation, thermodynamics, kinetics, etc. From this diverse field of phenomena we concentrate here on the energetic driving force for the surface segregation and its consequences on the formation of the protective oxide scale on Fe–Cr–Al alloys. To investigate experimentally the high temperature oxidation resistance of Fe–Cr–Al alloys as a function of Cr- and Al-contents four different alloys were prepared by induction melting. The alloys prepared, Fe–13Al, Fe–18Al, Fe–23Al, and Fe–10Cr–10Al, (at% used in alloy formulas) cover the interesting concentration range including the Al threshold corresponding to the onset of the high temperature oxidation resistance and the effect of Cr to turn a non-resistant Fe–Al alloy to a high temperature oxidation resistant alloy by enhancing the Al-oxide scale. The samples were oxidized at 1000 °C and the surfaces were investigated by Auger Electron Spectroscopy (AES), Atomic Force Microscopy (AFM), and Optical Microscopy. The experimental data are analysed in the light of computational data to deepen the understanding of the atomic view of the oxidation of Fe–Cr–Al.

2. Computational methods The calculations are based on the density functional theory [37,38] and performed using the Exact Muffin-Tin Orbitals (EMTO) method [39,40]. The EMTO method is an improved screened Korringa–Kohn–Rostoker method [41], where the one-electron potential is represented by large overlapping muffin-tin potential spheres. By using overlapping spheres, one describes the crystal potential more accurately, compared with the conventional nonoverlapping muffin-tin approach [42–45]. For the exchange-correlation density functional the generalizedgradient approximation [46] was used. We have tested earlier that this choice leads to the best overall equilibrium volume for the component metals of Fe–Cr–Al [47,48]. The EMTO basis set included s, p, d, and f orbitals. The one-electron equations were solved

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Fig. 1. The equilibrium Wigner–Seitz radius (rWS in units of Bohr radius) of Fe–Cr– Al alloys as a function of Cr concentration with Al-content as a parameter. For bcc alloys the lattice parameter (a) and the WS radius are related as a = (8p/3)1/3rWS.

within the scalar-relativistic and soft-core approximations. The EMTO Green’s function was calculated self-consistently for 32 complex energy points distributed exponentially on a semi-circular contour, which included states within 1 Ry (2  1018 J) below the Fermi level. In the one-center expansion of the full charge density, we adopted an l-cutoff of 8 and the total energy was calculated using the full charge density technique [43,40]. For each alloy the calculated equilibrium volume was used. The optimized Wigner– Seitz radii are shown in Fig. 1. The convergence of the total energy with respect to the number of k-vectors was tested. It was found that 506 (for bulk) and 231 (for surfaces) uniformly distributed kvectors in the irreducible part of the 3D and 2D Brillouin zones were enough for the present purposes. The alloys were described as substitutionally disordered ferromagnetic bcc alloys [49,50]. Since in the present investigation we map a large concentration region of Fe–Cr–Al with Cr- and Al-concentrations approaching zero the conventional supercell method would require enormously large and numerous supercells. Here, we resolve this difficulty by employing the Coherent Potential Approximation (CPA) [51,52]. Within the CPA, the alloy components are embedded in an effective medium, which is constructed in such a way that it represents, on the average, the scattering properties of the alloy. In this way, the original alloy problem reduces to the Schrödinger equation for the effective medium plus the real space Dyson equations written for each single impurity. In the present application, we adopt the CPA implemented within the frameworks of the EMTO method [39,40]. The EMTO approach has been applied successfully in the theoretical study of various structural and electronic properties of alloys and compounds [40], demonstrating the accuracy and efficiency needed for the present investigation. The semi-infinite bulk surface system is modelled by a slab consisting of 8 atomic layers parallel to the surface. To retain the periodicity of the model system an infinite array of the slabs separated by vacuum layers is constructed. The thickness of the vacuum layers is equivalent to 4 atomic layers. The differences of chemical potentials of a ternary A–B–C alloy are calculated as

lA  lB ¼

 1 dU  xC ¼ constant; N dxA 

ð1Þ

where lA, lB, U, N, xA, and xC are the chemical potentials of A and B components, the internal energy of the system, the number of atoms in the system, and the atomic fractions of the components A and C, respectively. One should note that since we require the

E. Airiskallio et al. / Corrosion Science 52 (2010) 3394–3404

Table 1 Calculated concentration profile of the surface at different temperatures. T (K)

0 100 200 300 400 500 600 900 1200 1500

Fe5Al

Fe5Cr5Al

Fe10Cr5Al

Fe15Cr5Al

(at%) Fe

Al

Fe

Cr

Al

Fe

Cr

Al

Fe

Cr

Al

13 16 19 22 25 27 30 37 43 49

87 84 81 78 75 73 70 63 57 51

1 6 9 12 15 18 21 28 34 39

0 0 0 0 0 0 0 0 0 0

99 94 91 88 85 82 79 72 66 61

0 0 0 1 4 7 10 19 27 33

0 0 0 0 0 0 0 0 1 1

100 100 100 99 96 93 90 81 72 66

0 0 1 3 6 9 12 19 26 32

0 0 0 0 0 0 0 0 0 0

100 100 99 97 94 91 88 81 74 68

total number of atoms of the system considered to be conserved the atomic fractions are related as xA + xB + xC = 1. The chemical potential differences are calculated for the bulk and for the surface. The surface concentrations for the selected alloys as a function of temperature T (Table 1) were obtained by minimizing the grand potential (X) of the system.

X ¼ U  TS 

X

li Ni ;

ð2Þ

i

where S is the entropy of the system, and Ni and li are the total number and the chemical potential of particles of type i. In our calculations only the configurational part was included in the entropy. Details of the calculational method are published elsewhere [53– 55]. While the surface profiles shown in Table 1 are calculated by varying the surface concentrations the rest of the surface calculations in the present paper refer to homogeneous bulk–surface systems. 3. Experimental procedure The alloys were prepared by induction melting under argon. Weighted amounts of pure metals (99.99%) were melted in an alumina crucible. Nominal compositions were Fe–13Al, Fe–18Al, Fe– 23Al and Fe–10Cr–10Al in atomic %. Ingots were slowly cooled to room temperature. Atomic % checked with EDX are shown in Table 2. Slides of about 3 mm thick were cut from the ingots. Samples were grounded with abrasive paper down to 1000 grit to flatten and remove the deformed layer of the surface. Final polishing was done with 6 and 1 lm diamond suspensions. Samples were then cleaned with acetone and methanol in ultrasonic bath, and dried with air flow. Oxidation at 1000 °C was performed under oxygen flow in a quartz tube. Heating of the sample was done by induction and temperature measured with infrared thermometer. Reaching 1000 °C took about 40 s. A PHI 610 spectrometer was used for depth profiling. Base pressure during the measurements was 1  109 Torr. Electron gun was run at 5 keV and 80 nA. Sputtering was done with 3 keV argon ions at an angle of 40 deg and rastered over 4  4 mm area. The sputtering rate was calibrated using a 100 nm thick Ta2O5 oxide layer as a reference sample. The estimated etching rate was 4 nm/min. Intensities were determined from the peak-to-peak heights of the differentiated spectra.

