N Bourenane

Y Hamlaoui
email: hamlaoui_youcef@yahoo.fr

C Remazeilles

F Pedraza

Effect of the pH of the electrolyte on the formation and on the corrosion properties of ceria based coating on carbon steel

Keywords: Cerium oxide coatings, Cathodic electrodeposition, Solution pH, Morphology, Corrosion products, Corrosion

The effects of the modification of the pH in 0.1M cerium nitrate solutions in the elaboration and corrosion resistance of ceria based coatings on carbon steel are investigated. Increasing the concentration of acetic acid impedes an efficient electrodeposition. At low concentrations, acetic acid seems to prevent the precipitation of Ce(OH)3 and the formation of even films of cerium oxides is favoured. The increase of pH through the addition of NaOH to the cerium nitrate solutions with 0.008M of acetic acid is shown to provide superior corrosion resistance for exposures in air and immersed in 0.5M NaCl for 30 days.

Introduction

Rare earth metals are among the best alternatives to the toxic hexavalent chromium either as conversion coatings or corrosion inhibitors. Hinton et al. [1] were the first authors to carry out the deposition of rare earth metals. Various coating techniques can be adopted regarding the application of these rare earth metals based coatings on different metal substrates (Al, stainless steels, zinc, etc.) to be protected. For instance, physical, chemical or electrochemical techniques like simple immersion, anodic or cathodic deposition, PVD, sol-gel, etc. have been reported to successfully elaborate such coatings [2][3][4][5][6][7][8][9][10][11][12].

The final coating is composed of two layers (inner and outer) with different chemistries. While it is established that the outer layer is mainly composed of cerium oxy-hydroxides [13], the chemistry of the inner layer depends on the metal substrate and on the application technique [12][13][14]. Indeed, in addition to the cerium (III) hydroxide, the inner layer of the coating is usually enriched with different amounts of the chemical dissolution products of the substrate due to the acidic pH of the deposition solution. For some materials, the presence of this type of corrosion products does not affect the quality of the coating. For instance, it is well known that the dissolution products of e.g. Al, and Cr are homogeneous, compact and adhere well to the surface of the substrate and consequently constitute a corrosion barrier to prevent aggressive media to penetrate through. In contrast, the corrosion products formed on zinc impair little protection [15,16].

In the case of mild steel, the Ce(NO3)3 concentrations employed to deposit cerium oxyhydroxides have pH values between 3.1 to 4.5 [17]. Such acidic pHs promote the formation of ferrous hydroxide Fe(OH)2 and carbonated green rusts that oxidize further in lepidocrocite [18]. Most of such iron-based corrosion products are known to be not protective due to their porosity and low adhesion to the steel substrate [19]. Yet, cerium oxides films have been reportedto act as effective cathodic coating and to restore the passive state of steel after being disturbed in NaCl 3.5% solution "self-healing mechanism" through the formation of secondary layer of of Ce(OH)4, CeO2, Fe(OH)2 and Fe2O3 in scratched areas [20].

Therefore, the use of additives in the electrodeposition bath is often recommended. Generally, small amounts of additives affect mainly the electrodeposition reaction kinetics either by adsorption or complexation [27][28]. Thus, an important part of metal and metal oxide electrodeposition is realized from baths containing organic additives [29][30][31][32][33]. In the particular case of an efficient cathodic electrodeposition of the cerium oxides, it is necessary to stabilize Ce 3+ ion to prevent the precipitation of Ce(OH)3 according to reactions (1) to (3) [34,35]:

2 Ce 3+ + 3 H2O ⇄ Ce(OH)3 + 6 H + (1) 2 Ce 3+ + 3 H2O ⇄ Ce2O3 + 6 H + (2) Ce2O3 + H2O ⇄ 2 CeO2 + 2 H + + 2e - (3)
In some of our previous works [36,37] we have shown that the coatings obtained from a bath containing cerium nitrate and an organic additive (polyethylene glycol, PEG) on electrogalvanized steel and in mild steel were somewhat free from corrosion products and showed a good resistance against corrosion for long immersion times in NaCl and Na2SO4 solutions. Ferreira et al. [38] obtained a uniform protective coating composed of CeO2 dominated by Ce2O3 on electrogalvanized steel through the addition of citric acid. Golden and Wang employed lactic, acetic, citric and oxalic acids and EDTA in the anodic deposition method of CeO2 [39]. They concluded that ligands with weaker formation constants, i.e. acetic and lactic acid were able to produce CeO2 films under certain experimental conditions. Indeed, ligands (chelants) such as EDTA and citric acid can strongly complex the metal [40] in which case, cerium would not be available to precipitate as cerium hydroxide. Therefore, in the present study, acetic acid is chosen since it forms metalorganic complex with Ce 3+ and dissolve well in aqueous medium [39]. The aim is to obtain electrodeposits on carbon steel substrate free from corrosion products with good corrosion protectiveness. Thus, cathodic electrodeposition of cerium oxide was carried out on carbon steel from a bath containing 0.1 M cerium nitrate and acetic acid. The effect of the pH of the solution was also studied by adding drops NaOH. The corrosion protection performance of the final electrodeposited films was evaluated in ambient air and in 0.5M NaCl solution for 30 days.

Experimental procedure

Materials

Round samples of 14 mm of diameter and thickness of 2 mm of A 366 cold-rolled carbon steel (Fe-0.13C-0.041Mn-0.04S-0.012N-0.55Cu, wt%-nominal composition) were polished with progressively finer grit of SiC till grade#4000, rinsed with distilled water and cleaned in ultrasonic bath of ethanol and dried with hot air immediately before electrodeposition.

Experimental set-up for cathodic electrodeposition and electrochemical tests

A 0.1M aqueous solution of Ce(NO3)3,6H2O (Aldrich,  99% pure) was employed as electrodeposition bath (pH 3.87). Different concentrations (0.008 to 0.1 M) of acetic acid were added as complexing agent. The pH of the solutions was modified by adding drops of NaOH every 5 min. The electrodeposition was performed using a classical three-electrode configuration, with the carbon steel sample being the cathode, a platinum grid as the counter electrode and a Saturated Calomel Electrode (SCE) as the reference one. The measurements were carried out using a potentiostat/galvanostat AutoLab PGZ 300. The same experimental set-up was used for the corrosion tests in 0.5M NaCl (Fluka,  99.5% pure). Cyclic voltammetry was recorded from -1.0 to -2.5 V/SCE at 10 mV/min as scanning rate. All the corrosion tests were repeated two or three times for reproducibility purposes. They were carried out at room temperature by magnetic stirring the solutions to obtain a slight vortex of the electrolyte. Prior to any electrochemical test, the time to stabilization of the open circuit potential was 30 min.

The Tafel polarization curves were obtained at a scanning rate of 10 mV/min around ±250 mV

with respect the open circuit potential (Eocp). Polarization resistance (Rp) was carried out at 10 mV/min of scan rate, and the derived values were obtained at 20 mV away (cathodic and anodic domains) of the corrosion potential (Ecorr).

Characterization of the films

The morphologies of the films were first investigated by optical microscopy (LEICA DM R-MN) then by Scanning Electron Microscopy (SEM) (JEOL 5410 LV). The structural features of the deposits were investigated by X-Ray diffraction (XRD) in the θ-2θ configuration using Cu Kα radiation =1.5406 nm (Bruker AXS D8-Advanced diffractometer). The different components were analyzed by Raman microspectroscopy (Jobin Yvon LabRam HR8000) using an incident beam of 632.82 nm emitted by a HeNe laser. Different spots ~3 μm on the surface were analyzed after focusing with the ×50 lens of the optical microscope (Olympus BX 41) attached to the apparatus. The resolution of the spectra is about 2 cm -1 at room temperature.

The characteristic bonds were identified by Fourier Transformed Infrared spectroscopy with a Thermo Nicolet FT-IR Nexus spectrometer using a KBr beamsplitter, a DTGS detector and a diffuse reflectance accessory. The powders of the deposit were mixed by grinding with a mirror.

The spectra were recorded with the Omnic software at a resolution of 4 cm -1 and an accumulation of 64 scans. The background was carried out with KBr. The deposits were also scraped then milled and further analyzed by differential scanning calorimetry (DSC) in a Q100 of TA instruments between 30 and 550°C with a heating rate of 10°C min -1 under nitrogen.

Experimental set-up for electrochemical tests

For the electrochemical tests, the same experimental set-up used in the electrodeposition tests was employed.

Results and Discussion

Characterization of the deposits

The influence of the addition of different concentrations of acetic acid on the deposition of cerium hydroxide was investigated. With the addition of 0.1M of acetic acid, the pH of the cerium nitrate solution dropped from 3.87 to 2.73. After several tests of electrodeposition at -0.5 mA/cm² for 20 min (1200 s), the pH remained unchanged (2.73), which suggested that no film had formed. This was confirmed by XRD, where no peaks related to CeO2 or Ce(OH)3 were detected (Fig. 1). Therefore, the concentration of acetic acid was decreased to (C1) 0.008 (pH 3.17), (C2) 0.01 (pH 3.15), (C3) 0.03 (pH 3.00) and (C4) 0.06 M (pH 2.87) (Table 1).

According to the literature [38,39], the reaction between the cerium salt and acetic acid occurs through equation <4> :

Ce +3 + xL - Ce𝐿 𝑥 (3-𝑥)+ (4),
where L is here the acetate ion (CH3COO -) This reaction is characterized by a single constant of formation with Log β = 1.68 [39] following equation ( 5):

Ce 3+ + 3 CH3COOH ⇄ Ce(CHCOO)3 + 3 H + (5),
Hence, for any concentration of acetic acid, e.g. 0.008M:

𝐶𝑒 3+ + 3𝐶𝐻 3 𝐶𝑂𝑂𝐻 → [𝐶𝑒(𝐶𝐻 3 𝐶𝑂𝑂) 3 ] + 3𝐻 + at t=0 0.1M 0.008M 0 0 at equilibrium 0.1-x 0.008-3x x 3x
Therefore,

𝐾 𝑓 = [𝐶𝑒(𝐶𝐻 3 𝐶𝑂𝑂) 3 ][𝐻 + ] 3 [𝐶𝑒 3+ ][𝐶𝐻 3 𝐶𝑂𝑂𝐻] 3  [𝐶𝑒(𝐶𝐻 3 𝐶𝑂𝑂) 3 ] = 𝐾𝑓 (0.1-𝑥) ([𝐶𝑒(𝐶𝐻 3 𝐶𝑂𝑂) 3 ]-3𝑥) 3 (3𝑥) 3
with 𝑙𝑜𝑔𝐾 𝑓 = 1.68 [39] ⟹ 𝐾 𝑓 = 10 1).

Since cathodic electrodeposition occurs through the local alkalinisation of the interface [9], the lowest concentration of acetic acid (0.008M) allowed to maintain a relatively high pH (3.17) that facilitated the homogeneous deposition. In contrast, the increase of the acetic acid concentration is detrimental to the cerium coating as observed through the absence of decrease of potential toward more negative (cathodic) values (Figure 2a). This behavior was confirmed by Raman spectroscopy (Figure 2b). Irrespective of the concentration of acetic acid, all the Raman spectra showed the symmetric vibration of Ce-O bond at 457 cm -1 whose intensity decreased with increasing acid concentration. In addition, the shift of this bond from 457 to 450 cm -1 with low acid concentration could be related to the evolution of the particle size of CeO2 and/or to oxygen vacancies in the deposits [41]. This hypothesis is supported by the band at 600 cm -1 typical of oxygen vacancies [35]. The peak at 712 cm -1 is related in the literature to carbonates [17] while the peaks at 740 and 1049 cm -1 can be ascribed to nitrates [18].

Therefore, the bath containing 0.1M of Ce(NO3)3,6H2O and low concentration of acetic acid (C1 = 0.008M, pH 3.17) was selected to study the effect of the increase of the pH solution on its chemical stability in time until 7 days by adding NaOH drops every 5 min. Between pH 5 and 8 the color of the solution did not evolve from transparent to green or yellow, which are respectively, typical of Ce(OH)3 and CeO2 [39], i.e. there was no precipitation of the cerium compounds in the solution.

The evolution of the potential with electrodeposition time and with increasing pH from 5 to 8 and applied current density of 0.5 mA/cm² is shown in Fig. 3. All curves display a sudden decrease of the potential in the first period (t < 250 s) followed by a stabilization period which is attributed to the formation of a film covering the whole surface of the substrate [35]. The shift of the potentials to more negative values is related to the increase of the film thickness [42]. These phenomena are more marked with increasing the pH of the solution. Moreover, the shape of the curves looks similar irrespective of the pH, which suggests that the growth mechanisms of the deposits are equivalent [43].

The SEM images show a surface totally covered with a deposit of acicular and laminated structure (Figure 4). The needle-like morphology is typical of hydroxide bonding [35] while the laminated is associated with the oxide [17]. The latter increases with increasing the pH, i.e.

the more alkaline the solutions the more favoured is the oxide. However, cracking of the coatings is also promoted with increasing pH. The cracks can be generated by the shear stresses between the coating and the substrate [44] or by the hydrogen evolution reaction (HER) that induce stresses in the coating itself [45].

The cyclic polarization curves obtained from cerium nitrate solution with and without acetic acid addition confirmed this hypothesis (Fig 5). The potential shifts towards more cathodic values and the hysteresis is more important with 0.008M than with 0.1 M of acetic acid.

Moreover, the addition 0.008M of acetic acid results in a hysteresis similar to the 0.1M Ce(NO3)3,6H2O solution. The shift of the potential and the hysteresis was already observed in acetic acid-free 0.1 M Ce(NO3)3,6H2O solutions [35] and was ascribed to the formation of the cerium hydroxide then to the cerium oxide by oxidation of the Ce 3+ compounds. However, the deposits contained cracks. Here, the low concentrations of acetic acid allowed to retain minute amounts of Ce (III) that can be released from the complex according to equation (5), i.e. the reverse of equation ( 4) [39]. The maximum concentration of [Ce 3+ ] that can be released is given by that of the cerium acetate (Table 1).

2 Ce(III)-L ⇄ 2 Ce 3+ + 2 L (5)

In contrast, with the high concentration of acetic acid (0.1M), the deposit cannot form because of the hydrogen evolution reaction (HER) following reactions ( 6) and ( 7) [46] :

Fe + HAc + 1e -⇄ FeH + Ac - (6)
FeH + Hads ⇄ H2 + Fe (7) where HAc and Ac -are, respectively, the acetic acid and the acetate anion.

Indeed, the release of gas bubbles was observed at the naked eye and the E vs. time curves appear noisy.

Irrespective of the pH and with 0.008M of acetic acid, all the patterns display the typical fluorite structure of CeO2 in addition of the peaks of the substrate (Figure 6). However, at pH 5, 6 and 7, two additional peaks appear at 10.5 and 22° which can be related to the carbonated green rust [47]. In contrast, such green rust is not observed at pH 8. J-P. Viricelle reports that the peak at 10.4° corresponds to Ce2(CO3)3,8H2O [START_REF] Viricelle | Etude de la transformation thermique de hydroxycarbonate de cérium (CeOHCO3) en dioxyde de cérium IV[END_REF]. Therefore, FT-IR analyses were performed for a more thorough assessment of the compound formed at pH 8 (Fig. 7). The peaks at 830, 1038 and 1385 cm -1 are associated with nitrates while those at 1320 and 1473 cm -1 correspond to carbonates when compared to the FT-IR spectra of Ce(NO3)3,6H2O and Ce2(CO3)3,8H2O [START_REF] Viricelle | Etude de la transformation thermique de hydroxycarbonate de cérium (CeOHCO3) en dioxyde de cérium IV[END_REF].

Therefore, it appears that the XRD peak at 10.5° corresponds rather to carbonate species instead of the green rust. In addition, the peaks located at 1320, 1750 and 2495 cm -1 can be assigned to C-O, C=O stretching vibration and C-H asymmetric stretching in CH3 in plane bending respectively [START_REF] Mani | [END_REF]50]. The strong and wide bond appearing at 3300 cm -1 is assigned to OH - stretching. All the latter are indicative of the presence of the molecular form of acetic acid whether adsorbed on the surface of the electrode or entrapped in the deposits.

The presence of acetic acid in the deposit was investigated by DSC (Figure 8). Table 2 gathers the main results of the thermal characterization such as humidity, amorphous phase, decomposition, transformation as well as the temperature ranges. In general, both curves display the same phenomena upon heating which consist first in the release of water and in the transformation from the amorphous hydroxide into crystalline ceria [17]. However, the exact temperature and the energy associated with the thermal events are different at pH 8 than at the solution pH. In particular, water seems to be released of the pH 8 deposits at 76°C while dehydration occurs at about 130°C in the deposits obtained at the solution pH. Similarly, the temperature at which the transformation from the amorphous to crystalline phase occurs is lower in the deposits obtained at pH 8. This is indicative of a more crystallized deposit at pH 8. Also, an additional intense exothermic peak at 257°C is observed at pH 8 that can be attributed to the decomposition of acetic acid [51] although J-P. Viricelle attributes the thermal evolution at 270°C to CeOHCO3 issued from the hydrolysis of Ce2(CO3)3,8H2O with acetic acid [START_REF] Mani | [END_REF]. In addition, the transformation of Ce(III) and Ce(IV) hydroxides to Ce(IV) oxide at pH 8 is accompanied by a greater absorption of energy which indicates that the oxidation of the deposits prepared at pH 8 occurs more slowly compared to those obtained at the solution pH (3.87). In summary, it appears that the deposits formed with acetic acid and pH 8 contain less water, are more oxidized, i.e. CeO2 and probably contain carbonates, nitrates and acetic acid.

Protection afforded by the deposits

The protection afforded by these electrodeposits was conducted by exposing the coated samples to ambient air for 30 days (720 h) and room temperature. As opposed to the coatings obtained at different pHs, the one at pH 8 did not show any corrosion product at the surface (Fig 9). In agreement with our previous work [52], we can conclude that the substrate dissolution occurs just after immersion of the electrode in the solution which leads to the formation of rust beneath the ceria based coating. which the establishment and dissolution of corrosion film on the substrate surface are concomitant [53]. Between 350 and 500 h, the Eocp evolved towards the substrate potential.

These results indicate that the sample is still under active corrosion process where the coating is unable to provide protection.

However, the Eocp for the coatings prepared at pH 8 is lower than that for coating obtained at solution pH. As with the solution pH, the Eocp evolves during the first hour but the potential value does not exceed the potential of the initial coating (-0.61V/SCE) and then stabilizes at -0.64V/SCE along the whole exposure. Since this Eocp is sufficiently far from the potential of the steel substrate, one can presume that the evolution during the first hours arises from the penetration of the aggressive solution through cracks and defects in the coating that ensures the self-healing ability of the ceria based coatings proposed by Ferreira et al. and Ma and collab. [54,55] In order to compare the appearance of the surface of both coatings after the immersion period, the evolution of the surface was examined by using optical microscopy (Fig. 11). It can be clearly seen that the coatings prepared at solution pH show full degradation from 2 days of immersion in NaCl solution where the red rust (substrate dissolution) covers the whole surface.

However, the coatings prepared with acetic acid and pH8 continue to be resistant against corrosion throughout the whole testing period.

Potentiodynamic polarization curves were performed during 30 days (720 h) of immersion in 0.5 M NaCl solution (Fig. 12). The corresponding corrosion current densities "Icorr" using the Tafel extrapolation are depicted in Fig. 13 and the data gathered in Table 3. Fig. 12 shows that the Ecorr shifts continuously towards the cathodic domain with immersion time. In addition, the shape of the anodic branches of all the curves are less affected with time compared to the cathodic ones, which indicates that the corrosion process is under cathodic control. At the beginning of the immersion, the Icorr for the coated substrate was around 39 µA/cm² which is 4 times lower than the value of the untreated substrate (111 µA/cm²). During the first 48h, the Icorr dropped dramatically, then tended to increase slowly till a sharp increase at 360 h (15 days). Thereafter, the Icorr dropped again and remained stable till 720 h (30 days). Such fluctuations of the Icorr values suggest that the coating remains active till completely healed with corrosion products. This hypothesis is confirmed by the evolution of the polarization resistance Rp with time (Fig. 13) that follows the opposite trend of Icorr.

Conclusion

The

  This positive result encouraged us to assess electrochemically the coating obtained from 0.1 M Ce(NO3)3 at pH 8. