Oxidation behavior of AISI 316 steel coated by hot dipping in an Al–Si alloy

Surface & Coatings Technology 236 (2013) 188–199

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Oxidation behavior of AISI 316 steel coated by hot dipping in an Al–Si alloy E. Frutos a, P. Adeva a, J.L. González-Carrasco a,b, P. Pérez a,⁎ a b

Departamento de Metalurgia Física, Centro Nacional de Investigaciones Metalúrgicas, CENIM-CSIC, Avda. Gregorio del Amo 8, 28040 Madrid, Spain Centro de Investigación Biomédica en Red de Bioingeniería, Biomateriales y Nanomedicina (CIBER-BBN), Spain

a r t i c l e

i n f o

Article history: Received 1 June 2013 Accepted in revised form 23 September 2013 Available online 4 October 2013 Keywords: Metal coating Stainless steel High temperature oxidation Cavity formation Microstructural evolution Oxidation mechanism

a b s t r a c t AISI 316 LVM stainless steel was coated by hot-dipping into a molten bath containing Al–31 at.% Si. Major advantage of the use of the coating would arise from the development of a protective alumina layer that would permit to raise the working temperature of this stainless steel in oxidizing environments. Uncoated AISI 316LVM alloy develops spinel or chromia layers when it is exposed at high temperatures in air, with the risk of Cr2O3 evaporation as volatile CrO3 at T ≥ 900 °C at high oxygen pressures. The development of an alumina scale would solve this problem. Thus, the high-temperature oxidation behavior of coated AISI 316 alloy was studied between 700 and 900 °C for 150 h and the results compared with the uncoated alloy. The oxidation resistance of coated AISI 316 LVM alloy was substantially improved at 900 °C by the formation of a thin protective alumina layer. Below 900 °C, however, the coating had no any beneficial effect on the oxidation resistance of the stainless steel. Oxidation process was characterized by void development into the coating and beneath the oxide scale as a result of the Kirkendall effect. Nitrides and also oxides were formed at the coating/AISI 316 LVM interface. A mechanism is proposed to explain the resulting oxide scale pattern. © 2013 Elsevier B.V. All rights reserved.

1. Introduction Hot dipping of steels, whereby the alloy is immersed in a bath of molten metal, is a low cost industrial technique successfully used to develop stable metallic coatings without harmful substances such as lead or hexavalent chromium. The bath determines the nature of the metallic coating in such a way that specific properties of the alloy can be improved by controlling its composition. For instance, hot dipping in molten zinc is widely used in industry for galvanizing machine elements and steel profiles for construction or automotive applications. Hot dipping in aluminum baths is effective in improving wear or corrosion resistance in all environments (urban, industrial and marine) [1,2]. Hot dipping in Al–Si alloys is successfully used to enhance magnetic properties of sheets made from ferritic steels by increasing the silicon content in the coating [3,4]. The requirement of increasing the engine efficiency to reduce exhaust gas pollutant emissions according to more recent environmental regulations has pushed automobile industry to raise maximum combustion temperature. Obviously, such increase requires an improvement on the oxidation resistance of exhaust gas systems [5]. Hot dipping in aluminum baths of low alloyed steels and cast irons is used to improve their high temperature oxidation resistance in extremely aggressive atmospheres like those found in exhaust gas systems [6,7]. More recently, hot dipping in Al–10Si alloys has been used to enhance the oxidation resistance of cold rolled steel able to withstand temperatures up to 650– 800 °C without scaling or delamination [8]. Typical applications are ⁎ Corresponding author. E-mail address: [email protected] (P. Pérez). 0257-8972/$ – see front matter © 2013 Elsevier B.V. All rights reserved. http://dx.doi.org/10.1016/j.surfcoat.2013.09.046

exhaust systems, thermal shields, heating equipments, boilers, frying pans, barbecues, baking trays, etc. Hot dip aluminizing of austenitic AISI 316 stainless steels, a material widely used for structural applications at high temperatures because its creep strength is higher than that of low-alloy steels, could constitute a good alternative for these uses [9,10]. Hot dipping of AISI 316 stainless steel would permit to raise the working temperature in oxidizing environments. Protection against oxidation in AISI 316 alloy is conferred by the establishment of spinel or chromia layers whose protectiveness tends to decrease with increasing the temperature. Furthermore, volatilization of chromia layers constitutes an additional limiting factor for using the alloy at temperatures higher than 900–1000°C in oxidizing atmospheres. Hot dipping in Al–Si baths would result in the development of novel (Al,Si)-rich intermetallic coatings tightly adhered to the AISI 316 substrate [11], which should provide enhanced oxidation resistance due to the formation of protective alumina and/or silica layers. This study deals with the evaluation of the oxidation behavior of novel (Al,Si)-rich intermetallic coatings on AISI 316 LVM alloy by hot dipping in a melted bath of hyper-eutectic Al–31Si alloy. The oxidation behavior of the coated material has been evaluated in air in the range of 700–900 °C and compared with the uncoated AISI 316 LVM alloy. 2. Experimental procedure Specimens with a size of 20 cm × 12 cm × 2 mm were cut off from rolled sheets of AISI 316 LVM stainless steel, whose chemical composition (given in atomic percentage) is: Cr 18.73, Ni 13.41, Mo 1.67, Mn 1.64, Si 1.05, C 0.11, Cu 0.06, N 0.24, S 0.002, and Fe in balance. The

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dipping process was performed in a Rhesca hot dipping simulator (ITMA, Spain). The simulator works under a reducing gas atmosphere (N2–5% H2) during preheating and keeps the bath under a protective N2 atmosphere. Hyper-eutectic Al–31 at.% Si (Al–31Si) alloy was used as the melted bath. Before dipping in Al–31Si alloy, the coupons were preheated up to 660 °C with a heating rate of 12 °C/s and held at this temperature for 60 s. Then, coupons were immersed during 120 s in the bath kept at 860°C. After immersion they were cooled to room temperature under a N2 flux. Small pieces were cut from the sheets and then thermally treated at 780 °C for 6 h. Finally, the samples were cooled to room temperature in 8 h. In the following scheme the different steps for the preparation of coated AISI 316 LVM samples are summarized:

189

a



660 C Immersion AI–31Si bath Preheated Coated AISI 316 LVM→ → 60 s 860  C; 120 s AISI 316 LVM AISI 316 LVM

Thickness, microstructural characterization and chemical composition of the coatings were determined using a field emission gun scanning electron microscope (Jeol JSM-6500F) coupled with an energy dispersive spectroscopy (EDS) system. Chemical microanalysis was performed on coated samples after and before the heat treatment. Phase identification was carried out on a Bruker AXS D8 diffractometer, in grazing incidence condition. In the setup used, the X-ray Co tube was equipped with Goebel mirror optic to obtain a parallel and monochromatic X-ray beam. X-ray diffraction (XRD) data were collected (normal incidence) between 30 and 120° with 0.03° as step size. After hot dipping the alloy exhibits an outer layer of Al–Si, remnant of the melted alloy, and an inner intermetallic layer, hereafter named the coating. The outer layer was removed by chemical etching followed by mechanical abrasion. All major surfaces were abraded on successively finer silicon carbide papers, then mechanically polished with 1 μm diamond paste and cleaned with ethanol. Uncoated samples were polished using the same procedure. Determination of the oxidation kinetics was carried out by continuous isothermal thermogravimetry tests in the temperature range of 700–900 °C for exposures up to 150 h using a CI PRECISION MK2-M5 microbalance. Oxidation tests were carried out under dry synthetic air (dewpoint below −40 °C). Characterization of the oxidation products was performed on samples used for mass-gain determination as well as on samples isothermally oxidized for different exposure times. Cross-sections were prepared by conventional metallographical techniques. To prevent scale loss during the metallographic preparation of the sample, the surfaces were successively coated first with a thin gold layer (by sputtering) and then with a thicker layer of copper (electrolytically deposited). Surfaces and cross-sections of the oxidized specimens were studied by scanning electron microscopy (SEM). Phase identification of the oxide scale was performed by X-ray diffraction (XRD) and energy dispersive X-ray microanalysis (EDX). 3. Results

b Element Concentration (%)

Coated AISI 316 LVM Heat treatment : → for oxidation tests 780  C; 6 h

70

Steel

Fe Al Si Cr Ni

A layer

B layer

60 50 40 30 20 10 0 -10

0

10

20

30

40

50

Depth (μm) Fig. 1. (a) Cross-sectional view of the coating generated after hot dipping for 120 s in Al–31Si bath and annealing at 760 °C for 6 h. (b) Depth profiling of the elements constituting the coating.

Fig. 1b shows the concentration profiles through the coating thickness, determined by EDX microanalysis. Compositional changes in the atomic concentrations of aluminum, silicon, iron and chromium through the profile clearly indicated the existence of more than one phase. The average concentrations of the outermost A layer, expressed in atomic percentage, were 49% Al, 27% Si, 17% Fe and 5% Cr. Aluminum and nickel contents increased to 55% and 4%, respectively in layer B compared to layer A, but silicon percentage decreased to 17%. Silicon in the EDX profiles was an indication that ternary Fe–Al–Si intermetallic compounds were present in the coating. It is interesting to note that nickel was not found in the coating when the AISI 316 LVM alloy was immersed in the eutectic Al–12 at.% Si alloy [11,15]. Identification of the phases coexisting in the outermost layer (layer A) was carried out by XRD. The results revealed that, at least, four intermetallic phases coexist in layer A; Al3FeSi2, Al3FeSi, (Si,Al)2Cr and CrSi2, as shown in Fig. 2. The existence of additional intermetallic phases could not be ruled out because of the presence of unidentified peaks at 56.4 and 60.8°.

3.1. Coating characterization 3.2. Mass gain curves A cross section of coated AISI 316 LVM stainless steel is shown in Fig. 1a. Four different layers adjacent to the AISI 316 LVM stainless steel could be distinguished. The coating was well adhered to the AISI 316 LVM substrate, being the total thickness of the coating about 75 μm. Unlike the undesirable tongue-like morphology developed in the hot dipping processes in pure aluminum baths [12–15] the coating was flat, continuous and free of delaminating signs. The thicknesses of the outermost layers (A & B) were 36 and 18 μm, respectively. These layers contained another second phase homogeneously distributed across their thickness.

Mass gain curves in the temperature range of 700–900 °C for coated and uncoated AISI 316 LVM alloy are presented in Fig. 3. The curves proved that an intermetallic Al3FeSi2-based coating modifies the oxidation behavior of the AISI 316 LVM alloy. The influence of the coating on the oxidation resistance of the alloy was slightly detrimental below 900 °C. Thus, mass gains of coated samples at 700 and 800 °C were two times higher than those corresponding to the uncoated material (Table 1). Nevertheless, the opposite was true at 900 °C. At this temperature, the oxidation resistance of the coating increased substantially

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Fig. 2. XRD pattern of coated AISI 316 LVM stainless steel. ▲ CrSi2 (JCPDF 35-0781), ● Al3FeSi2 (JCPDF 52-0917), ○ (Si,Al)2Cr (JCPDF 35-0781), △ Al3FeSi (JCPDF 20-0032), ? unknown phase.

(mass gain was half of the AISI 316 LVM alloy). The curves were characterized by a first transient stage which was followed by a steady state, except for the coated material during initial oxidation at 800 °C and over all the exposure at 900 °C. Most of the mass gain took place during the transient stage, especially for the coated sample. Once the steady state was achieved the kinetics obeyed parabolic laws. It is interesting to note that oxidation exponent n remained practically constant for the coated material. However, n gradually deviated from the parabolic behavior (n=0.5) towards almost-parabolic laws (n=0.65) as the temperature was raised, indicative of the less protectiveness conferred by the oxide scale formed at 800 and 900 °C. Although mass gain of the coated material at 900 °C was smaller than that found for the uncoated alloy, the curve was complex. The curve reached a first steady stage.