Auger line energies used for the concentration determination were: O 510 eV, Cr 530 eV, Fe 705 eV and Al 1395 eV. For aluminum, high energy line was used as there is no interference with other lines and also the effect of the chemical environment on the peak shape is smaller. Sensitivity factors for the metallic components were measured from pure Cr, Fe and Al samples. Oxygen sensitivity factor was determined from an Al2O3 layer [29]. Calibration samples were first sputter cleaned and then measured using the same experimental parameters. For aluminum oxide, the sputtered surface concentration was supposed to be the same as in bulk [56]. 4. Results 4.1. Theoretical surface profiles and chemical potentials of alloys The calculated surface profiles (Table 1) can be compared with the LEED measurements of the top-layer concentration of Fe1xAlx(1 0 0) [57–59]. In these measurements the surfaces were annealed not beyond about 1200 K before quenching. For the Fe– 5Al alloy the surface Al-concentration according to the abovementioned LEED measurements is about 57 at%. Our surface profile at 1200 K is in good agreement with these experimental results. The difference of the chemical potentials of Fe and Cr (lFe  lCr) as well as that of Fe and Al (lFe  lAl) were calculated using Eq. (1) and they are shown in Figs. 2 and 3. The bulk data was calculated with a dense concentration mesh (xCr, xAl) but since the computational cost increases considerably for surface calculations the surface mesh is sparser. Due to the smaller number of surface calculations (16 calculations for (lFe  lCr)surface and 15 calculations for (lFe  lAl)surface) the difference of the surface chemical potentials was first fitted to a second order two-dimensional polynomial. This polynomial with fixed Al-concentrations (1, 5, and 10 at% Al) is shown as surface data in Figs. 2 and 3. Since at this point our model system used for Fe–Cr–Al alloys is a homogeneous slab the surface and bulk concentrations in Figs. 2 and 3 are the same. Comparison of Figs. 2 and 3 shows that for lFe  lCr the bulk and surface curves intersect, as in the Fe–Cr case [55], whereas for lFe  l Al the surface curves are far above the bulk curves. Especially the (lFe  lCr)bulk curves have a steep slope within the range

45 bulk 1% Al 5% Al 10% Al surface

40 35

μFe- μCr [mRy]

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30 25 20 15 10 5 0

0

5

10

15

20

25

Cr (at. %)

Table 2 Concentrations of the alloys determined by EDX. Alloy

Fe (at%)

Al (at%)

Cr (at%)

Fe–13Al Fe–18Al Fe–23Al Fe–10Cr–10Al

87.4 82.1 77.4 79.9

12.6 17.9 22.6 9.8

– – – 10.3

Fig. 2. Bulk and surface chemical potential differences (lFe  lCr) of Fe–Cr–Al, (Al percentages for surface potentials from top to bottom: 1 at%, 5 at% and 10 at%). Surface data is taken from two-dimensional polynomial fit and the calculated bulk values (shown by symbols) are connected by spline curves. The surface has the same composition as the bulk. The zero level of lFe  lCr corresponds to 443.865 Ry in absolute scale. Atomic chemical potentials are presented, hence atomic Rydberg units (Ry = 2.17987  1018 J) are used for convenience.

E. Airiskallio et al. / Corrosion Science 52 (2010) 3394–3404

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70

μFe- μAl [mRy]

60 bulk 1% Al 5% Al 10% Al surface

50 40 30 20 10 0

5

10

15

20

25

Cr (at. %) Fig. 3. Bulk and surface chemical potential differences (lFe  lAl) of Fe–Cr–Al, (Al percentages for surface potentials from top to bottom: 1 at%, 5 at% and 10 at%). Surface data is taken from two-dimensional polynomial fit and the calculated bulk values (shown by symbols) are connected by spline curves. The surface has the same composition as the bulk. The zero level of lFe  lAl corresponds to 2059.943 Ry in absolute scale (Ry = 2.17987  1018 J).

Fig. 4. AES surface depth profiles of Fe–18Al oxidized at 1000 °C and 1 atm for 5 min.

of 0–10 at% Cr. This can be related to the transition from the Cr-miscible to the Cr-immiscible region in the Fe–Cr bulk phase diagram [55]. It is also interesting to note that adding Al to the alloy shifts the chemical potential difference curves down in Figs. 2 and 3. This can be related to the effect of increasing the volume of the alloy with increasing Al-content (Fig. 1). Similar effect has been obtained also in Fe–Cr–V alloys. [60] The trends of the bulk chemical potentials in Figs. 2 and 3 are also consistent with the experimentally observed Al partitioning in Fe–Cr–Al alloys which shows the depletion of Al from the Cr-rich phases [61–63]. 4.2. Experimental Auger profiles of oxidized surfaces The prepared ingots of Fe–13Al, Fe–18Al, Fe–23Al, and Fe– 10Cr–10Al were exposed to oxidizing atmosphere and the resulting oxide scales were investigated by repeated Ar sputtering and Auger measurements. Our Auger measurements of the surface depth profiles of oxidized Fe–13Al, Fe–18Al, Fe–23Al, and Fe– 10Cr–10Al show that Fe–13Al does not form the protective oxide scale. Instead, the surface was almost entirely covered with a rough and thick Fe-oxide layer without any traces of Al. In striking contrast to this the other three alloys form thin protective oxide layers consisting mainly of Al-oxides. Our experimental AES-results are shown in Figs. 4–6. As shown in the figures the amount of Al (Fe) increases (decreases) in the oxide scale in the sequence of Fe–18Al, Fe–23Al, and Fe–10Cr–10Al. This is also the sequence of the improved oxidation resistance, since the Al-oxide gives better protection against oxidation at high temperature compared with the Cr- and Feoxides.

Fig. 5. AES surface depth profiles of Fe–23Al oxidized at 1000 °C and 1 atm for 5 min.

5. Discussion 5.1. Scaling losses To take a full advantage of our electronic structure calculations in the quest for veiled atomic phenomena related to the oxidation of Fe–Cr–Al we first survey the published data focusing on the formation of oxide scales on the Fe–Cr–Al surfaces. The rates of the scaling loss (weight lost by oxidation) of Fe–Cr–Al alloys at 1200 °C, reviewed in Ref. [4], are first converted into a two-dimensional plot (Fig. 7). This representation shows in parallel how both

Fig. 6. AES surface depth profiles of Fe–10Cr–10Al oxidized at 1000 °C and 1 atm for 5 min.

alloying components (Cr and Al) affect the formation of the oxide films. One should note that the experimental data that we have used in our fitting procedure, is limited to the region where the

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E. Airiskallio et al. / Corrosion Science 52 (2010) 3394–3404

contours at levels: 0.5, 1, 2, 5, 10, 50 40

40 threshold

200

35

180 35

30

160

30

Cr (at. %)

140 120 100

Cr (at. %)

25

80

25 20 15 10

60 20

5

40

0

20 15

5

10

15

20

25

Al (at. %) Fig. 8. Fe–Cr–Al alloys rated as good (filled symbols) or poor (open symbols) with respect to the high temperature (P1000 °C) oxidation resistance. The data is from Ref. [4] (triangles pointing downward) and from Ref. [20] (triangles pointing upward). The dividing threshold curve between the poor and good oxidation resistance domains is the 5 g/(m2h) scaling loss curve adapted from Fig. 7.

10

5

0

0

0

0

5

10

15

20

Al (at. %) Fig. 7. Experimental scaling losses (g/m2h) of Fe–Cr–Al alloys at 1200 °C after 240 h oxidation. The colour coding is from 0 to 200 g/m2h. To illustrate the groups of alloys having similar oxidation rates five equiscaling contours are shown in the plot. The plot is obtained by fitting to the data shown in Ref. [4].

oxidation rate is 10 g/(m2h) or less. Therefore, the data shown in the lower left corner of Fig. 7 is a result of extrapolation. However, one should note that at high temperatures (1000 °C) J 10 at% Al (p. 84 in [4]) or J 14 at% Cr (p. 126 in [2]) are reported to give sufficient protection against corrosion in Fe–Al and Fe–Cr, respectively. It is interesting to note that for these two threshold values our fitting function in Fig. 7 gives scaling losses approximately of the same order in magnitudes. This clearly suggests that the fitting procedure works well and the obtained function can also be used within extrapolated regions. As Fig. 7 shows, at 1200 °C in most cases Al is about 1.5–2 times as effective oxidation retardant as Cr if scaling losses in Fe–Al and Fe–Cr are compared. Nevertheless, considering Fe–Cr–Al one can define a line in Fig. 7 where Al and Cr are equally effective, i.e. the reduction in scaling loss is the same for the same amount of added solute Al or Cr. This dividing line is obtained by joining the points where the normal of an equiscaling contour is at 45 deg to the horizontal direction. This happens approximately for the concentrations (cCr, cAl) satisfying

cCr  0:7cAl  2;

ð3Þ

where at% are assumed for the concentrations. For alloys above this line in Fig. 7 adding Al while keeping Cr content fixed reduces scaling losses more than doing vice versa (adding Cr with Al-content fixed). Below this line the effectivity of Al and Cr is reversed. 5.2. Protective oxide scales The reported high temperature (P1000 °C) data concerning the rating of the oxidation resistance of Fe–Cr–Al as good or poor [4] is shown in Fig. 8. Although the results scatter due to dissimilar experimental conditions and rating criteria used in different investigations one can divide the diagram approximately into two parts, showing good or poor oxidation resistance depending on the con-

centrations of the alloys. For instance, if the 5 g/(m2h) contour from Fig. 7 is replotted in Fig. 8 it approximately splits the alloys in two groups possessing good or poor high temperature oxidation resistance. According to the review article of Thomaszewicz and Wallwork [4] the Al-concentration in bulk Fe–Al alloy should be at least 10–15 at% in order to the oxidation resistance to be effective at high temperatures. They also report that adding 10 at% Cr into Fe–Al allows one to reduce the Al-content down to 3 at% without weakening the corrosion resistance of the alloy. The abovementioned 5 g/(m2h) contour starts from the Fe15Al and goes to about 4 at% Al level when there is 10 at% Cr in Fe–Cr–Al. Therefore, the 5 g/(m2h) oxidation rate curve seems to give an approximate lower limit for the combined Al and Cr concentrations in high-temperature corrosion resistant Fe–Cr–Al alloys. Correspondingly, we suggest that for Fe–Cr–Al alloys to have good high temperature oxidation resistance their concentrations should be chosen according to the following relation, corresponding to the dividing threshold curve in Fig. 8,

cAl > 4 þ

150 ; 8 þ cCr

ð4Þ

where concentrations are given in at%. Additionally one should note that the oxidation resistance of Fe–Al and Fe–Cr–Al alloys seems to correlate with the Al-concentration in the surface layer. According to Fig. 8 the Fe–10Cr–5Al and Fe–15Al alloys have approximately similar oxidation resistances. Table 1 shows that for the Fe–10Cr–5Al the calculated Al concentration in the surface at 1200 K is 72 at% and the LEED measurements [59] give about 74 at% Al for the top layer of Fe–15Al, annealed not beyond 1200 K and quenched. 5.3. Oxide type and chemical potentials We are now at the point when we can compare the calculated chemical potentials (Figs. 2 and 3) with experiments. For that purpose we present the difference of the chemical potentials (lFe  lAl)bulk  (lFe  l Al)surf and (lFe  lCr)bulk  (lFe  lCr)surf in two-dimensional plots (Figs. 9 and 10). In these figures we also show the results of the experimental characterization of the main component of the surface oxides of Fe–Cr–Al alloys above 1000 °C [4] and at 800 °C [64].

E. Airiskallio et al. / Corrosion Science 52 (2010) 3394–3404 -70 -68 -66 -64 -62 -60 -58 -56 -54 -52 CrO AlO (Cr,Al)O FeO

15 10

Cr (at. %)

20

5 0 0

5

10

15

20

25

Al (at. %) Fig. 9. Calculated chemical potential difference (lFe  lAl)bulk  (lFe  lAl)surf (in mRy) as a function of Cr- and Al-concentrations. The results of the experimental characterization of the main component of the oxide scale grown on Fe–Cr–Al above 1000 °C [4] are shown by symbols. Gray triangle upward: Cr2O3, blue triangle downward: Al2O3, green square: mixed Cr2O3 and Al2O3, black circle: Fe2O3. (For interpretation of the references to colour in this figure legend, the reader is referred to the web version of this article.)

II

20

IV

15 10

Cr (at. %)

15 10 9 8 7 6 5 4 3 2 1 0 −1 −2 −3 −4 −5 CrO AlO (Cr,Al)O FeO

5

I 0

III 5

0 10

15

20

25

Al (at. %)

Fig. 10. Calculated chemical potential difference (lFe  lCr)bulk  (lFe  lCr)surf (in mRy) as a function of Cr- and Al-concentrations. The results of the experimental characterization of the main component of the oxide scale grown on Fe–Cr–Al above 1000 °C [4] are shown by symbols. Gray triangle upward: Cr2O3, blue triangle downward: Al2O3, green square: mixed Cr2O3 and Al2O3, black circle: Fe2O3. The major types of scaling behaviours fall into four regions I–IV, separated by solid black lines [64]. The latter data refer to 800 °C and the major components of the oxide scales are, I: Fe2O3, II: Cr2O3, III: Al2O3 with Fe2O3 nodules, IV: Al2O3. (For interpretation of the references to colour in this figure legend, the reader is referred to the web version of this article.)

Somewhat surprisingly, it is observed that although Al-oxide is the protective component in the oxide scale, there is no correlation between the calculated Fe–Al chemical potentials and the measured composition of the oxide scale (Fig. 9). The reason for this result could be that the chemical potential difference in Fig. 9 is within the range from 70 to 50 mRy (1.5  1019 to 1.1  1019 J). This means that the driving force for the Al diffusion between the bulk and the surface regions due to chemical potentials is always directed from the bulk to the surface and the relative variations in the Al driving force are small within the concentration region considered. On the contrary, as Fig. 10 shows, the Fe–Cr chemical potential changes sign within the investigated region implying significant changes to the chemical potential induced driving force of the Cr diffusion. The Fe–O dominated scales appear in the lower left corner of the plot (Fig. 10), where the driving force for Cr is from the surface to the bulk and there is practically no Cr at the equilibrium surfaces of these alloys. The experimental borderline of the changing of the Fe– O dominance to Cr–O or Al–O dominance lies approximately at the line where the chemical potential difference, i.e. the driving force of Cr diffusion changes its sign. The Cr–O dominance is changed to

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Al–O dominance at about 5 at% Al, close to the saddle point of the Fe–Cr chemical potential difference plot. The reason for the suggested importance of Cr on the oxidation of Fe–Cr–Al in general could be the fast formation of Cr2O3 on the surface, a phenomenon that would give more time for a-Al2O3 to form up, provided there is enough Cr at the surface at the initial stage of the oxidation. This is in line with the grazing angle X-ray diffraction (XRD) measurements of Fe–Cr–Al. In contrast to the commonly held view that corundum-type oxides only form above 900 °C [25], there is a clear evidence in the XRD results of Fe–Cr–Al that a-Al2O3 forms even at 700 °C [29]. This was proposed to be due to the fact that the corundum nucleation is greatly facilitated on a surface consisting of the isostructural Cr2O3. 5.4. Tuning the chemical potentials by Cr substitution Next we discuss how the substitution of Al and/or Fe by Cr in bulk Fe–Cr–Al affects the chemical potentials of the components and how this is expected to be reflected in the state of the Fe– Cr–Al surfaces. The Cr substitution in bulk has direct effects on both bulk and surface quantities. Starting with bulk properties we note that lFe  lAl and lFe  lCr both decrease steeply with increasing Cr content, within the region from 0 up to 10–15 at% Cr (Figs. 2 and 3). In contrast to the situation in bulk the chemical potential of the surfaces changes rather slowly within the same concentration region. This implies increasing chemical potential induced driving force for Al atoms to diffuse from the bulk to the surface, with increasing Cr content in bulk. Fig. 2 also shows that the driving force of the Cr diffusion reverses from the surface to bulk direction to the bulk to surface direction at about 10 at% Cr in bulk Fe–Cr–Al. Summarising from above, the Cr addition to Fe–Al has a double-edged effect: the equilibrium surfaces will contain more Al than they would have otherwise and within a certain concentration region also Cr appears in the surface. This effect can be clearly seen in Table 1 where the 10 at% increase in the bulk Cr content increases the Al-concentration in the surface layer by about 20 at%. Table 1 also shows that traces of Cr are expected to be found in the surface at high temperature when the Cr concentration in bulk reaches the 10 at% level. Increasing the Cr content in Fe–Cr–Al up to 15 at% shows that the Cr-driven surface enrichment with Al is saturated. This is due to the flattening of the lFe  lAl (Fig. 3) when the Cr content in bulk is increased beyond 10 at%. Because the changes in the relative chemical potentials in the bulk seem to be the key quantities in determining the surface concentrations of Fe–Cr–Al alloys at equilibrium conditions it is interesting to analyse the mechanisms through which the substi tutional Cr affects the chemical potentials in more detail. For that purpose we calculated for Fe–Cr–Al the density of states (DOS) at Fe, Cr, and Al sites as a function of Cr content. Comparing the site projected DOSs, relative to the common Fermi energy, reveals that the peaks in the occupied part of the Fe DOS shift down with the increasing Cr concentration. Therefore, the projected band energy in Fe sites is reduced and since Fe, Cr, and Al sites have the same Fermi energy this lowering of the Fe bands is consequently transferred to (lFe  lAl)bulk and (lFe  lCr)bulk which decrease similarly with increasing Cr content (Figs. 2 and 3). To some extent it is also possible to relate the chemical potentials to atomic pair energies. To compare our results with these quantities we refer to investigations on FeAl with Cr impurities. A study on the effect of Cr on the antiphase boundaries of FeAl alloy [65] reports the interatomic energies. The interatomic energies of Al–Al, Cr–Al, and Fe–Al pairs are of the same order of magnitude, approximately from 0.4 to 0.6 eV. The Fe–Fe interaction energy is about from 1.2 to 1.3 eV and the Cr–Fe interaction energy is about 1.6 eV. This suggests that adding Cr into Fe–Al increases

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the average bonding of Fe to the alloy matrix which is in line with our result of lowering the band energy of Fe. The site resolved mixing energies in FeAl show slight preference for Cr to substitute Al, i.e. Cr prefers Fe neighbors [66] suggesting stronger bonding in Fe– Cr than in Fe–Al pairs. Comparing the Cr and Ni additions in FeAl, it has been found that contrary to Cr, Ni prefers to substitute atoms in the Fe sublattice [67] which is in line with the fact that Ni tends to destroy the oxidation resistance of Fe–Al [4], quite an opposite effect compared to that of Cr. 5.5. The role of Cr in the oxidation of Fe–Cr–Al The Cr induced changes in the state of the Fe–Cr–Al surfaces suggest important consequences on the formation of the protective oxide scale on the surfaces of Fe–Cr–Al. The initial oxidation depends of course on the state of the alloy surface prior to oxidation. The Cr addition has both direct and indirect effects. The predicted increased Al-content in the surface suggests direct improvement in the formation of protective Al-oxide scales. On the other hand, Cr itself in the surface may induce transient oxidation containing Cr-oxide patches which act as oxidation retardants and nucleation centers for a-Al2O3 [29]. Since the formation rate of Cr-oxide is higher than that of Al-oxide [68] and Cr2O3 has the same crystal structure as a-Al2O3 the above facts are suggested to be beneficial to the growth of the protective a-Al2O3 scale on the Fe–Cr–Al surfaces. 5.6. Oxidation experiments Now we are ready to analyse our experimental observations in the light of our theoretical results and the published data. We start with the characterization of the oxide scales of Fe–Al and Fe–Cr–Al alloys. Analysing Fig. 4 we can distinguish four different zones within the Auger profile of Fe–18Al. Although the division of the Auger profiles into different zones is to some extent arbitrary it is, however, a very useful concept in investigating the oxidation mechanism in general and in comparing the oxidation of different alloys. Here we omit the approximately 5 nm thick topmost surface layer which we expect to have been affected by atmospheric contaminants. Based on the observed atomic fractions of Fe, Al and O we define the different zones in sequence from the surface to the base alloy as: 1. Surface zone. Within this about 75 nm thick zone the oxygen content is practically constant while the Fe content decreases from 5 at% to zero and the Al-content increases from 35 to 40 at% correspondingly. 2. Saturation zone. Here the Al-content reaches its maximum and there is practically no Fe in this zone. The oxygen content is at the same level as in the previous zone. The thickness of the saturation zone is about 40 nm. 3. Inversion zone. In the next, about 60 nm wide section the Fe content increases and the Al-content drops down, both approaching their bulk values. The right hand side end of the inversion zone is defined here to be at the inflexion point of the Al-concentration curve. 4. Deflection zone. This about 50 nm thick zone is found below the inversion zone. Within this zone the atomic concentrations are still deflected from their bulk values to some extent. One should note that different sputtering yields of the chemical elements may lead to somewhat biased values for the atomic fractions in the profiles. Especially the result for Al fraction might be slightly underestimated. However, in our investigations we used such Auger lines which are not very surface sensitive [69]. Additional uncertainty to the atomic fractions in the alloy

side is expected to be due to the different sensitivity factors in the oxide and metal matrices as already discussed in the experimental procedure section. The four zones described above are clearly discernible also in the Auger profiles of Fe–23Al (Fig. 5). In this case the first two zones are almost identical to those found in the Fe–18Al case. A slight difference is found at the top of the surface zone: it suggests that the Fe content decreases and correspondingly the Al-content increases when the Al-concentration in the base alloy is raised from 18 to 23 at%. The zones of Fe–23Al are slightly thinner than the zones of Fe–18Al. Compared to the Fe–23Al case, the zone structure of Fe–10Cr– 10Al (Fig. 6) differs more from the zone structure of Fe–18Al. From top to bottom the thickness of the first three zones in Fe–10Cr– 10Al are about 85, 50, and 80 nm. The thickness of the deflection zone cannot be determined reliably in this case because our measurement does not extend to a depth which would show the Alcontent to have reached the bulk value. However, we estimated this zone to be about 60 nm thick. In addition to the fact that the thicknesses of the zones are somewhat larger in Fe–10Cr–10Al compared to those in Fe–18Al the heights of the concentration profiles within the zones differ from those of Fe–18Al as well. The concentration variations between the surface and saturation zones in Fe–10Cr–10Al are smaller than in Fe–18Al. Another clear difference is that in Fe–10Cr–10Al the Al-concentration does not have a dip in the deflection zone as it has in the Fe–18Al and Fe–23Al cases. This is in line with our earlier prediction of the Cr-induced Al-pump [70]. Furthermore, in Fe–10Cr–10Al there is a Cr contribution in the inversion and deflection zones. No Cr in the surface and saturation zones were detected which is in agreement with the differences between the diffusion constants of the components. The integrated interdiffusion constant of Cr is much lower than that of Fe and Al. The approximate ratios are 1:4:5 for Cr, Fe, and Al in (Fe, Cr)3Al2 [71]. The observed weak maximum in the Cr concentration could be a trace of the initial surface since Cr-rich oxide has been related to the position of the initial surface before oxidation [28]. Comparing the Auger profiles of Fe–10Cr–10Al with those of Fe–18Al and Fe–23Al reveals that if the Fe and Cr profiles are added up the obtained sum profile within the inversion zone is very similar to the Fe profiles of Fe–18Al and Fe–23Al. Also the Al profiles of Fe–10Cr–10Al, Fe–18Al, and Fe–23Al are quite similar. One should note that it is the Al-oxide that is considered to block the atomic diffusion at high temperatures in these alloys. Therefore, substituting 1/9 of Fe by Cr in the Fe–10Al alloy seems to raise the Al-content in the oxide scale to the same level as that found in the Fe– 18Al and Fe–23Al alloys. 5.7. Initial oxidation So far we have not discussed much the important initial stage of the oxidation [2,3] of Fe–Al alloys just at the oxidation resistance borderline, i.e. alloys containing 10–15 at% Al [4]. Therefore, we now try to get better understanding of the high temperature oxidation of Fe–13Al alloys. For that purpose we re-examined our Fe– 13Al sample at the initial stage of oxidation. Instead of 5 min oxidation the sample was oxidized only 1 min at normal pressure and 1000 °C temperature. After oxidation the surface was scanned by AFM and a surface with nodules randomly scattered on an otherwise smooth base was observed (Fig. 11). Our observation is in agreement with earlier results, see e.g. [64,72], which show randomly positioned Fe-oxide and Fe-sulfide nodules on a smooth Al-oxide surface of Fe–Al and Fe–Cr–Al alloys. Analysing the composition of the nodules reveals that they consist almost exclusively

E. Airiskallio et al. / Corrosion Science 52 (2010) 3394–3404

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Fig. 11. AFM image of an Fe–O nodule on the smooth surface oxide scale of Fe–13Al oxidized at 1000 °C and 1 atm for 1 min.

of iron and oxygen. The nodular structure has been observed also earlier for Fe–Al and Fe–Cr–Al alloy surfaces [20,27,73,64]. Between the nodules we found a smooth thin oxide scale containing both Fe- and Al-oxides. A typical Auger profile measured in this smooth surface region between the nodules is shown in Fig. 12. Comparing Fig. 12 with Fig. 4 one realizes that the thin smooth oxide scale in Fe–13Al shows similar features as the protective scale in Fe–18Al. The significant differences between Fe– 13Al and Fe–18Al are found in the surface zone, i.e. in the topmost part of the oxide scale. In Fe–13Al the surface zone contains more Fe than Al and also the oxygen content, compared to the metal atoms, is higher than the ratio 3/2 which was found in Fe–18Al. Actually, the total iron content within the surface zone of Fe– 13Al is several magnitudes higher than in the Fe–18Al case. This suggests that at the initial stage the Al-oxide scale in Fe–13Al leaks and there exists a substantial Fe flux from the base alloy to the surface through the Al-oxide scale during the first moments of the oxidation. Our investigations suggest that the surface of Fe–13Al is partially covered by the protective Al-oxide scale. This Al-oxide scale shows two blocking fronts (Fig. 12). The blocking front for the Fe diffusion from the bulk alloy to the surface is located about 125 nm below the surface and the blocking front for the O diffusion from the surface to the bulk alloy is found about 50 nm below the surface. However, the protective Al-oxide scale in Fe–13Al is only a transient feature. We already know from our earlier oxidation experiments that after 5 min of oxidation most of the surface is covered with a non-protective thick Fe-oxide scale. The protective Al-oxide scale of Fe–13Al can be destroyed principally in two ways. The destructive Fe-oxidation may start from the Fe-oxide rich nodules and proceed sideways along the surface into the protective

Fig. 12. AES surface depth profiles of the smooth oxide scale of Fe–13Al oxidized at 1000 °C and 1 atm for 1 min.

Fig. 13. Optical microscope image of Fe–O nodules at the surface of Fe–13Al oxidized at 1000 °C and 1 atm for 1 min.

oxide scale, eventually breaking it. There can be also hidden defects in the protective scale through which Fe atoms can gradually diffuse from the bulk alloy to the surface and O atoms can diffuse from the surface to the bulk alloy. This kind of leak process would, however, trigger new Fe-rich nodules on the smooth oxide scale. To obtain information on the large scale properties of the nodular structure of the oxidized Fe–13Al surface optical microscope was used. As Fig. 13 shows the Fe-oxide nodules are randomly scattered throughout the otherwise smooth surface. Since after 1 min of oxidation all the nodules have the same symmetrical circular shape and they all are of the same size, we conclude that they all have emerged at the very first moments of the oxidation from point defects scattered randomly on the initial alloy surface. During the oxidation the nodules grow and after 5 min they practically cover the whole surface. To improve the oxidation resistance of low-Al Fe-alloys it is important to investigate what are those point defects that nucleate the nodular growth of Fe-oxide and what are the possibilities to inactivate them. The main features of our Auger depth profiles agree with those of other similar measurements on Fe–Al and Fe–Cr–Al alloys [73– 78]. Summarizing from above, the following conception of the oxidation process of the investigated alloys is proposed. (i) For Fe–Al alloys under oxygen flow at 1000 °C and 1 atm the threshold value of the Al-content for the protective Al-oxide scale to form up and hold for a longer time is between 13 and 18 at% in the base alloy. (ii) In the main part of the oxide scale the oxygen and metal atom concentrations are 60 and 40 at%, respectively, corresponding to the atomic fractions in the stoichiometric oxides Al2O3 and Fe2O3. This suggests that in the surface and saturation zones practically all oxygen and metal atoms are combined into oxides. (iii) The building up of the Fe-free saturation zone can be understood in the following way. At the early stages of oxidation there presumably exist both Al- and Fe-oxides in the saturation zone region. However, the oxygen flow from the gas phase is decreased step by step as the oxide layer grows. If there is no free oxygen available and the Al flux from the base alloy brings more Al to this region the higher oxygen affinity of Al [79], compared to that of Fe, leads to the process 2Al + Fe2O3 ? 2Fe + Al2O3 which combined with the subsequent Fe escape due to diffusion saturates this zone with Al-oxide.

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(iv) When the Al-concentration of an Fe–Al alloy is close to the threshold value of the onset of the high temperature oxidation resistance both protective and non-protective surface oxide scales coexist at the surface. The non-protective oxide scale grows by the rapid generation of the Fe–O nodules during the first moments of the oxidation. These nodules are nucleated by the surface point defects. While the oxidation continues the nodules grow at the expense of the protective scale. (v) Replacing Fe by Cr in the base alloy pushes more Al to the surface region leading to the formation of the protective Al-oxide scale with considerably reduced Al-content in the base alloy. In the oxide scale of Fe–10Cr–10Al the concentration of Al is even higher than in the oxide scales of Fe–18Al and Fe–23Al alloys.

Ds ΔEs

ΔEb Energy

3402

Db

0

1

2

3

n (Distance form the surface) 5.8. Atomic scale mechanisms To shed some light on the underlying mechanisms of the surface oxidation of Fe–Al alloys we reconsider our experimental data in the light of our computational results. The computational results refer to the (0 0 1) surface. However, the oxidation experiments for FeAl above 500 °C lead to the formation of well-ordered Al2O3 films on all low-index surfaces (0 0 1), (0 1 1), and (1 1 1) [34]. Moreover Al segregation is observed for all low-index surfaces [69]. We therefore expect that the theoretical results for the (0 0 1) surface can be justifiably used in analysing our experimental data. From the theoretical point of view we can describe two features that affect the oxidation resistance of Fe–Al alloys at high temperatures. The first one is a straightforward mechanism to improve the oxidation resistance simply by increasing the Al-concentration in the base alloy. This accelerates, in a natural way, the formation of the protective Al-oxide scale which then reduces both internal oxidation as well as the growth of the outer oxide scale. However, there exists another, a more complex, mechanism that is also beneficial to the oxidation resistance of Fe–Al alloys. This can be explained by considering the calculated chemical potentials shown in Fig. 3. As seen in Fig. 3 there exists a strong chemical potential induced driving force for Al to diffuse from bulk to the surface which is in line with the surface energies of the components [80,81] and the results of the surface segregation investigations [82]. It is also useful to relate the data shown in Fig. 3 to the diffusion coefficients in Fe–Cr–Al. Considering either direct or concerted Al–Fe pair exchange process [83] the corresponding diffusion coefficient for the Al moving towards the surface (Ds) or Al moving towards the bulk (Db) are approximated as

Ds ¼ D0 eðDEs =ðkB TÞÞ ;

ð5Þ

Db ¼ D0 eðDEb =ðkB TÞÞ ;

ð6Þ

where T is the temperature and kB is the Boltzmann’s constant, the activation energies are

DEs ¼ Esp  EAlðnþ1Þ;FeðnÞ ; sp

AlðnÞ;Feðnþ1Þ

DEb ¼ E  E

:

ð7Þ ð8Þ

Here D0 and Esp are the pre-exponential factor and the energy of the system when the Al–Fe pair is in its saddle-point configuration along its migration path, both assumed to be independent of the direction of the exchange process. EAl(n + 1),Fe(n) and EAl(n),Fe(n+1) are the energies of the system in the initial states corresponding to Al–Fe pair having the Al end in the bulk side or in the surface side, respectively. The index n refers to the atomic layer, n = 0 refers to the surface layer and deeper layers towards the bulk are indexed

Fig. 14. Schematic figure of the energetics of the exchange processes of Al–Fe pairs. The red wavy curve represents the energy of the system at different stages of the pair exchange of an Fe–Al pair. The configuration Al at n = 1 and Fe at n = 2 is shown by the green sphere on the left and the configuration Fe at n = 1 and Al at n = 2 is shown by the green sphere on the right. On average Al atoms in Fe–Cr matrix will diffuse from right to left, i.e. from the high Cr region (bulk side, high potential) to the low Cr region (surface side, low potential). (For interpretation of the references to colour in this figure legend, the reader is referred to the web version of this article.)

with successively increasing n. From Eqs. (5)–(8) one obtains the ratio of the diffusion coefficients n

n

Ds eðlFe lAl Þ=ðkB TÞ ¼ nþ1 nþ1 : Db eðlFe lAl Þ=ðkB TÞ

ð9Þ

According to our calculations (Table 1) the Cr concentration at the surface is lower than that in the bulk. Therefore, it is expected that nþ1 in the region between the bulk and the surface lnþ1 is less Fe  lAl than lnFe  lnAl (in line with Fig. 3) leading to Ds/Db > 1, i.e. the above model predicts the driving force of the Al diffusion to be towards the surface. The energetics of the diffusion is illustrated in Fig. 14. Due to the anisotropy of the Al diffusion, the Al-concentration in the surface region is expected to exceed that in the bulk. This enrichment of the surface with Al continues as long as the minimum of the free energy has been achieved or the surface has become saturated with Al. When this stage is reached there is no net driving force for Al diffusion to any direction. In Table 1 we show the calculated equilibrium surface concentration of Fe–5Al alloy as a function of temperature. We see that at 0 K the surface in the equilibrium state contains 87 at% Al. Using the data in Fig. 3 we estimate that in the case of Fe–13Al, Fe– 18Al, and Fe–23Al the equilibrium surface at 0 K would contain about 94, 98 and 100 at% Al, respectively. However, at high temperatures the entropy contribution in the free energy becomes significant and consequently at about 1000 °C the corresponding equilibrium Al surface concentrations for Fe–13Al, Fe–18Al, and Fe–23Al are 67, 72 and 77, respectively. It is interesting to note that the Al surface concentrations at 1000 °C of Fe–13Al and Fe5Cr5Al are approximately the same and also according to Fig. 7 the experimental scaling losses of these alloys are of the same order. Another similar pair of alloys is Fe–18Al and Fe10Cr5Al suggesting for a connection between the Al-concentration at the surface layer and the scaling loss in oxidation atmosphere. What is even more interesting here is the role of Cr in enhancing the oxidation resistance of Fe–Al alloys. The effect of Cr on the driving force of Al diffusion is clearly seen in Fig. 3. The chemical potential difference (lFe  lAl)bulk decreases substantially when Cr content in the base alloy is increased from 0 to 10 at%. This decrease means that in the bulk the chemical potential of Al is raised com-

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pared to that of Fe. Because the same Cr addition does not change the chemical potential difference at the surface appreciably, the driving force of Al from bulk to the surface is therefore considerably enhanced due to the Cr addition as verified computationally and shown in Table 1. The Cr-enhanced oxidation resistance discussed above is effective mainly at the initial stage of oxidation or during the healing of a broken oxide scale, i.e. in cases where an open metal surface is exposed to oxidizing environment. However, this is not the whole story of the beneficial effect of Cr in Fe–Al alloys. Cr improves the oxidation resistance of Fe–Al alloys also during the steady state of the oxide scale growth. Comparison of Figs. 5 and 6 shows that an Fe– 10Cr–10Al alloy has a better supply of Al in the oxide scale than an Fe–23Al alloy even though the Al-content in the latter is more than twice of that of the first one. In the oxide scale of Fe–10Cr–10Al the Al-content is higher and correspondingly the Fe content lower than in the oxide scale of the Fe–23Al alloy. This can be understood on the basis of the calculated chemical potentials shown in Fig. 3. During the initial oxidation also Cr oxidizes leading to a Cr-depleted zone [84,15] under the oxide scale. Due to the slow mobility of Cr [71] this Cr-depleted zone is stable. If the Al-concentration in bulk is below the critical level, then due to the high oxygen affinity of Al [79] the exposure of the surface to oxygen results in a formation of an Al-depleted zone under the oxide scale, which could eventually lead to the breakdown of the protective oxide scale [4]. The important issue within the maintaining of the protective oxide scale is the recovery of the Al level in the depleted zone. At this point Cr plays an important role. As seen in Fig. 3 the Cr depletion in Fe–Cr–Al raises the chemical potential of Fe relative to that of Al and consequently helps in restoring the Al to Fe ratio in the depleted zone to the level that is needed to maintain the protective oxide scale. Using the chemical potentials shown in Fig. 3 we can give a theoretical estimate, based on our calculations, on the effectivity of Al and Cr on the improvement of the high temperature oxidation resistance of Fe–Cr–Al alloys. To get the same improvement in the oxidation resistance the Al addition should be about twice of the amount of the Cr addition. This result is in line with our Auger measurements since comparing Figs. 4–6 one observes that the oxide scale of Fe–10Cr–10Al contains less Fe than that of Fe– 23Al. Therefore, adding 10 at% Cr is effectively more than adding 13 at% Al. Also the experimental results of the scaling losses (Fig. 7) show that if one considers, for instance, Fe–10Al which has the scaling loss of about 50 g/(m2h), there are two extreme ways to improve its oxidation resistance. One can either substitute Fe by Cr or Al. If the Cr concentration is increased from 0 to 5 at% we arrive at the Fe–5Cr–10Al alloy which has the scaling loss of about 2 g/(m2h). In order to achieve the same scaling loss level by Al addition we have to add Al twice of the Cr amount, which leads to the Fe20Al alloy having approximately the 2 g/(m2h) scaling loss. As Fig. 2 shows, adding Cr into Fe–Al alloys affects also lFe  lCr. As above, also in this case the drop of lFe  lCr in the bulk is much larger than at the surface. Therefore, considering FexCr10Al alloys one observes that for small x the driving force for Cr (relative to that of Fe) is from the surface to the bulk. Increasing the Cr concentration this tendency gradually weakens and at about 10 at% Cr in the base alloy there is no relative driving force anymore and Cr and Fe have equal probabilities to occupy bulk and surface sites. This phenomenon of the increasing of the probability of Cr, relative to that of Fe, to occupy surface sites has also a beneficial effect on the formation of the protective oxide scale on the surface.

6. Conclusions Ab initio electronic structure calculations have been performed for Fe1xyCrxAly alloys, 0 6 x 6 0.25, 0 6 y 6 0.25. The measured

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AES surface depth profiles are compared with the calculated atomic chemical potentials and surface concentration profiles. To conclude we list the main results found in our investigation. 1. Increasing the Al-content of Fe–Al from 13 to 18 at% makes the alloy resistant to oxidation at 1000 °C. Further increase of the Al content to 23 at% slightly improves the protecting Al-oxide scale on the surface. At the borderline of the onset of the oxidation resistance in Fe–Al alloys the nodular growth of Fe-oxide is nucleated at randomly scattered points along the surface. During the oxidation these nodules grow in size but not in number. Further investigation of the nucleation and the early kinetics of these nodules is expected to be useful in improving the quality of Fe–Al surfaces. 2. Substituting Fe by Cr in Fe–Al alloys clearly increases the driving force of Al to diffuse from the bulk to the surface. Because the increase of the Al driving force is mainly due to bulk properties, this increase is not very sensitive to the type of the surface, whether it is metal or oxide, or the type of the oxide. A driving force for the Al diffusion based on the Cr concentration gradient in Fe–Cr–Al is predicted. 3. Comparing the measured chemical composition of the oxide scales with the calculated chemical potentials of the atoms in the bulk and at the surface of the alloy a correlation between the composition of the oxide and the energetic balance between Fe and Cr in the bulk and at the surface (namely (lFe  lCr)bulk  (lFe  lCr)surf) can be seen. In the case of Fe and Al a similar kind of correlation is not found. 4. With increasing Cr content the Fe bands shift down, relative to the Cr and Al bands, which can be related to the calculated changes in the chemical potentials in Fe–Cr–Al. Increasing the Cr concentration in the bulk results in two effects on the surface of the Fe–Al alloys: the Al-content at the surface increases and Cr is also found at the surface. 5. Corrosion resistant surfaces of Fe–Cr–Al are quite robust because their formation and self-healing properties are to a large extent determined by the bulk part of the alloy. Acknowledgements The computer resources of the Finnish IT Center for Science (CSC) and Mgrid project are acknowledged. Financial support from the Academy of Finland (Grant No. 116317), Outokumpu Foundation (E.A.), National Graduate School in Materials Physics (E.N.), the Swedish Research Council, the Swedish Iron Office (B.J., L.V.), and the Carl Tryggers Foundation (M.P.) are acknowledged. References [1] S.L. Case, K.R. van Horn, in: F.T. Sisco (Ed.), Aluminium in Iron and Steel, Alloys of Iron Research Monograph Series, John Wiley and Sons Inc., New York, 1953. [2] A.S. Khanna, Introduction to High Temperature Oxidation and Corrosion, AMS International, Materials Park, OH, 2002. [3] A.S. Khanna, High temperature oxidation, in: M. Kutz (Ed.), Handbook of Environmental Degradation of Materials, William Andrew Publishing, Norwich, NY, 2005, pp. 105–152. [4] P. Tomaszewicz, G.R. Wallwork, Rev. High Temp. Mater. 4 (1978) 75–105. [5] B.A. Gordon, W. Worrell, V. Nagarajan, Oxid. Met. 13 (1979) 13–23. [6] G.B. Abderrazik, G. Moulin, A.M. Huntz, E.W.A. Young, J.H.W. de Wit, Solid State Ionics 22 (1987) 285–294. [7] F.H. Stott, Rep. Prog. Phys. 50 (1987) 861–913. [8] F.H. Stott, F.I. Wei, Oxid. Met. 31 (1989) 369–391. [9] R. Prescott, M.J. Graham, Oxid. Met. 38 (1992) 73–87. [10] J.H. DeVan, P.F. Tortorelli, Corros. Sci. 35 (1993) 1065–1071. [11] I. Gurrappa, S. Weinbruch, D. Naumenko, W.J. Quadakkers, Mater. Corros. 51 (2000) 224–235. [12] C. Schwalm, M. Schütze, Mater. Corros. 51 (2000) 161–172. [13] N. Babu, R. Balasubramaniam, A. Ghosh, Corros. Sci. 43 (2001) 2239–2254. [14] D.B. Lee, G.Y. Kim, J.G. Kim, Mater. Sci. Eng. A339 (2003) 109–114. [15] I.G. Wright, R. Peraldi, B.A. Pint, Mater. Sci. Forum 461–464 (2004) 579–590. [16] J.A. Nychka, D.R. Clarke, Oxid. Met. 63 (2005) 325–352.

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