Henri Buscail

Raphaël Rolland

Sébastien Perrier

INFLUENCE OF WATER VAPOUR ON 316L OXIDATION AT HIGH TEMPERATURE -IN SITU X-RAY DIFFRACTION

Water vapour mixed with air (10 vol.% H 2 O) has little effect on the AISI 316L stainless steel oxidation under isothermal conditions at 800 and 900 °C. The oxide scale is composed of Cr 2 O 3 and Mn 1.5 Cr 1.5 O 4 located at the external interface. Results show that the presence of water lowers the specimen's mass gain. At 1000 °C, the effect of water vapour on the scale structure is deleterious. In situ X-ray diffraction permits to show that the breakaway is due to iron oxidation starting after about 30 hours. This leads to an external Fe 2 O 3 scale growth and an internal multilayered FeCr 2 O 4 scale formation.

Résumé -L'air contenant 10 vol.% H 2 O a peu d'effet sur l'oxydation de l'alliage AISI 316L à 800 et 900°C. La couche d'oxyde est alors constituée de Cr 2 O 3 et de Mn 1.5 Cr 1.5 O 4 qui est situé à l'interface externe. Les résultats montrent que la présence de vapeur d'eau diminue la vitesse de prise de masse. A 1000 °C la vapeur d'eau a un effet désastreux sur la structure de la couche d'oxyde. La diffraction des rayons X in situ à 1000°C sous air humide a permis de montrer que l'oxydation du fer débute après environ 30 heures d'oxydation. Ceci conduit à la formation d'une couche externe de Fe 2 O 3 et une couche interne stratifiée de FeCr 2 O 4 .

INTRODUCTION

The AISI 316L austenitic stainless steel is widely used in industrial applications due to its good creep properties and oxidation resistance provided by the high alloy chromium content [1][2][3][4][5][6][7][8]. Several studies proposed that, at high temperature, chromia-forming alloys such as AISI 316L SS exhibit an oxide growth leading mainly to Cr 2 O 3 and Mn 1.5 Cr 1.5 O 4 formation [9][10][11]. It has also been demonstrated that yttria addition leads to an increased oxidation resistance at 800°C [12]. The role of molybdenum on the AISI 316L oxidation mechanism has been already discussed. It has been demonstrated that molybdenum plays a similar protective role as the one provided by silicon. In dry air, it hinders iron oxidation and leads to a better keying of the chromia scale in the alloy [13][14][15]. Water vapour is present in many industrial gases. It is a constituent of combustion gases and can affect high temperature corrosion of heat exchangers or engines. Many investigations have been conducted to increase solid oxide fuel cells performance and increase the corrosion resistance of metallic interconnects [16,17].

Water vapour is also present in power-generation systems and increasing service temperatures leads to corrosion problems. Water vapour effects have already been described in the literature [18][19][20][21]. Several studies were conducted on chromia forming alloys [22][23][24][25][26]. Generally, water vapour has a major effect on the growth rate of oxide scales. Increasing the amount of water vapour in the atmosphere reduces the time to the onset of breakaway oxidation [27]. Galerie et al. have explained that the breakaway on a Fe-15%Cr alloy was induced by rapid growth of hematite at the metal/chromia interface, at 800-1000 °C, in Ar-15 vol.% H 2 O [23]. But in many cases, the chromia scale failure was related to the formation of the CrO 2 (OH) 2 and Cr(OH) 3 volatile hydroxides [28][29][30]. Evaporation can lead to the alloy chromium depletion resulting in the formation of an iron-rich non-protective scale. Young proposed that water vapour affects chromia scale growth similarly to what is observed with reactive metal elements [31]. Both segregate to oxide grain boundaries. Grain boundary cation diffusion is then largely suppressed increasing the oxygen internal diffusion and refining the oxide grain structure. On the other hand, he also explained that in short-term laboratory experiments at temperatures above 800 °C, chromium volatilization could usually be neglected, even in O 2 /H 2 O gas mixtures. Refining of the oxide grain structure under wet conditions has also been observed by Jacob et al. after pure chromium oxidation [32]. Recently, Othman et al. proposed that water vapour has no significant effect if a large chromium reservoir induces the growth of a continuous Cr 2 O 3 -scale layer [33]. The ability of Fe-Cr alloys to form a protective oxide scale was related to the chromium concentration at the alloy surface. When the concentration of chromium is too low, iron-rich oxides form, inducing breakaway oxidation. Rahmel et al. suggested that an H 2 /H 2 O gas mixture can be formed in voids within the scale then facilitating rapid inward transport of oxygen [18]. From the literature review it can be stated that owing to its relatively high chromium content and alloying elements, AISI 316L is oxidation resistant at temperatures up to 1000 °C in dry environments and in isothermal conditions.

The aim of the present work is to determine the effect of water vapour on the oxidation mechanism of the AISI 316L alloy compared with dry air experiments. In situ X-ray diffraction with flowing moist air (10 vol.% H 2 O) will provide useful information on the oxidation process at operating temperatures. It will be determined if a change in the structural composition of the layer occurs during the oxidation test and if any oxide phase transitions occur during cooling to room temperature.

EXPERIMENTAL

The material used in the present work is an AISI 316L austenitic stainless steel. The alloy composition is given in table I. The specimens, provided by APERAM Alloys Imphy, were polished on SiC paper up to 800 grade, then washed with ethanol and finally dried prior to isothermal oxidation at 800, 900 and 1000°C. The specimens are 1.5 mm thick and show a total surface area of about 6 cm 2 .

The kinetic results obtained under isothermal conditions were recorded during 96 h using a Setaram TG-DTA 92-1600 microthermobalance. The in situ characterization of the oxide scales was carried out in a high temperature MRI chamber adapted on an X-ray Philips X' PERT MPD diffractometer (Copper radiation, λ α K = 0.154 nm). The diffraction patterns are registered every hour at 1000 °C in 10 vol.% H 2 O flowing air.

X-ray diffraction (XRD) conditions were: Θ-2Θ scan, step 0.02°, range from 10 to 80°, 3 s counting time. The morphology of the external interface as well as the cross sections was observed using a Scanning Electron Microscope (SEM) JEOL JSM-7600F. The analysis of the scale was performed with a LINK energy dispersive X-ray spectroscopy (EDXS). The EDXS point analyses were performed with an electron probe focused to 1 µm spot. The water vapour experiments were performed in air containing 10 vol.% H 2 O. The flowing air (8 L/h flow rate) is saturated through boiling water. Then, the water bath humidifier maintained at 50 °C controls the water vapour ratio by condensation of the excess water vapour. All the connecting tubes are maintained at 100°C in order to avoid any water condensation out of the furnaces. A schematic drawing of the rig has been provided in a previous paper [34]. Similar gas humidifiers were used either for TGA and wet in situ DRX (MRI chamber). Dry air testing was performed in flowing air at atmospheric pressure. The weight change of the samples was determined after each oxidation test on a balance with an accuracy of 0.01 mg. II. At 800 °C, the oxidation rates are similar under dry and wet air. At 900 °C, the oxidation rate is slightly higher in dry air than in wet air. At 1000 °C, a parabolic regime is followed in dry air whereas in wet air, 10 vol.% water vapour induces an important increase of the oxidation rate after 30 hours oxidation. At 1000 °C, the oxidation rate recorded before the breakaway regime is similar in wet air and dry air. An apparent activation energy (E a ) was calculated from the slope of the Arrhenius plots of the k p values, E a = 208 ± 30 kJ mol -1 under dry conditions and E a = 241 ± 30 kJ mol -1 in wet air. 2) in order to identify the elements constituting the oxide scale. After 96 h oxidation in dry and wet air at 800, 900 and in dry air at 1000 °C, the scale is adherent. EDXS analyses performed on the cross section show that manganese chromite is located at the scale/gas interface. The middle part of the oxide scale is mainly composed of chromia. Some internal silicon oxidation is observed along the alloy grain boundaries at temperatures higher than 800 °C. No influence of water vapour is observed on the silicon internal oxidation process. Nevertheless, it appears that in wet air the oxide scale is more convoluted. At 1000 °C, in wet air, the outer scale corresponds to hematite. This oxide subscale spalled off during cooling to room temperature (30 °C/minute) and is not observed on the cross section. The inner scale is multilayered and shows alternatively FeCr 2 O 4 scales and non-oxidized nickel-rich metallic layers. EDXS analyses performed in the metallic substrate, under the oxide scale, show that in the upper 2 µm depth the chromium content is 14.2 ± 0.3 wt.%. The 17.7 wt.% bulk content is reached at 15 µm deep inside the alloy. The alloy chromium ratio close to the interface is the same whatever the temperature and gaseous environment. It then appears that water vapour does not lead to important chromium depletion inside the metallic substrate.

In situ X-Ray Diffraction results in wet air at 1000°C

The oxide scale formed on AISI 316L was analysed by in situ XRD on the metallic substrate at 1000°C, in wet air (10 vol.% H 2 O) during 65 hours. alloy surface and Cr 2 O 3 peaks are still observed. After 65 hours oxidation, the hematite scale is thick and the FeCr 2 O 4 inner oxide is hardly detected. After cooling to room temperature, figure 3 shows that the XRD pattern obtained does not show the presence of hematite anymore. FeCr 2 O 4 and the alloy are the only phases observed on the diffraction pattern due to the complete hematite scale spallation during cooling. In situ XRD results at 1000°C, in dry air, were already presented in a previous paper [13]. It was shown that Mn 1.5 Cr 1.5 O 4 and Cr 2 O 3 were detected on the alloy surface all along the oxidation test in dry air and after cooling to room temperature.

DISCUSSION

Kinetic results have been obtained on AISI 316L between 800 at 1000 °C under isothermal conditions (figure 1). The parabolic behaviour was always followed in this temperature range under dry conditions and under wet conditions at 800 and 900 °C. This permitted the calculation of the parabolic rate constants k p at each temperature (table II). The parabolic rate constant at 1000 °C in moist air was calculated from the kinetic curve obtained before 24 hours oxidation. At 1000 °C, in dry air, the k p value is calculated from the data obtained during 96 h oxidation. An apparent activation energy (E a ) was calculated from the slope of the Arrhenius plots of the k p values, E a = 208 ± 30 kJ mol -1 under dry conditions and E a = 241 ± 30 kJ mol -1 in wet air. These values are close to the one generally proposed in the literature when Cr 2 O 3 is formed on the alloys in this temperature range [32,35]. The apparent activation energies are in good agreement with the SEM cross-section observation (figure 2) showing that the scale was mainly composed of a continuous chromia scale, acting as a diffusion barrier. According to SEM and XRD analyses, the good kinetic behaviour at 800 and 900°C, in dry and moist air, appears to be mainly due to the absence of iron containing oxides in the scale. Figure 1 also shows that water vapour slightly decreases the oxidation rate relative to those registered in dry conditions and the apparent activation energy is higher in wet air. As proposed by other authors it could be due to weight losses related to the formation of the volatile hydroxide CrO 2 (OH) 2 [28][29][30]. At 1000 °C, the scale remains protective during the oxidation test in dry air, (figure 1 and figure 2). Mn 1.5 Cr 1.5 O 4 and Cr 2 O 3 are present all along the oxidation test. In wet air, oxidation at 1000 °C firstly leads to Mn 1.5 Cr 1.5 O 4 and Cr 2 O 3 formation. Nevertheless, breakaway corrosion occurred after 30 hours oxidation. In situ X-ray diffraction permit to analyse the oxides formed on AISI 316L during isothermal oxidation at 1000 °C in moist air (figure 3). The breakaway oxidation is due to the iron oxidation starting after 30 hours oxidation. This leads to an external hematite scale growth and an internal multilayered FeCr 2 O 4 scale formation. Some authors have observed that water vapour can induce a breakaway during high temperature oxidation of iron-based chromia-forming alloys [20,23]. Then, it was observed that it depends upon the oxidation time, testing conditions and the chromium alloy content [23]. As discussed above, chromium volatilization by formation oxy-hydroxides CrO 2 (OH) 2 can be envisaged [36,37]. Some authors also envisaged the transport of H 2 O into the oxide scale by the diffusion of hydrogen as H + , OH -or H 2 O in the oxide scale [38,39]. In this study, manganese chromite and chromia were observed during the initial oxidation stage at 1000°C under wet air. Nevertheless, after about 30 hours oxidation, the protective chromia scale fails and breakaway appeared. It is proposed that water dissociation occurred on the specimen surface [23,40]. This dissociation is more important at 1000 °C than at lower temperature because water dissociation is endothermic. Then, the diffusion of hydrogen-containing species could occur. [42]. The low H + and OH -radii permit their diffusion through the oxide scale. Then, increase of protons in the scale changes the oxide defect chemistry and could be responsible for the change in the growth mechanism. It can then be considered that the reaction of H + (and/or H 2 O) with the metallic substrate under the chromia scale form H 2 and wüstite [36,37]. Thermodynamic data indicate that, at 1000 °C, wüstite is more stable than H 2 O. Mikkelsen et al. also explained that addition of water vapour favours the formation of wüstite [43]. When FeO and chromia are present together they react to form FeCr 2 O 4 wich explains that wüstite does not accumulate at the internal interface and is not detected by XRD [44]. As observed by in situ XRD (figure 3), iron ions also diffuse through the scale to form Fe 2 O 3 at the external interface, then inducing the breakaway oxidation.

CONCLUSION

In situ X-ray diffraction was used to identify the oxides formed on the AISI 316L stainless steel during isothermal oxidation at 1000 °C in moist air (10 vol.% H 2 O). Our results show the presence of the Mn 1.5 Cr 1.5 O 4 and Cr 2 O 3 chromia growing simultaneously on the specimen surface before 30 hours oxidation. After about 30 hours oxidation, Fe 2 O 3 is growing on the specimen surface at the external interface and internal corrosion leads to a multilayered FeCr 2 O 4 scale. Hematite, located at the external interface, spalled off easily during cooling to room temperature. Water vapour clearly promots iron oxidation at 1000°C. At 800 and 900°C, 10 vol.% H 2 O in air has no detrimental effect on the AISI 316L stainless steel isothermal oxidation. The oxide scale is always composed of Cr 2 O 3 with a small amount of Mn 1,5 Cr 1,5 O 4 at the external interface. The lower oxidation rate can be related to chromia evaporation in wet air.

Figure 1 .

 1 Figure 1. Mass gain curves obtained after 96h isothermal oxidation of AISI 316L SS under dry air and wet air (10 vol.% H 2 O) at 800, 900 °C, and 1000°C.

3. 3 .

 3 X-Ray Diffraction results after isothermal oxidation XRD patterns obtained from the AISI 316L SS sample after 96 h isothermal oxidation at 800, 900, 1000 °C in dry air reveal the presence of Mn 1.5 Cr 1.5 O 4 (JCPDS 33-0892) and Cr 2 O 3 (JCPDS 38-1479) after oxidation at all temperatures. The relative intensity of the alloys peaks is smaller at 1000 °C due to the higher oxide thickness. After wet air oxidation, Mn 1.5 Cr 1.5 O 4 and Cr 2 O 3 are present on the alloy surface after oxidation at 800 and 900°C. At 1000°C, XRD analysis of the remaining oxide only shows the presence of FeCr 2 O 4 (JCPDS 34-0140) because the external Fe 2 O 3 scale totally spalled off during cooling at 30°C/minute. Fe 2 O 3 scale pieces have found fallen down out of the crucible.

Figure 3

 3 shows the presence of the Mn 1.5 Cr 1.5 O 4 (JCPDS 33-0892) and Cr 2 O 3 (JCPDS 38-1479) during the 30 first hours of oxidation. It appears that both oxides nucleate simultaneously on the specimen surface during the first hour of oxidation and remain present together on the surface. No other oxides were formed during the 30 first hour oxidation. After 36 hours oxidation, hematite Fe 2 O 3 (JCPDS 24-0027) is detected on the

Figure 2 .

 2 Figure 2. SEM cross-section on the AISI 316L SS oxidized at 800, 900, 1000°C in dry air or wet air (10 vol.% H 2 O) for 96 h.

Figure 3 .

 3 Figure 3. In situ XRD patterns obtained during the oxidation of the 316L alloy at 1000°C under wet air (10 vol.% H 2 O).

Table I .

 I Composition of the AISI 316L SS (weight %)The mass gain curves per unit area are given in figure 1. All isothermal thermogravimetric analyses were carried out during 96 h at 800, 900 and 1000 °C in ambient air and wet air (10 vol.% H 2 O). A parabolic regime was observed throughout the 96 h oxidation tests except when testing at 1000 °C in wet air. The

	AISI 316L	Fe	Cr	Ni	Mo	Mn	Si	C
	Content	Balance 17.7	10.9	2.16	1.57	0.55	0.023
	3. RESULTS
	3.1. Isothermal kinetics

p k values are given in table

Table II .

 II k p values obtained on the AISI 316L SS after dry air and wet air (10 vol.% H 2 O) oxidation between 800 at 1000 °C.

	Temperature	k p (

g 2 cm -4 s -1 ) Dry air k p (g 2 cm -4 s -1 ) Wet air

	800°C	4.8 ± 0.1 10 -14	3.5 ± 0.1 10 -14
	900°C	5.8 ± 0.110 -13	2.6 ± 0.1 10 -13
	1000°C	1.9 ± 0.1 10 -12	2.5 ± 0.1 10 -12 (before 24h)
	3.2. Scale morphology
	SEM examinations were carried out on the specimen cross section (figure

  Transport of H 2 O into the oxide scales is performed by the incorporation of hydrogen in the form of H + or H 2 O [38, 39, 41]. It is assumed that protons (H + ) from water vapour might dissolve in the oxides and diffuse at grain boundaries. Essuman et al. proposed a model to explain the internal oxidation of chromium alloys [20, 21]. Ardigo et al. performed gold and deuterium marking experiments and have evidenced that Fe 3 O 4 and Cr 2 O 3 scales are permeable to hydrogen


Henri Buscail

Raphaël Rolland

Christophe Issartel

INFLUENCE OF WATER VAPOUR ON THE ZIRCALOY-4 OXIDATION AT HIGH TEMPERATURE -STRESS DETERMINATION BY IN SITU X-RAY DIFFRACTION

    Influence of water vapour on the zircaloy-4 oxidation at high temperature -Stress determination by in situ X-Ray diffraction

INTRODUCTION

Increased attention has been paid to the vulnerability of the Spent Fuel Pools (SFPs) since the Fukushima accident. This vulnerability is of major concern for nuclear safety because SFPs which are large water-filled structures are generally placed outside the reactor containment building so that the fuel clad is the only barrier against fission product release in case of dewatering. To gain knowledge in the above mentioned areas and in order to better evaluate the safety margins, the French Institut de Radioprotection et de Sûreté Nucléaire (IRSN) in collaboration with partners from French universities has launched the experimental DENOPI project. Most of the kinetic investigations on the Zr-based cladding material oxidation in air at high temperature have used bare cladding samples. As far as spent fuels are concerned, the presence of corrosion oxide scale formed during normal operation in reactor has to be considered. In steam atmosphere a protective effect of the pre-oxide is generally observed. Zirconium alloys oxidation resistance is an important factor affecting their use as fuel cladding material in nuclear plants. It is generally accepted that during zirconium alloys oxidation, a dense and protective oxide layer first forms on the metal, giving rise to a reaction rate decreasing with time with a parabolic law [START_REF] Chevalier | French Activity on High Temperature Corrosion in water Vapour[END_REF]2]. However, after a certain time, a kinetic acceleration appears, corresponding to the breakaway phenomenon. This is related to the loss of protectiveness of the first formed oxide scale. Breakaway can occur several times, leading to corrosion kinetics showing a cyclic behaviour [START_REF] Chevalier | French Activity on High Temperature Corrosion in water Vapour[END_REF]2]. Post-transition phenomena are also encountered at temperatures expected in accidental situations, either in steam [3,4] or oxygen and air atmospheres [5,6]. The oxidation resistance of a zircaloy, in usual operation conditions as well as in accidental situations, is controlled by the resistance to breakaway of the zirconia scale. The protectiveness loss is attributed to formation of diffusion shortcuts, for example micro-cracks, which were evidenced both for temperatures relevant to in-service conditions [5] and to accidental situations [6]. The breakaway phenomenon can be explained by mechanical considerations [7], but the transition is a complex phenomenon involving coupled chemical, mechanical, microstructural, crystallographic aspects. It has been mentioned that the transition from a protective to a more porous oxide could be associated to a phase transitions from tetragonal (t-ZrO 2 ) to monoclinic (m-ZrO 2 ), close to the metal/oxide interface [8]. From the literature, it has been experimentally demonstrated by X-ray diffraction (XRD) [9][10][11][12][13], Raman spectroscopy [14][15][16][17] or transmission electron microscopy [18,19] that t-ZrO 2 can be detected at the metal/oxide interface in pre-transition scales while it is hardly detected in the scale during the post-transition regime. Various factors such as a deviation from stoichiometry, nanometric grain size, chemical doping, and compressive stresses can explain the stabilization of the t-ZrO 2 for temperatures below its thermodynamic stability domain. A high amount of t-ZrO 2 at the metal/oxide interface is often correlated with the presence of high local compressive stresses [10,12,14,15,17]. It has been confirmed by thermodynamic calculations that compressive stresses can stabilize the t-ZrO 2 form of zirconia [20], due to the high Pilling-Bedworth ratio of the Zr/ZrO 2 system. It is generally observed that zirconia scales exhibiting the highest compressive stresses and thus a better stability of the t-ZrO 2 compound exhibit a better corrosion resistance [10][11][12]18]. Most of the authors consider that the t-ZrO 2 to m-ZrO 2 transformation have a detrimental effect on the corrosion resistance. Nevertheless, El Kadiri and co-workers [16] suggest that this phase transformation has a beneficial effect. It is due to the lattice expansion that accompanies the transformation and it contributes to maintain a compressive stress state in the scale and thus delays the oxide breakdown, which occurs when the stress sign change from compressive to tensile firstly close to the external interface. Whatever is the exact role of the t-ZrO 2 to m-ZrO 2 transformation on the breakaway phenomenon, the local stress field that develops in oxide scales formed on Zr alloys appears to be an important factor for the understanding of their oxidation resistance. This explains why stress development in zirconia scales has been studied in the past. In case of loss of primary coolant accident in a light water reactor, the zirconium alloys fuel cladding would be oxidized in air and steam mixtures at high temperatures. From the recent literature review of Guerain and co-workers [21] it appears that no growth stress determinations have been performed on Zircaloy-4 (Zy-4) in dry air or in air/steam gas mixtures close to operating conditions at 500°C. The key parameter for the understanding the effect of water vapour on the oxidation mechanism is the determination of the stress level in the zirconia scale and in the Zy-4 substrate at high temperature. In order to determine the structure and microstructure of the oxide layers, X-ray diffraction (XRD) analyses have been performed in situ under dry air and wet air environments at 500°C on Zy-4 flat sheet samples. This temperature has been chosen in order to be close to the operation conditions in nuclear power plants. The aim of the present work is to show the influence of water vapour on the growth stresses developed at high temperature in the oxide scale as well as inside the alloy. The stress evolution during cooling to room temperature will also be determined in order to collect data on the influence of water vapour on the stress relaxation mechanism.

EXPERIMENTAL

The material used in the present work is a low tin Zircaloy-4 (Zy-4) alloy. The alloy composition is given in Table 1. The specimens, were cut by electro-erosion from a SRA (stress relieved annealed at 480°C) flat sheet. Flat specimens are 0.43 mm thick and show a total surface area of about 2.6 cm 2 . The in situ characterization of the oxide scales was carried out in a MRi high temperature chamber adapted on an Xpert'3 PANalytical Θ-Θ Bragg-Brentano diffractometer (cobalt radiation, λ α k = 0.179 nm). The diffraction patterns are registered for five days (120 h) at 500°C and at 400, 300, 200, 100 and 20°C during cooling to room temperature. Each stress determination on the alloy or on the oxide scale takes about 5 hours. One should take care that for the 40 first hours oxidation in wet air the scale thickness is too low to permit a correct stress determination (the oxide peak intensity is two low). Then, no values are reported on the graphs concerning the growth stress in the oxide during this period. Concerning stress determination in the alloy, as long as the oxide scale is thin enough, the stress determination remains possible, which is the case at 500°C after 5 days. The water vapour experiments were performed in air containing 15 vol.% H 2 O. The laboratory compressed flowing air (8 L/h flow rate) is saturated through boiling water. Then, the water bath humidifier maintained at 55 °C controls the water vapour ratio by condensation of the excess water vapour. All the connecting tubes are maintained over 65°C in order to avoid any water condensation. A schematic drawing of the rig was provided in a previous paper [START_REF] Chevalier | French Activity on High Temperature Corrosion in water Vapour[END_REF]. Dry air testing was performed in flowing air at atmospheric pressure. Calibration of the sample temperature was performed with the help of a thermocouple welded on a Zy-4 specimen and placed on the heating element in the high temperature chamber.

Stress levels in m-ZrO 2 (ICDD 37-1484) were determined by in situ X-ray diffraction at 500 °C under atmospheric pressure conditions. Stress determinations were made by using ψ goniometer technique, [23,24] with Co-K αl = 0.179 nm radiation, based on the (10-4) planes of the monoclinic m-ZrO 2 structure (84.62° 2Θ). No stress determination is possible on tetragonal zirconia with this technique. The ψ angle was varied in steps of 2 ° between -38 ° and +38 °. The mean stress σ φ in the plane of the oxide is obtained from the slope (1 + ν). σ φ /E of the linear relationship between ε and sin 2 ψ, where ε is the strain, given by ∆d/d, d being the interplanar spacing. The oxide was assumed to be isotropic with no stress gradients parallel to the surface components. The m-ZrO 2 Young's modulus, E, was taken as 203 GPa at room temperature (174 GPa at 500°C) [25] and Poisson's ratio ν as 0.32. The technique used is not able to determine the eventual stress gradient inside the oxide scale. Only a mean value is obtained from the whole surface and thickness analyzed. Stress determination in the metal is based on the (105) planes of the hexagonal Zr structure (135.6° 2Θ-Co-K αl ) (ICDD 65-3366). The ψ angle was varied in steps of 2 ° between -45 ° and +45 °. The mean stress σ φ in the plane of the oxide is obtained from the slope (1 + ν) σ φ /E of the linear relationship between ε and sin 2 ψ, where ε is the strain, given by ∆d/d, d being the interplanar spacing. The Zy-4 alloy was assumed to be isotropic with no stress gradients parallel to the surface components. Ζy-4 Young's modulus, E, was taken as 96 GPa at 500°C and Poisson's ratio ν as 0.33. At 500°C and even after 120 hours oxidation, the oxide scale is thin enough to permit the alloy analysis. On the specimen exposed to dry air, the oxide scale is about 6 µm thick and the analysed depth in zirconia is about 10 µm which permits to analyse the underlying Zy-4 all along the test.

3.RESULTS

3.1. Determination of the stress evolution in the Zy-4 at 500 °C and during cooling to room temperature in dry air.

In situ growth stress determination has been performed on a Zy-4 specimen in dry air at 500 °C for 5 days (120 hours). It appears that in dry air at 500 °C the stress in the Zy-4 alloy is very low (-20 +/-20 MPa) and does not vary for the 120 h oxidation test. During cooling the stress remains the same.

3.2. Determination of the stress evolution in the m-ZrO 2 oxide scale at 500 °C and during cooling to room temperature in dry air.

Figure 1 shows the in situ growth stress determination in the m-ZrO 2 on a Zy-4 specimen in dry air at 500 °C for 5 days (120 hours). Before 20 hours oxidation, stress determinations where only performed in the alloy because the scale is not thick enough to collect valuable data concerning the m-ZrO 2 stress state. As soon as the determination is possible in the m-ZrO 2 scale, the compressive growth stress registered is relatively high (-1800 +/-200 MPa). It remains constant all along the oxidation test. At 300 °C and at lower temperatures the compressive stress decreased to reach -850 +/-200 MPa. The stress decrease indicates that a relaxation process occurred during cooling. On the specimen face exposed to dry air, the oxide scale is about 6 µm thick.

3.3. Determination of the stress evolution in the Zy-4 at 500 °C and during cooling to room temperature in wet air (15 vol.% H 2 O ).

In situ growth stress determination has been performed on a Zy-4 specimen in flowing wet air (15 vol.% H 2 O) at 500 °C and for 5 days (120 hours). For the 40 first hours oxidation, stress determinations where only performed on the alloy because the scale is not thick enough to collect data concerning the m-ZrO 2 stress state. In wet air, results show that at 500 °C the stress in the Zy-4 alloy remains very low (30 +/-20 MPa) and does not vary for the 120 h oxidation test. No stress variation is registered during cooling from 500 °C to room temperature. The metallic behaviour is similar in dry and wet air during cooling to room temperature. Figure 2 shows the in situ growth stress determination in the m-ZrO 2 oxide scale in wet air (15 vol.% H 2 O) at 500 °C and during cooling to room temperature. After 40 hours oxidation the stress determination is possible in the m-ZrO 2 scale and the growth stress close to (-1700 +/-200 MPa). Compressive stress decreases after the initial stage oxidation to reach a value close to (-1000 +/-200 MPa). Then, it remains constant along the oxidation test. On the specimen surface exposed to wet air the oxide scale is about 2 µm thick. The scale is finer than the one obtained in dry air. Nevertheless, results show that the compressive growth stress registered in the oxide scale is lower in wet air compared to dry air environment. At 400 °C and for lower temperatures the compressive stress remains close to -1000 +/-200 MPa. This indicates that thermal stresses are not relaxed during cooling by crack formation or creep due to the specific mechanical properties of the oxide induced by the presence of water in air. It should also be noticed that the residual stress values obtained after cooling to room temperature in dry and wet air are close to each other (-1000 +/-200 MPa). This is due to the stress decrease after dry air oxidation and stress value remaining constant after wet air oxidation. Nevertheless, in situ growth stress determinations have shown that the compressive stress is much higher in dry air oxidation conditions at 500 °C.

DISCUSSION

In dry air, the present results show that the (-1800 +/-200 MPa) compressive growth stress in the oxide is close to the average residual stress magnitude found in the oxide scale formed at 500°C on Zr and other alloys (M5, Zy-2) [26,27]. In the literature it is shown that the stress magnitude reaches a maximum value for a scale thickness ranging between4 and 6 µm on M5 at 550 °C [28]. In our study the growth stress determined by XRD remains constant all along the 120h isothermal test because the scale thickness remains below 6µm and the analytical technique gives a mean stress value for the whole oxide scale. In situ growth stress determination in the oxide scale indicates that the compressive stress is larger in dry air compared to wet air environment at 500 °C (-1000 +/-200 MPa). The lower stress state observed in wet environment is in accordance with the results of Pétigny [10,11]. A comparative study on the effect of oxygen and steam was performed by this author at 470°C by using X-ray diffraction. It appears that the average stress magnitudes are larger for oxygen compared to steam. These results are consistent with the values obtained by synchrotron X-ray diffraction by Béchade et al. at 400°C [29,30].As observed in the present work, it has been proposed that under wet atmospheres, the average stress values are at a maximum at the very early oxidation stage and then decreases. This behaviour may be related to stress relaxation phenomena taking place early during the formation of the scale [21]. This could explain the relatively low stress values obtained in wet air in the present work even though the stress determination could not be performed at the very beginning of the test (before 40 hours) due to the too low scale thickness. In dry air, the compressive growth stress registered in the m-ZrO 2 scale at 500 °C is relatively high (-1800 +/-200 MPa) (Figure 1). At 300 °C and at lower temperatures the compressive stress decreases to reach -900 +/-200 MPa at room temperature. The progressive stress decrease indicates that a relaxation process occurred during cooling due to cracks initiation and propagation in the zirconia scale during cooling. In wet air, the compressive growth stress registered at 500 °C is relatively low (-1000 +/-200 MPa) in the m-ZrO 2 scale (Figure 2). At 400 °C and for lower temperatures the compressive stress remains the same (-1000 +/-200 MPa). It indicates that thermal stresses are not relaxed during cooling by crack formation or creep due to the specific mechanical properties of the oxide induced by the presence of water in air. It has been observed that the residual stress values obtained after cooling to room temperature in dry and wet air are close to each other (-1000 +/-200 MPa). This is due to the stress relaxation after dry air oxidation and the constant stress value registered after wet air oxidation. This indicates that no important stress relaxation by scale cracking occurred in zirconia scales obtained under wet air conditions during cooling, because cracking has already occurred under isothermal conditions after the early stage oxidation period (about 60 hours). It has also been proposed that the incorporation of hydrogen can occur in the metal in the presence of steam [30]. Hydrogen is located in interstitial positions in the Zr hexagonal lattice; or if the solubility limit is exceeded, it forms Zr hydride precipitates. In both cases it leads to an increase of the volume of the metal which may modify the stress evolution [31,32]. However, a comparison of the results obtained under dry and wet conditions show that no detectable influence of hydrogen has been registered in the mean stress state of the alloy even if hydrogen dissolution is expected during wet air oxidation.

CONCLUSION

In situ growth stress determination in the oxide scale has shown that the compressive stress is larger in dry air oxidation conditions compared to wet air environment at 500 °C. The lower stress state observed in wet environment is in accordance with results performed in the same temperature range on other zirconium alloys. The lower growth stress can be explained by growth stress relaxation due to cracking of the oxide in wet air at 500 °C. This will be confirmed by acoustic emission experiments. The oxide creep can also be favoured in wet air due to proton doping of zirconia or hydrogen diffusion in the alloy. Analysis of the hydrogen content after wet air oxidation will be performed in the near future. After dry air oxidation, the compressive stress decreased during cooling from -1800 +/-200 MPa to reach -900 +/-200 MPa at room temperature. It indicates that a relaxation process occurred during cooling due to cracks initiation and propagation in the zirconia scale. After wet air oxidation, the compressive stress remains close to -1000 +/-200 MPa even after cooling to room temperature. This indicates that the growth stress state in the oxide scale is not high enough to induce detectable stress relaxation during cooling due to the specific mechanical properties of the oxide induced by the presence of water in air.

Figure 1 .

 1 Figure 1. In situ growth stress determination in the m-ZrO 2 scale in dry air at 500 °C and during cooling to room temperature.

Figure 2 .

 2 Figure 2.In situ growth stress determination in the m-ZrO 2 oxide scale in wet air (15 vol.% H 2 O) at 500 °C and during cooling to room temperature.

Table 1 :

 1 Composition of the alloy used for stress determination.

	Alloy	Composition (wt%)				ppm	ppm	ppm
		Zr	Sn	Fe	Cr	O	C	Nb	H
	Zy-4	Balance	1.32	0.21	0.11	0.12	125	< 40	< 3

Acknowledgement. The DENOPI project is a work performed in the frame of the DENOPI project, funded by the French government as part of the "Investment for the Future" Program reference ANR-11-RSNR-0006.


Henri Buscail
email: henri.buscail@udamail.fr

Christophe Issartel
email: christophe.issartel@udamail.fr

Frédéric Riffard
email: frederic.riffard@udamail.fr

Raphaël Rolland

Sébastien Perrier
email: sebastien.perrier@udamail.fr

Alexandre Fleurentin
email: alexandre.fleurentin@cetim.fr

Hot Corrosion and Degradation in Corrosive Atmospheres Focus

Keywords: Lanthanum, Sol-gel coating, Chromia, Alloy 330 (Fe-35Ni-18Cr-2Si)

The influence of a lanthanum sol-gel coating on the chromia scale adherence has been studied on the 330 alloy (Fe-35Ni-18Cr-2Si) oxidized at 900°C, in air. Argon annealing of lanthanum sol-gel coatings have been performed at various temperatures. Kinetic results show that lanthanum sol-gel coatings lead to a lower oxidation rate compared to blank specimens. On blank 330 specimens scale spallation is observed after cooling to room temperature. On the non-annealed sol-gel coated specimen and the argon annealed specimens at 600, 800 or 1000°C, the oxide scale formed at 900 °C is adherent after 48 hours isothermal oxidation. The adherent oxide scales are convoluted. It results from o an anionic and a cationic mixed diffusion process in the chromia scale Thermal cycling tests on lanthanum the sol-gel coated specimen show that the oxide scale remains adherent after 250 cycles. It is concluded that argon annealing of the lanthanum sol gel coating is not necessary to improve the scale adherence.

INTRODUCTION

Metal dusting is a catastrophic carburisation phenomenon that occurs at temperatures of 450-850°C in atmospheres of high carbon activity. Even though some high-silicon alloys resist to metal-dusting environments [1], carburization is the main reason of conveyor units degradation due to a quick alloy weakening. The resistance of alloys to corrosion, including metal dusting, relies on the formation of a dense, adherent oxide layer that separates the alloy from the corrosive environment. For such an oxide layer to be protective, it must achieve full surface coverage, be crack-free and be established before significant material degradation has occurred. An adherent oxide scale can act as a carbon diffusion barrier and improve the alloy carburization resistance [2][3][4][5]. Chromia Cr 2 O 3 is considered as a protective oxide formed at high temperature. Stevens et al. have shown that a Fe-35Ni-18Cr-2Si alloy can be used as a resistance heating-element material and it has been studied in the 1000 and 1300°C temperature range [6]. Nevertheless, at temperatures higher than 1100°C the chromium oxide CrO 3 evaporation occurs. In this case, chromium depletion is observed in the oxide scale and in the alloy. The concentration of elements in the protective oxide was found to be both temperature and time dependent. Chromia is stable at low temperature (< 1100 °C) and was found to evaporate at higher temperatures and longer oxidation time, thereby increasing nickel and iron concentrations. On the basis of changes in the chemical composition of the oxide film, together with scaling, the failure of heating-element wire by the development of hot spots can be explained. At failure both long-term (low temperature < 1000°C) and shorter-term tests (temperature > 1000°C) were found to have oxidized to the same depth. This can be explained on the basis of a depletion of chromium, which occurred at the oxide-alloy interface, and becomes insufficient for a healing process. A previous study has shown that in the 800-1000°C temperature range, the niobium containing 330Cb alloy oxidation leads to a chromia scale acting as a good diffusion barrier under isothermal conditions [7]. The external manganese chromite subscale has no influence on the diffusion properties of the main chromia scale. It is expected that the manganese chromite can limit the Manuscript Click here to view linked References chromia scale evaporation at temperatures higher than 1000°C [8]. Nevertheless, after oxidation and cooling to room temperature, important scale spallation generally occurs [7]. At 900°C, spallation only appeared on the niobium and silicon rich areas. At 1000°C, scale spallation is present on the entire specimen surface indicating that this temperature is too high to obtain an adherent chromia scale on the alloy surface. The comparisons between the 48 and 160h oxidation duration have shown that at 800 and 900°C a short oxidation duration (46h) is favourable to the oxide scale adherence. It then appeared that the chromia scale never shows a perfect adherence on the alloy under the tested conditions. In order to improve the chromia scale adherence previous works have proposed to introduce lanthanum on the alloy surface as a La 2 O 3 coating [9][10][11][12][13] or as an alloying element [14]. The present work deals with the effect of lanthanum sol-gel coatings applied on the specimen surface before oxidation. Previous works have already indicated that the argon annealing could improve the effect of the coating on the oxide scale adherence [15][16][17]. These works where performed on alumina forming alloys. Lanthanum sol gel coating alone was not improving the alumina scale adherence. The aim of the present work is to check if the argon annealing could improve the effect of the coating on chromia forming alloy. Then, the sol-gel coating argon annealing as been tested at various temperatures on a model 330 chromia forming alloy.

EXPERIMENTAL PROCEDURES

The model 330 alloy is an austenitic stainless steel elaborated at the Ecole des Mines de Saint Etienne, France. Table 1 gives the 330-alloy composition in weight %. 1 mm thick cylindrical specimens of around 13 mm diameter were abraded up to the 120-grit SiC paper, then degreased with ethanol and finally dried. This surface finishing was chosen in order to fit the surface state of the alloy used in industrial applications. The oxidation tests were performed at 900°C. Kinetic results are obtained by thermogravimetric measurements under isothermal conditions. High temperature oxidation was performed during 46 hours in ambient air at the atmospheric pressure using a Setaram TGDTA 92-1600 microthermobalance. The oxide characterization was realized by in situ X-ray diffraction (XRD) tests. In situ XRD patterns were obtained by use of a high temperature MRI chamber installed in a Philips X'pert MPD diffractometer (copper radiation, k  = 0.15406 nm). The XRD conditions were 2 scan, step 0.05° ranging from 10 to 80°, 2.5 s counting time. Each in situ XRD pattern is obtained within one hour in the 10-80° 2 range. In situ XRD permit to follow the evolution of the oxide nature with time at the testing temperature. The oxide scale surface and cross-section morphologies have been observed in a JEOL 6400 scanning electron microscope (SEM) coupled with a LINK energy dispersive X-ray spectroscopy (EDXS). The EDXS point analyses were performed with an electron probe focused to a 1m spot. The lanthanum sol-gel coating was performed using a 0.4 mol.l -1 lanthanum hydroxide solution obtained from the precipitation of lanthanum hydroxide in 0.27 mol 28% ammonia. After filtration, the dissolution of the precipitate is obtained in the 69% nitric acid at 60°C [18]. The lanthanum sol-gel coatings are carried out by dipping the specimens during 2 s in the solution. Immediately after dipping, the specimens are dried in warm flowing air at 40°C during 1 min. The coating thickness is close to 0.5 µm. After solgel coating, the argon annealing was performed in a horizontal quartz reaction tube. A heating ramp is applied at a rate of 5°C/min. until the 600, 800 or 1000°C annealing temperature is reached. The temperature is maintained during 2 h. Then, the specimen is cooled down at 5°C/min to room temperature. Before oxidation the XRD pattern indicates that the alloy is detected alone on the surface on non-annealed specimens. It means that the sol-gel coating is not sufficiently crystallised before the heating up of the specimen to be detected by XRD. Thermal cycling was performed in static laboratory air at atmospheric pressure in a horizontal tubular furnace. Two samples were placed at the same time into an alumina crucible to estimate the reproducibility. The thermal cycle consists in a 30 minutes exposure at high temperature, followed by 30 minutes at room temperature after air quenching. This frequency corresponds to industrial conditions in conveyor units. The weight change of the samples was determined after each cycle on a balance with an accuracy of ± 0.01 mg.

RESULTS

Kinetic results indicate that all lanthanum sol-gel coatings show a lower oxidation rate compared with blank specimens in Figure 1. The lowest oxidation rate is observed with non-annealed specimen, oxidized at 900 °C. The other sol-gel coated specimens show similar kinetic behaviours. Figure 2 shows the mass change curves versus the number of 30 minutes cycles at 900°C for blank specimens and lanthanum sol-gel coated specimens. The comparison of the kinetic curves show that lanthanum sol-gel coatings largely improve the oxide scale adherence. On blank specimens spallation started after 50 cycles. On lanthanum sol-gel coated specimens the oxide scale is still adherent after 250 cycles.

In situ X-ray diffraction patterns obtained on the blank 330 specimen oxidized, in air at 900°C indicate that the oxide scale is only composed of Mn 1.5 Cr 1.5 O 4 (JCPS 33-0892) and Cr 2 O 3 (JCPDS 38-1479). Chromia is formed at the beginning of the oxidation process. No phase transition is detected during the 46h oxidation and during cooling to room temperature. The relatively high Mn 1.5 Cr 1.5 O 4 peaks intensities indicate that it is present at the external interface even though it corresponds to a low amount of oxide grains observed on the SEM micrographs. The alloy peaks are still detected after 46h oxidation because the scale is thin enough to permit its detection. Figure 3 shows the in situ XRD patterns obtained during the 46 first hours oxidation at 900°C, in air, of the non-annealed sol-gel coated alloy 330. Before oxidation, XRD shows that the alloy is detected alone on the surface. It indicates that the sol-gel coating is not crystallised before the heating up of the specimen. At 900°C, XRD analyses show the initial formation of Mn 1.5 Cr 1.5 O 4 . The oxide Cr 2 O 3 appears rapidly at the beginning of the oxidation. A lanthanum containing oxide LaCrO 3 (JCPDS 33-0701) is also observed from the beginning of the oxidation process. This phase is stable all along the test and is not transformed into another oxide. No phase transition is detected after cooling to room temperature.

In situ XRD patterns, obtained during the 46 first hours oxidation at 900°C, in air, of the sol-gel coated 330, with a 600°C argon annealing, show that LaCrO 3 is present on the surface. This oxide is present on the surface until the end of the oxidation test. At 900 °C, XRD patterns show the growth of Mn 1.5 Cr 1.5 O 4 and Cr 2 O 3 . Similar results are obtained on 800 or 1000 °C argon annealed specimens. After argon annealing chromia and LaCrO 3 are present on initial surface. Cr 2 O 3 and the manganese chromite grow all along the oxidation test. LaCrO 3 is less detected with time because it is incorporated inside the growing chromia scale. Figure 4 shows the specimen's surface scale morphologies (SEI images x 3000). Blank specimens oxidized during 46h are covered by a chromia scale and octahedral manganese chromite grains. Spalled areas are found on about 10% of the blank specimens surface. Where the oxide scale spalled off, the underlying substrate surface is silicon rich (Figure 4a). With sol-gel coated specimens, no scale spallation is observed on the surface (Figure 4b). The oxide scale is convoluted and the grain size is small. The oxide grains are composed of chromia and lanthanum chromite. The surface micrograph obtained after the 46 hours oxidation at 900°C in air of the sol-gel coated alloy 330, argon annealed at 600°C is presented on Figure 4c. The biggest pyramidal grains correspond to manganese chromite. No spallation is observed on the specimen's surface. The convoluted scale is mainly composed of chromia. SEM, surface micrographs obtained after the 46 hours oxidation at 900°C in air of the sol-gel coated alloy 330, argon annealed at 800°C, are presented on Figure 4d. No spallation is observed on the surface after oxidation at 900°C. The chromia scale is convoluted. Surface micrographs obtained after the 46 hours oxidation at 900°C, in air, of the sol-gel coated 330, argon annealed at 1000°C are presented on Figure 4e. In this case, little spalled areas are found on the surface. Lanthanum is still found on the specimen's surface. It is also observed that the scale is convoluted. It is then concluded that the 1000°C temperature argon annealing permits the incorporation of lanthanum inside the oxide scale.

The scale cross-section morphologies are presented on Figure 5. Blank specimens, oxidized during 46h, are covered by a 3m thick chromia scale (Figure 5a). On the sol-gel coated 330 specimen no scale spallation is observed on the cross-section (Figure 5b). The oxide scale is 2m thick and convoluted. LaCrO 3 is located at the external interface (white areas on the BSE micrograph). The cross-section micrograph obtained after the 46 hours oxidation at 900°C in air of the sol-gel coated 330, argon annealed at 600°C is presented on Figure 5c. No spallation is observed on the specimen's cross-section. The very convoluted scale is mainly composed of a 2m thick chromia scale. Figure 5c shows the presence of lanthanum (white areas) inside the chromia scale. The cross-section micrograph obtained after the 46 hours oxidation at 900°C in air of the sol-gel coated alloy 330, argon annealed at 800°C is presented on Figure 5d. Lanthanum is found inside the thin chromia scale. The 800°C temperature argon annealing permits the incorporation of lanthanum inside the oxide scale. Cross-section micrographs obtained after the 46 hours oxidation at 900°C in air of the sol-gel coated 330, argon annealed at 1000°C is shown on Figure 5e. On this specimen, lanthanum is present close to the oxide/gas interface of the chromia scale.

DISCUSSION

Kinetic results have shown that lanthanum sol-gel coatings always lead to a lower oxidation rate compared to blank specimens. All heat treatments lead to the same kinetic effect because lanthanum has been incorporated into the oxide scale in the same way leading to similar scale structures. Paúl already showed this effect on austenitic steels. This author proposed that a lanthanum sol-gel coating lowers the parabolic rate constants during oxidation at 900°C [19,20]. It was also demonstrated that cerium ion implantation followed by an annealing process reducing the defaults present in the alloy structure, leads to a lower oxidation rate at 900°C [19]. On the 330 alloy the oxide scale protective character is confirmed by in situ XRD analysis (Figure 3), which shows the chromia scale formation.

It also exhibits the manganese chromite growth from the beginning of the oxidation process at 900°C. SEM surface analyses confirm the presence of Cr 2 O 3 and Mn 1.5 Cr 1.5 O 4 . On all sol-gel coated specimens, the chromia scale growth occurs with the presence of LaCrO 3 from the starting point of the XRD in situ test. This oxide formation is thermodynamically explained by a reaction between Cr 2 O 3 and La 2 O 3 according to the Cr 2 O 3 -La 2 O 3 phase diagram [22]. This oxide has also been obtained by mechanochemical synthesis by grinding constituent oxides by Zhang [23]. The LaCrO 3 formation at high temperature on chromia forming alloys has been described by other authors [13,[24][25][26][27]. The present results indicate that the reaction between Cr 2 O 3 and La 2 O 3 occurs at the beginning of the oxidation process on all coated specimens. As observed by SEM, argon annealing also lead to the lanthanum incorporation inside the chromia scale during oxidation. In situ XRD shows that no structural changes occur along the test and during cooling to room temperature. It is also demonstrated that the lanthanum chromite LaCrO 3 relative peaks intensity is the same before an after cooling, indicating that this oxide remains unchanged after cooling. On figure 5, it is observed that LaCrO 3 is mainly located close to the oxide/gas interface and that inward oxygen diffusion contributes to the chromia scale growth. This scale structure develops with time during the oxidation process and is not the result of a deformation during cooling to room temperature. The convolution mechanism formation corresponds to the relaxation process of the lateral growth stresses generated in the chromia scale during the oxidation process. It results from the formation of new oxide inside the chromia scale. Moreover, Figure 2 shows that on blank specimens spallation started after 50 cycles whereas on lanthanum sol-gel coated specimens the oxide scale is still adherent after 250 cycles. As proposed by other authors a convoluted scale structure suggests that the cationic diffusion in chromia scale is lowered by the presence of lanthanum [28]. With lanthanum, the oxide scale is adherent and shows a very convoluted aspect. This scale structure develops with time during the oxidation process and is not the result of a deformation during cooling to room temperature [29][30][31][32][33]. The convolution mechanism formation is not fully understood but it is proposed that during the oxidation process, lateral growth occurred in the chromia scale [33]. It results from the formation of new oxide inside the chromia scale [33]. It can also be due to growth stresses generated inside the chromia scale leading to a plastic deformation of the scale and the underlying metallic substrate. Then, the growth stress relaxation in a thinner scale can explain the better scale adherence. Hou et Al. proposed a mechanism explaining a scale convolution on a lanthanum implanted Fe-Cr-Al alloy [34]. They have shown the distribution of local stress components. It appears that the convolution results from ox  which is the compressive stress inside the oxide scale and n  which is the tensile stress normal to the alloy/oxide interface on the convoluted oxide scale. According to Evans et al [35], n  increases with the curvature radius increase. Pint indicated that during the scale growth, stresses develop inside the scale and at the oxide/alloy interface [28]. If the scale adherence is maintained, growth stresses can be high enough to plastically deform the metallic substrate and also influence the scale morphology. Interfacial voids can also play an important role on the stress development and scale morphology. With, compressive stresses in the scale, an interfacial void can lead to a stress contribution perpendicular to the alloy/oxide interface. On a hard metallic substrate, cavities can develop without scale buckling during its thickening. On the contrary, a smooth alloy can accommodate the stress increase by plastic deformation and buckling of the oxide scale. Some authors have shown that an yttrium sol-gel coating promotes a continuous silica scale formation at the internal interface on the alloy AISI 304 [36]. The silicon presence hinders the iron oxidation and the formation of non-protective iron oxides [37]. Even though alloy 330 contains 2.15 wt.% silicon no continuous silica scale formation was observed within the 46h oxidation at 900°C, in air because the oxidation duration is perhaps to short. The bad scale adherence on uncoated specimens can be also explained by the presence of a too high silicon amount in the alloy [34]. The absence of iron oxides in the scale can also be due to the fact that the alloy 330 is a 33.8% nickel-rich alloy.

CONCLUSION

Kinetic results have shown that lanthanum sol-gel coatings always lead to a lower oxidation rate compared to blank specimens. In situ XRD has shown that the lanthanum coating leads to the formation of LaCrO 3 by a reaction between La 2 O 3 and Cr 2 O 3 . The scale structure suggests that the cationic diffusion in a chromia scales is lowered by the presence of lanthanum. On blank 330 specimens scale spallation is observed after cooling to room temperature. It appears that the 2-wt% silicon present in the alloy has no beneficial effect on the chromia scale adherence and that no continuous silica layer is formed. On the non-annealed sol-gel coated specimen and on the argon annealed specimens at 600, 800 and 1000°C, the oxide scale formed at 900°C, in air, is adherent after 48 hours isothermal oxidation. SEM micrographs revealed that the adherent oxide scales are convoluted and that lanthanum is close to the external interface. It results from the formation of new oxide inside the chromia scale due to an anionic and a cationic mixed diffusion process. Convolution can also be the result of growth stresses relaxation by plastic deformation of the scale. It appears that, argon annealing does not provide a significant improvement compared to non-annealed lanthanum coated specimens. Thermal cycling tests, performed on blank 330 and lanthanum sol-gel coated specimens without argon annealing, show that the scale adherence is largely improved by this coating. The results show that argon annealing of the lanthanum sol gel coating is not necessary to improve the scale adherence.

Figure captions:

FIGURE 1 .

 1 FIGURE 1. Mass gain curves obtained on the alloy 330, oxidised in air, at 900°C. Comparison of the kinetic curves obtained on the specimens: Blank specimens, lanthanum sol-gel coated without annealing, and lanthanum sol-gel coated with argon annealing at 600, 800 or 1000°C.

FIGURE 2 .

 2 FIGURE 2. Mass change curves versus the number of cycles (0.5h cycles), at 900°C; on blank and lanthanum solgel coated 330 alloy specimens without argon annealing specimens.

FIGURE 3 .

 3 FIGURE 3. In situ XRD patterns obtained for the 46 first hours oxidation at 900°C in air of the lanthanum sol-gel coated alloy 330.

FIGURE 4 .

 4 FIGURE 4. SEM surface micrographs of the specimens oxidized at 900°C, for 46 h, in air (SEI x 3000), effect of Lanthanum coatings. a) Blank alloy 330, b) Alloy 330 sol-gel coated without argon annealing, c) Alloy 330 sol-gel coated with 600°C argon annealing, d) Alloy 330 sol-gel coated with 800°C argon annealing, e) Alloy 330 sol-gel coated with 1000°C argon annealing.

FIGURE 5 .Figure 1 :

 51 FIGURE 5. SEM cross-section micrographs of the specimens oxidized at 900°C, for 46 h, in air (BSE). a) Blank alloy 330, b) Alloy 330 sol-gel coated without argon annealing, c) Alloy 330 sol-gel coated with 600°C argon annealing, d) Alloy 330 sol-gel coated with 800°C argon annealing, e) Alloy 330 sol-gel coated with 1000°C argon annealing.

Figure 1 -Figure 2 :

 12 Figure 1

Figure 2

 2 Figure 2

Figure 4 :

 4 Figure 4: In situ XRD patterns obtained for the 46 first hours oxidation at 900 °C in air of the lanthanum sol gel coated alloy 330.

Figure 8 :

 8 Figure 8: SEM surface micrographs of the specimens oxidized at 900°C, for 46 h, in air (SEI x 3000), effect of Lanthanum coatings. a) Blank alloy 330, b) Alloy 330 sol gel coated without argon annealing, c) Alloy 330 sol gel coated with 600 °C argon annealing, d) Alloy 330 sol gel coated with 800 °C argon annealing, e) Alloy 330 sol gel coated with 1000 °C argon annealing.

Figure 9 :

 9 Figure 9: SEM cross-section micrographs of the specimens oxidized at 900°C, for 46 h, in air (BSE). a) Blank alloy 330, b) Alloy 330 sol gel coated without argon annealing, c) Alloy 330 sol gel coated with 600 °C argon annealing, d) Alloy 330 sol gel coated with 800 °C argon annealing, e) Alloy 330 sol gel coated with 1000 °C argon annealing.

Table 1

 1 330 alloy composition (Fe-34Ni-21Cr-2Si) weight %.

	Alloy 330	Fe	Ni	Cr	Si	Mn	Mo	Cu	C
	wt. %	Bal. 33.80 21.10 2.15 0.70 0.20 0.20 0.05

Table 1

 1