After some time, the scale loses its protective nature during a short period in which the mass gain increased rapidly until a new steady state obeying parabolic kinetics was again attained. This behavior was also found at 800 °C during the initial stage of oxidation, but after the first breakdown event the kinetics obeyed parabolic laws over the entire exposure. 3.3. Surface morphology 3.3.1. Uncoated material Morphologies of the oxide layers formed on AISI 316 LVM alloy after 150h of exposure in the 700–900°C temperature range are presented in Fig. 4. Oxide layer formed after exposure for 150 h at 700 °C was very uniform and consisted of a fine-grained oxide overgrown by crystals with two different morphologies; numerous small octahedron-like crystals of about 0.5–0.8 μm and scarce long prismatic crystals up to 4 μm long and 0.5 μm thick (Fig. 4a). The external scale established after exposure at 800 °C was more uniform and exclusively constituted by small like-octahedron crystals of about 1–1.5 μm (Fig. 4b). After oxidation at 900 °C for 150 h, the appearance of the oxide surface was analogous to that found at 800 °C but the uniform scale was frequently dotted by thicker nodules, as shown in Fig. 4c. Nodules were constituted by coarse faceted crystals up to 4–5 μm in size, often with large pores in the middle of the crystal (Fig. 5a). This morphology was different to octahedron-like crystals found on the uniform scale (Fig. 5b). 3.3.2. Coated material The surface of the sample oxidized at 700 °C was more irregular, exhibiting smooth elevated regions surrounded by others depressed (Fig. 6). This could be indicative that texture and/or phase distribution in the coating (mainly Al3FeSi, Al3FeSi2 and (Si,Al)2Cr) could favor oxide growth in certain preferred orientations. No differences were found among the oxide formed on top and bottom regions. The external scale consisted of uniform fine-grained oxide overgrown by numerous lenticular particles of about 1μm, as shown in the inset of Fig. 6. However, the oxide morphology after exposure at 800 °C for 150 h was totally different (Fig. 7). The surface consisted of flakes, growing preferentially

Table 1 Values of the mass gain and oxidation exponent for the AISI 316 LVM stainless steel and the coated AISI 316 LVM stainless steel after oxidation at 700, 800 and 900 °C for 150 h. T (°C) Mass gain (mg/cm2)

Oxidation exponent (n)

AISI 316 LVM Coated AISI 316 LVM AISI 316 LVM Coated AISI 316 LVM

Fig. 3. Mass gain curves after oxidation in air at 700, 800 and 900 °C for 150 h. (a) AISI 316 LVM alloy. (b) Coated AISI 316 LVM alloy.

700 800 900

0.06 0.17 1.72

0.11 0.30 0.9

0.5 0.6 0.64

0.55 0.45 0.51 (steady state)

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Fig. 5. Morphology of the oxide over the nodules (a) and over the uniform scale (b) after exposure for 150 h at 900 °C.

3.4. Cross sections 3.4.1. Uncoated material The thinness of the oxide layer formed on AISI 316 LVM after 150h at 700 °C precluded cross section observations. After 150 h of exposure at 800 °C, the scale was about 0.5 μm. Isolated coarse particles up to 1 μm in size could be observed over a continuous thin layer, as shown in Fig. 9. EDX microanalyses of the outermost isolated crystals proved

Fig. 4. Morphologies of the scales formed on the AISI 316 LVM alloy after oxidation for 150 h at different temperatures: (a) 700 °C, (b) 800 °C and (c) 900 °C.

normal to the metal surface. In addition, some cracks were also found irregularly dispersed on the scale. Flakes probably arose from further growth of lenticular particles found after exposure at 700 °C. These flakes constituted the outermost part of the oxide layer, as deduced by the existence of isolated regions, smooth dark areas, in which the flakes were absent. After oxidation at 900 °C, the morphology of the scale was identical to that observed during oxidation at 800 °C. Oxide flakes were already present after 45h of exposure, as shown in Fig. 8a. No changes in the scale morphology were observed after more prolonged exposure. Spalling during cooling revealed the development of numerous large cavities in the coating just beneath the oxide, i.e. at the coating/oxide layer interface, as shown in Fig. 8b.

Fig. 6. General view of the scale formed on coated AISI 316 LVM stainless steel after oxidation at 700 °C for 150 h. The inset shows the morphology of the scale.

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Fig. 7. General view of the scale formed on coated AISI 316 LVM stainless steel after oxidation at 800 °C for 150 h. The inset shows the platelet-like morphology of the external alumina scale.

that they were enriched in chromium (about 20 at.%), with smaller manganese and iron contents (about 10 at.% of each element). The thin innermost oxide layer was enriched in chromium and iron, concentrations around 26 at.% and 12–14 at.%, respectively. Manganese was practically absent of this continuous layer, about 1–2 at.%. Compared to 800 °C, the oxide layer considerably thickened after exposure at 900 °C for 150 h (Fig. 10). Three regions could be identified: (a) The

Fig. 8. (a) Morphology of metastable aluminas on the surface of the scale formed on coated AISI 316 LVM stainless steel after oxidation in air at 900 °C for 45 h. (b) General view of spalled regions formed on coated AISI 316 LVM oxidized in air at 900 °C for 150 h. The inset shows cavities developed at the scale/coating interface.

Fig. 9. Cross sectional view of the layer formed on AISI 316 LVM stainless steel oxidized for 150 h at 800 °C. At the bottom, detail of the inner continuous layer overgrown by isolated particles.

first one is presented in Fig. 10. It corresponded to a uniform thin scale of about 5–6 μm. EDX microanalyses indicated a complex oxide constituting the outermost part of the oxide scale whose average composition was 19 at.% of chromium, 8–9 at.% of manganese and 10–12 at.% of iron. This composition remained almost constant throughout the thickness, although it changes towards the stoichiometry of a binary chromiumrich oxide (28–29 at.% Cr and 6–7 at.% Fe) in the region close to the scale. (b) The second region corresponded to a thick uniform scale of about 15–20 μm, as shown in Fig. 11. Fig. 11c presents the depth profiling of this thick layer. The outermost part of about 7–8 μm consisted of an iron-rich layer (about 28–30 at.% of Fe) with minor amounts of chromium (about 10 at.%). The central part of the scale was about 7–8 μm. It was also enriched in iron, but the concentration was below, 22 at.%, of that found at the external part. Also the chromium content was slightly reduced from 10 to 8 at.%. Iron and chromium were partially replaced by nickel (7 at.%) and manganese (3 at.%) in the oxide. The innermost layer was very thin, 1–2 μm, and it consisted of a chromium-rich oxide (27–28 at.% Cr) containing lower amounts of iron, around 7–8 at.%. It is interesting to note that the composition of this layer was very close to that of the innermost layer of the scale found in the first region. (c) The third region corresponded to nodules irregularly distributed on the scale, as seen in Fig. 12. Nodules were thicker than the uniform oxide layer. They could be divided into two layers. The composition of the external part of the nodule, up to 25–30 μm in thickness, coincided with an iron oxide (34 at.% Fe) with low amounts of chromium

Fig. 10. Cross sectional view of the scale formed on AISI 316 LVM stainless steel oxidized for 150 h at 900 °C. The inset shows a close magnification of the thin uniform oxide layer.

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(2 at.%) and manganese (0.8 at.%). According to this stoichiometry, this layer should correspond to Fe2O3. The composition changed in the internal part of the scale, dissolving the iron oxide larger amounts of chromium (11–12 at.%). The internal layer was slightly thinner than the outer layer, up to 15–17 μm.

Fig. 11. (a) Cross sectional view of the thick scale formed on AISI 316 LVM stainless steel oxidized for 150 h at 900 °C. (b) Magnification of the scale in the region marked by the square. (c) Compositional depth profiling of the oxide layer along the line drawn in (b).

Fig. 12. Cross sectional view of the nodules formed on AISI 316 LVM stainless steel oxidized at 900 °C for 150 h.

3.4.2. Coated material Cross sectional views of the coated material were more complex due to the microstructural evolution of the coating in the course of the oxidation and the formation of oxides at the coating/AISI 316 LVM interface. After 150 h of exposure at 800 °C, the coating can be divided into four main regions, as observed in Fig. 13, whose borders are very irregular. The outermost region or Zone I consisted of a ternary Fe–Al–Si phase whose average composition is; 43–45 at.% Fe, 31–33 at.% Al, 20– 22at.% Si, and 2–4at.% Cr. This phase was slightly depleted in aluminum, around 30 at.%, in the region underlying the external oxide because this phase was the aluminum source for alumina formation. The next region or Zone II was constituted by different phases. Predominant phase (dark phase in micrograph) was a quaternary Al–Fe–Ni–Si compound with average composition; 53–55 at.% Al, 30–32 at.% Fe, 6–7 at.% Ni, 5–8 at.% Si and 2–3 at.% Cr. At the boundaries of this Al–Fe–Ni–Si phase, small particles of a white/bright phase could be identified. The composition of this phase, according to EDX microanalyses, was a ternary Cr–Si–Fe alloy with large compositional range; 34–43 at.% Cr, 30–34 at.% Si, 16– 20 at.% Fe, 5–13 at.% Al and 1–4 at.% Mo. The next two regions of the coating, Zone III and Zone IV were very similar. Major difference regarded the large volume fraction and size of heavy/bright particles, whose size and volume fraction increased in the direction of the coating/AISI 316 LVM interface. Bright particles corresponded again to Cr–Si–Fe particles, in which the chromium content was raised and iron concentration was diminished from the outermost part of the coating to the particles in the vicinity of the coating/AISI 316 LVM interface. The predominant dark phase corresponded to the Fe–Al–Si previously described, although the silicon content decreased up to 10–11 at.%. Such decrease was balanced by the increase in nickel content up to 10–12 at.%. Moreover, nickel concentration became up to 20–21 at.% in particles close to the coating/alloy interface, which was accompanied by the reduction of iron and silicon contents up to 36–37 and 5–7 at.%, respectively. At 800 °C the external oxide scale consisted of a thin aluminum-rich layer of about 1 μm, as shown in the inset of Fig. 13. The thinness of the oxide and EDX results suggested that this layer was alumina (Al2O3). The coating just beneath the oxide layer contained many cavities/ voids. The walls of cavities underlying the external oxide scale were

Fig. 13. Cross sectional view of coated AISI 316 LVM stainless steel showing the four regions developed in the coating after oxidation at 800 °C for 150 h. The inset shows a detail of the external alumina scale.

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also covered by a thin alumina layer in such a way that the entire cavity was sealed by the alumina layer, as seen in the inset of Fig. 13. This alumina layer surrounding cavity walls, however, was absent in cavities as further away of the oxide/coating interface, which were filled by mixed Al–Si oxides. At the coating/AISI 316 LVM interface, the formation of a discontinuous oxide layer up to 2–3 μm thick was noticed (Fig. 14). It is interesting to note that this oxide became even thicker than that formed externally. Isolated plate- or cubic-like particles (depending only on the orientation of the plates with respect to the metallographic plane) were present at this region. This phase could appear even deeper into the underlying alloy. These particles were enriched in nitrogen and aluminum, suggesting the formation of aluminum nitrides. The microstructure of the AISI 316 LVM alloy was also modified at the subsurface layer beneath the coating/AISI 316 LVM interface as a consequence of longterm exposure. Thus, the predominant dark phase corresponded basically to binary Fe–Cr phase containing minor amounts of other elements (70 at.% Fe, 15 at.% Cr, 4 at.% Si, 6 at.% Ni, 0.5 at.% Mo and 3 at.% Al). The white phase in micrographs was a Fe–Cr phase enriched in molybdenum and nickel. The composition of white phase was 58–60 at.% Fe, 17–18 at.% Cr, 9 at.% Si, 4–5 at.% Ni, 7–8 at.% Mo and 1–2 at.% Al. In addition, less intense white particles could be identified, with a composition not so different to those found for white particles, although with lower amounts of molybdenum and higher chromium contents (59 at.% Fe, 24 at.% Cr, 7 at.% Si, 5 at.% Ni, 3 at.% Mo and 1 at.% Al). Significant changes were observed after exposure at 900 °C for 150 h in air. Massive spalling during cooling down from oxidation temperature only left attached the external oxide layer in some isolated regions, as shown in Fig. 15. Beneath the scale, large cavities up to 10 μm in size were formed. The structure of the coating was homogeneous throughout the thickness. Different phases could be discriminated: (a) Fe–Si– Cr phase (40, 28 and 15 at.% of Fe, Cr and Si, respectively). The silicon content decreased slightly from the outer to the inner part of the coating. (b) Cr–Fe–Si phase whose composition was relatively uniform throughout the coating thickness (52 at.% Cr, 20 at.% Fe, 25 at.% Si, and minor amounts of aluminum and molybdenum). (c) Fe–Al–Si (50 at.% Fe, 27 at.% Al, 12 at.% Si, 7 at.% Ni, 4 at.% Cr). New phases were detected close to the coating/AISI 316 LVM interface. In the coating, nickel and molybdenum were concentrated in two phases. Ternary Ni–Al–Fe (36 at.% Ni, 32 at.% Al, 27 at.% Fe, 3 at.% Si, 1 at.% Cr) appearing as small bright regions at the innermost part of the coating and ternary Fe–Mo–Si (45 at.% Fe, 22 at.% Mo, 22 at.% Si, 7at.% Cr, 4at.% Al). Two types of oxides were found there. (a) Coarse isolated dark islands were located just at the coating/AISI 316 LVM

Fig. 14. Al-rich oxides and aluminum nitrides formed at the coating/AISI 316 LVM interface after oxidation at 800 °C for 150 h.

Fig. 15. Cross sectional view of the coated AISI 316 LVM stainless steel after oxidation at 900 °C for 150 h. The inset shows a detail of the external alumina layer.

interface, as observed in Fig. 16. It corresponded to an Al-rich oxide (58 at.% O, 34 at.% Al, 6 at.% Fe, 1 at.% Cr, 1 at.% Si). (b) A little more inside the subsurface layer, a bright nearly continuous oxide scale was formed (see Fig. 16). EDX microanalyses revealed that this phase was a complex oxide (62–63 at.% O, 17–18 at.% Cr, 9– 10 at.% Fe, 9 at.% Mn). It is worth noting the presence of large cavities at both sides of this oxide. A nearly homogeneous narrow band of an Fe-rich phase was present just above the oxide. This phase would be probably the alloying element source to build up the oxide. Its composition was 59 at.% Fe, 14 at.% Cr, 11 at.% Si, 8 at.% Ni, and 7 at.% Mo. Also the subsurface layer in the metallic substrate, underlying the oxide, changed with respect to the composition of AISI 316 LVM alloy. This is an indication that the coating was growing into the alloy as a consequence of diffusion, outwards or inwards, of the different elements. Basically two types of phases have been found. The first one is depleted in nickel and, especially, in chromium. A typical composition is Fe70Cr12Ni10Si5 at.% (the contribution of other elements has not been considered). The second phase is characterized by higher chromium contents with Fe68Cr21Si7Mo3 at.% as composition. In some regions this phase contained also nickel and higher chromium contents (Fe60Cr24Ni5Mo3Si7 at.%). It is interesting to note the relatively high silicon content of these phases, indicative of rapid inward silicon diffusion from the external coating. As found after exposure at 800 °C, numerous long plates or cubic particles were noticed in the subsurface

Fig. 16. Cross sectional view of the coating/AISI 316 LVM interface after oxidation at 900 °C for 150 h.

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region of the AISI 316 LVM alloy. EDX analyses proved that this phase was an aluminum nitride.

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other non-identified phases, especially after oxidation at 800 °C (see non-identified peaks in Fig. 18b). 4. Discussion

3.5. XRD 4.1. Uncoated alloy XRD patterns of samples exposed at 700 °C, taken with an incidence angle of 1°, provided mostly peaks corresponding to the metallic substrate. Intense oxide peaks were only detected after oxidation at higher temperatures. Thus, two different oxides were clearly identified in the scale formed over the AISI 316 LVM after oxidation at 800 °C for 150 h (see Fig. 17a); corundum type M2O3 (M = Cr,Fe) and spinel type MnCr2O4. M2O3 peaks were sited between those of pure α-Fe2O3 and Cr2O3, both corundum type oxides. This was indicative that iron and chromium are present in the composition of the oxide as (Fe,Cr)2O3, which is in agreement with EDX data. In addition, also the peaks corresponding to the face-centered cubic structure of the stainless steel (fcc-Fe) were found at 800 °C. The same oxides were detected in XRD patterns after oxidation at 900 °C for 150 h, as shown in Fig. 17b. However, the peaks corresponding to the metallic matrix coincided with those of the body-centered cubic iron (bcc-Fe) instead of those belonging to fcc-Fe. Also peaks associated with NiFe2O4 were found. Analysis of XRD patterns of the coated material was much more complicated because they contained many peaks which could not be ascribed to oxides (see Fig. 18). Thus, among all peaks found for the sample oxidized at 700 °C for 150 h only very less intense peaks were due to Al2SiO5 (see Fig. 18a). At higher temperatures, compositional changes inside the coating result in the appearance of new intermetallic phases. In agreement with EDX results, the scale formed after oxidation at 800 °C for 150 h should be composed by different types of alumina. Peaks associated with four kinds of alumina were found (γ-, δ-, θ- and α-Al2O3). Other reflections attributed to Al2SiO5 were also identified. After oxidation at 900 °C for 150 h two types of alumina were found composing the scale; θ- and α-Al2O3. XRD patterns also provided information about the coating. The intermetallic phases originally constituting the coating were absent in XRD patterns after long-term oxidation at 800 and 900 °C, being replaced by new intermetallic phases such as CrFe8Si, FeSi, Fe3Si, Fe9Si, Cr3Si or Fe5Si3. The coating should also contain

The oxidation behavior of AISI 316 LVM alloy is substantially modified when it is coated by hot dipping process in the Al–31Si alloy. The intermetallic coating exhibits good oxidation resistance at 700 and 800 °C, but still slightly inferior to that of the uncoated AISI 316 LVM alloy. At 900 °C, however, oxidation behavior is substantially improved because of the formation of a continuous protective alumina layer. Mass gain curves of the AISI 316LVM alloy agree rather well with the data reported in the literature for austenitic AISI 316L [16–18] or AISI 304 alloys [17–21] and ferritic FeCr9 alloy [22]. Kinetics tend to deviate from parabolic laws as the temperature is increased, indicating a gradual decrease in the protectiveness provided by the oxide controlling the oxidation. The protective character of the oxide layer is conferred by different kinds of oxides; (i) spinels appearing externally as isolated particles at 800 °C and as continuous layer at 900 °C, basically Mn1.5Cr1.5O4 and FeCr2O4 [19] and (ii) internal Cr2O3 (chromia) layer [19]. The composition of the external particles is more complex because they contain chromium, as predominant element, and equivalent amounts of iron and manganese, with Cr/(Fe + Mn) ≈ 1. In spite of the low manganese content in the alloy, this element is also concentrated in the outermost particles. This could be related either to relatively rapid manganese diffusion in the austenitic matrix, which could replace iron in the iron oxide, or rapid manganese segregation to the surface due to its high mobility in Cr2O3 lattice [23]. Formation of iron-rich oxides has been reported during the initial stages of oxidation at 700 °C in Fe–9Cr alloy, although the outermost oxide became also enriched in manganese and chromium in the course of the oxidation [22], as observed in the present study. This suggests that iron oxides are probably formed during the earliest stages of oxidation as a result of the high iron content in the stainless steel. This, added to the close packed structure of the austenitic matrix, should prevent the rapid establishment of a protective chromia (Cr2O3) layer. Then, other elements like chromium

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2θ (°) Fig. 17. XRD patterns of AISI 316 LVM samples oxidized for 150 h. (a) 800 °C (incidence angle of 2°), (b) 900 °C (incidence angle of 10°). □ MnCr2O4 (JCPDF 75-1614), ■ Cr2O3 (JCPDF 38-1479), ● Fe2O3 (JCPDF 33-0664), ○ NiFe2O4 (JCPDF 54-0964).

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2θ(°) Fig. 18. XRD patterns of coated AISI 316 LVM samples oxidized for 150 h. (a) 700 °C (incidence angle of 1°), (b) 800 °C (incidence angle of 2°), (c) 900 °C (incidence angle of 10°). △ Al2SiO5 (JCPDF 72-1447), □ (Si,Al)2Cr (JCPDF 35-0781), ■ CrSi (JCPDF 51-1356), Al3FeSi (JCPDF 20-0032), ○ Al3FeSi2 (JCPDF 52-0917), ● Al3Fe (JCPDF 45-1178), Cr3Si (JCPDF 7-0186), ✱ θ-Al2O3 (JCPDF 86-1410), ✚ δ-Al2O3 (JCPDF 16-0394), ✖ γ-Al2O3 (JCPDF 10-0425) ▲ α-Al2O3 (JCPDF 85-1337), Fe3Si (JCPDF 65-0994), ❶ FeSi (JCPDF 38-1397), ❷ Fe5Si3 (JCPDF 65-3593), ❸ Fe9Si (JCPDF 65-9130), ❹ Fe8CrSi (JCPDF 65-5584), ? unknown phase.

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and manganese can dissolve in the lattice of the iron-rich oxides because of the identical corundum structure (Me2O3) of manganese, chromium and iron oxides. The structure of the corundum oxide transforms into the spinel as manganese and chromium are gradually dissolved in the Fe2O3 lattice, as deduced from intense peaks found in XRD patterns. Formation of this complex oxide, whose structure agrees with that of spinel-like oxide, would favor subsequent development of a protective Cr2O3 layer through the well known second-getter effect. Microanalyses revealed, that this layer is (Cr,Fe)2O3 layer, with a Cr/Fe ratio close to 2, rather than pure Cr2O3. In any case, (Cr,Fe)2O3 is effective slowing down the oxidation rate up to very low values. Formation of spinel-like oxides is restricted to isolated particles on the surface developed at 700 °C, whose growth is exacerbated during oxidation at 900 °C. Thus, the typical octahedron-like morphology of spinel oxides covers the entire surface of the uniform scale developed after exposure at 900 °C. This uniform scale is only disrupted by the presence of numerous thick nodules. This clearly indicates that the large thickness of the outermost spinel-like oxide with the increase of the oxidation temperature could be related to higher diffusivity of manganese and chromium in the metallic matrix in such a way that outwards manganese and chromium atom flows are enough to form the spinel layer. Furthermore, the chromium spent for building up the spinel layer also diminishes chromium concentration in the underlying metal matrix, preventing further growing of protective chromia layer. Consequently, longer times should be required to develop again the

protective Cr2O3 layer at 900 °C. Only when the thickness of the spinel layer reaches a critical value which reduces the oxygen entering in the alloy, the chromium atom flow towards the spinel/metal interface is enough to promote the formation of a thin (Cr,Fe)2O3 layer. Similar evolution has been found during oxidation of Fe–9Cr alloy at 700 °C [22]. This mechanism should result in the formation of the predominant uniform oxide layer. It is interesting to note that in the subsurface region underlying the thin uniform oxide layer, numerous small Si-rich particles (probably SiO2) are present. Their formation can be explained considering that the oxygen partial pressure fixed by the protective (Cr,Fe)2O3 layer at the oxide layer/AISI 316 LVM interface is higher than that required for SiO2 formation. Another two types of oxide patterns, which are usually disrupting the continuity of the thin uniform oxide layer, can be found in the uncoated AISI 316 LVM oxidized at 900 °C. This would be indicative that oxidation could proceed in different ways over other regions of the alloy. The first pattern appears as thick nodules consisting of thick external Fe2O3 and internal (Fe,Cr)2O3 layers. Unlike the thin uniform layer, the iron content is higher than chromium content in the internal (Fe,Cr)2O3 layer. Negligible concentrations of manganese were dissolved in these nodules irrespective of the nature of the oxide. The second pattern corresponds to thick spinel regions, similar to that formed in the predominant uniform oxide layer. The spinel is located outwards, over a relatively thick oxide of about 5–6μm enriched in iron and nickel. Beneath this intermediate layer, a thin layer was developed, being the outer part

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enriched in iron and the inner part chromium-rich. EDX data suggest that the external part could be (Fe,Cr)2O3 and the internal would be (Cr, Fe)2O3. Discrepancies in the type of oxide could result from the mass flow of the different elements (chromium, iron, manganese) in the austenitic regions towards the different regions being oxidized. Formation of the protective spinel requires high chromium and manganese diffusion from the austenitic matrix. It seems that such matter flow could not be kept over the entire surface so some regions could remain depleted in manganese and chromium. Unfavorable orientations might delay diffusion in such a way that the alloy could not supply enough manganese and/or chromium to induce there the rapid formation of protective spinels/chromia. The result is the establishment of less protective oxides. Subsequent enrichment of the subsurface zone in nickel should promote the formation of NiFe2O4, as shown in depth profile of Fig. 11. Probably, the higher mass gain associated with these non-protective regions should account for the deviation from parabolic kinetics measured in mass gain curves. 4.2. Coated alloy According to the results, the oxidation of coated AISI 316 LVM proceeds according to three main stages, as illustrated in Fig. 19. (1) Formation of an external alumina layer which is accompanied by cavity development at the alumina/coating interface, within the coating and at the coating/AISI 316 interface. (2) Nitride precipitates are formed inside the AISI 316 substrate close to the coating/AISI 316 interface. (3) Oxides are formed at the coating/AISI 316 interface. The good oxidation resistance of the intermetallic coating alloy is related to the ability to form a protective alumina layer at the 700–900 °C temperature range. The aluminum concentration on the surface of the intermetallic coating is high enough to induce the formation of alumina layers. As found in other alumina-forming alloys, the type of alumina depends on the temperature. At 700 °C, XRD patterns indicated that γ-Al2O3 was exclusively formed. Low mass gains proved the

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protective nature of γ-Al2O3, as reported during the oxidation of chromium-containing Ni3Al alloys [24,25]. At 800 and 900 °C, different alumina morphologies could be distinguished on the surface of oxidized samples. Needle-like morphology observed in the surface indicated that the external surface was constituted by metastable δ- and θ-Al2O3. Underlying these metastable aluminas should be located γ-alumina. These aluminas will tend to transform gradually into α-Al2O3 with increasing temperature and/or exposure time. In fact, the external needle-like alumina formed initially at 900 °C thickened for long-term exposures, being converted into platelet-like crystals. Although the γ-Al2O3 is not usually considered as protective as the α-Al2O3, it can provide excellent oxidation resistance, almost comparable to that of α-Al2O3, between 700 and 900 °C [25]. Other factor, affecting enormously the oxidation behavior of the coated material, was microstructural changes occurring in the coating. The long-term exposure at 800 and 900 °C, temperatures comparable or even higher to that used for producing the coating, induced that new phases could appear/disappear in the course of the oxidation as a consequence of intense mass flux of the different elements constituting the coating and the stainless steel. Mass transport can take place through two different phenomena: 1) Diffusion induced by concentration gradients existing in the coating and alloy during long-term exposure at the oxidation temperature. Elements constituting the AISI 316 LVM stainless steel should tend to diffuse outwards. On the other hand, coating is enriched in aluminum and silicon, so these elements should tend to diffuse towards the alloy. Thus after exposure at 800 °C, the double layered structure of the coating is kept, but some changes regarding the composition of the phases and the appearance of new phases are viewed at the outer and inner parts of the coating (see for example the XRD pattern presented in Fig. 18b). At 900 °C, the coating is quite homogeneous throughout the thickness. Moreover, the thickness is increased because the coating grows several microns into the alloy. Silicon and aluminum diffusion should be very fast because

Fig. 19. Oxidation mechanism of the coated AISI 316 LVM stainless steel. (a) Initial state of the coating prior to oxidation. (b) Cavities developed at the alumina/coating interface, inside the coating and at the coating/AISI 316 LVM interface. (c) Nitride formation inside the AISI 316 LVM substrate close to the coating/AISI 316 LVM interface. (d) Oxide formation at the coating/ AISI 316 LVM interface which is accompanied by compositional homogenization of the coating throughout the thickness.

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noticeable amounts of these elements are detected in the vicinity of the coating/metal interface inside the stainless steel. 2) Diffusion induced by the formation of the oxide scale. Formation of alumina layer could deplete the subsurface layer beneath the scale in aluminum. Thus, aluminum should also tend to diffuse outwards from neighbor regions in the coating to balance such depletion. In addition, oxygen or even nitrogen can diffuse into the coating, as will be further discussed. Atom fluxes of the different elements are accompanied by net flux of vacancies in the opposite sense (Kirkendall effect). Vacancies could coalesce and condense as small voids which should act as effective sinks for subsequent vacancy flux [26]. The result is the gradual development of large cavities at the scale/coating interface. Furthermore, cavities were also observed embedded inside the coating in regions adjacent to the oxide scale, as a result of massive coalescence of vacancies at phase boundaries. In the same way, some cavities were formed at the coating/metal substrate interface in a similar way to that reported during oxidation of hot-dipping Al–Si coating on low carbon steel [2]. Obviously, cavity generation is minimized in the stainless steel because of its larger melting point compared to that of the coating, so diffusivity was much slower. Presence of cavities at these two regions affects oxidation resistance in two ways; (a) Cavities act as effective stress concentrators, reducing the adherence of the oxide scale to the substrate. If growth stresses are high enough and they cannot be released by plastic deformation in the intermetallic coating, cracking of the scale and further detaching can occur in the course of oxidation. This agrees with crosssectional observations in which parts of the alumina scale can be found piled. Moreover, some parts of the coating are located at the external surface, indicating that detachment occurred within the coating. The bare regions are newly oxidized, resulting in accelerated mass gain until the protective alumina layer is again formed. This can be visualized in mass gain curves at 800 and 900 °C as short periods where significant mass gain takes place, followed by long periods in which mass gain is considerably slowed down. At 900°C, large thermal stresses generated during cooling result in massive spalling. (b) Oxygen and/or nitrogen can condensate inside them to form oxides/nitrides within the coating/alloy, as observed in cross sectional views. The nature of the oxide present in cavities depends on their location. Some cavities formed on the surface are totally filled by oxides other than alumina, usually enriched in silicon and aluminum, because the initial formation of the external alumina layer causes aluminum depletion on the subsurface layer. On the contrary, other cavities are empty and a thin alumina layer grows on their walls. In this case, the oxygen pressure in the cavity, sealed from the outside by the external alumina layer, could rise until the alumina layer is formed. On the other hand a mixed Al–Si oxide was developed on cavities embedded inside the coating. Fig. 20 displays the different stages leading to formation and subsequent oxidation of large cavities developed inside the coating. More surprising is the formation of aluminum nitrides inside the alloy, beneath the coating/alloy interface. Thermodynamical stability of aluminum nitrides is below that corresponding to aluminum oxides, as demonstrated by the formation of alumina on the surface. Aluminum nitride formation implies that nitrogen should be transported from the gas atmosphere to the coating/metal interface faster than oxygen. This could be only possible if the oxygen diffusing into the alloy is consumed before it can arrive to the coating/metal interface. Gradual condensation of oxygen gas in the cavities developed at the outermost part of the coating will reduce the net oxygen flux into the metal. Moreover, oxygen consumption inside the cavities by forming alumina or Al–Si oxide contributes still more to decrease the oxygen flux into the alloy. Such reaction does not occur in the case of nitrogen, so this element can diffuse freely into the coating towards the coating/alloy interface. There nitrogen can condensate in microvoids generated close to this interface, reacting with aluminum to form nitrides. At 900 °C, the process

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Fig. 20. Scheme illustrating oxide formation within the cavities developed in the coating. (a) Oxygen and nitrogen diffuse inwards towards the coating/alloy interface. (b) Oxygen condensates within the cavities while nitrogen continues diffusing inwards. (c) Once the cavity is filled by an Al–Si oxide, the oxygen continues to diffuse inwards.

is accelerated, but a major difference is the generation of oxides at the coating/metal interface. This oxide coincides with the spinel formed over the uncoated alloy. Although aluminum arising from the coating is near this interface, its concentration was not enough to form an alumina layer. Moreover, previous aluminum nitride formation reduces still more the aluminum content at these regions, hindering alumina formation. The different stages for the formation of aluminum nitrides are illustrated in Fig. 21. 5. Conclusions From the present study the following conclusions can be drawn: Hot-dipping into a molten bath containing Al–31 at.% Si of AISI 316 LVM stainless steel improves the oxidation resistance only at 900 °C but not at lower temperatures. The rapid development of chromia layers confers excellent oxidation resistance to the uncoated AISI 316 LVM alloy while the good oxidation resistance of the intermetallic coating is related to the ability to form a protective alumina layer. At 700 and 800 °C, the chromia layer is more protective than the scale consisting predominantly of less protective aluminas (γ-, δ- and θ-Al2O3) established on the coated alloy. At 900 °C, however, the rapid transformation of less protective aluminas into the protective α-Al2O3 enhances considerably the oxidation resistance of the coated material compared to the uncoated AISI 316 LVM alloy. The oxidation rate of the AISI 316 LVM uncoated stainless steel becomes two times that of the coated alloy because the chromia layer cannot form over the entire surface of the material.

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Fig. 21. Sequence of nitride and oxide formation at the coating/AISI 316 LVM interface. (a) Nitrogen arrives to the coating/AISI 316 LVM interface. (b) Aluminum nitrides are formed at the coating/AISI 316 LVM interface. (c) Coating continues to grow into the stainless steel while oxygen arrives to the coating/AISI 316 LVM interface. (d) Oxygen condensates at cavities. (e) Oxide is formed at the coating/AISI 316 LVM interface. Simultaneously, nitrogen diffuses inwards leading to the formation of new aluminum nitrides into the AISI 316 LVM alloy.

In the course of the oxidation, especially at 800 and 900 °C, the microstructure of the coating evolves continuously. Inward diffusion of aluminum and silicon promotes the development of cavities inside the coating in the region beneath the alumina scale. Oxidation inside the cavities slows down oxygen flux towards the coating/ AISI 316 alloy interface in such a way that nitrogen can arrive first to the coating/AISI 316 alloy interface instead of oxygen, promoting the formation at this region of aluminum nitrides not observed during the oxidation of the uncoated AISI 316 LVM alloy. Acknowledgments The authors wish to express their thanks for the financial support of the Spanish MAT 2009-14695-C04-02-04. References [1] [2] [3] [4]

Y.Y. Chang, C.C. Tsaur, J.C. Rock, Surf. Coat. Technol. 200 (2006) 6588. C.J. Wang, S.M. Chen, Surf. Coat. Technol. 200 (2006) 6601. T. Ros-Yañez, Y. Houbaert, J. Appl. Phys. 91 (2002) 7857. T. Ros-Yañez, Y. Houbaert, O. Schneeweiss, J. Asensio-Lozano, M. Prado-García, Rev. Metal. Madrid 36 (2000) 339. [5] Y.L. Yang, Z.Y. Cao, Y. Qi, Y.B. Liu, Adv. Mater. Res. 97–101 (2010) 530.

[6] M.B. Lin, C.J. Wang, A.A. Volinsky, Surf. Coat. Technol. 206 (2011) 1595. [7] C.J. Wang, J.W. Lee, T.H. Twu, Surf. Coat. Technol. 163–164 (2003) 37. [8] J.R. Davies, Aluminium and Aluminium Alloys, ASM Speciality Handbook, ASM International, Materials Park, OH, 1993. 181. [9] J.R. Davis, Stainless Steels, ASM Speciality Handbook, ASM International, Materials Park, OH, 1994. [10] G.Y. Lai, High-temperature Corrosion of Engineering Alloys, ASM International, Materials Park, OH, 1990. [11] E. Frutos, J.L. González-Carrasco, C. Capdevila, J.A. Jiménez, Y. Houbaert, Surf. Coat. Technol. 203 (2009) 2916. [12] H. Glasbrenner, K. Stein-Fechner, J. Konys, Fusion Eng. Des. 51–52 (2000) 105. [13] H. Glasbrenner, K. Stein-Fechner, J. Konys, J. Nucl. Mater. 283–287 (2000) 1302. [14] A. Bahadur, O.N. Mohanty, Mater. Trans. JIM 32 (1991) 1053. [15] E. Frutos, R. Martínez-Morillas, J.L. González-Carrasco, N. Vilaboa, Intermetallics 19 (2011) 260. [16] V. Shankar Rao, Oxid. Met. 77 (2012) 9. [17] A. Paúl, R. Sánchez, O.M. Montes, J.A. Odriozola, Oxid. Met. 67 (2007) 87. [18] H. Buscail, S. El Messki, F. Riffard, S. Perrier, C. Issartel, Oxid. Met. 75 (2011) 27. [19] J. Botella, C. Merino, E. Otero, Oxid. Met. 49 (1998) 297. [20] C. Issartel, H. Buscail, E. Caudron, R. Cueff, F. Riffard, S. Perrier, P. Jacquet, M. Lambertin, Corros. Sci. 46 (2004) 2191. [21] A.M. Huntz, A. Reckmann, C. Haut, C. Sévérac, M. Herbst, F.C.T. Resende, A.C.S. Sabioni, Mater. Sci. Eng. A447 (2007) 266. [22] H. Asteman, M. Spiegel, Corros. Sci. 50 (2008) 1734. [23] A. Vesel, M. Mozetic, A. Drenik, N. Hauptman, M. Balat-Pichelin, Appl. Surf. Sci. 255 (2008) 1759. [24] D.F. Susan, A.R. Marder, Oxid. Met. 57 (2002) 131. [25] P. Pérez, P. Adeva, Oxid. Met. 65 (2006) 15. [26] J.D. Kuenzly, D.L. Douglass, Oxid. Met. 8 (1974) 139.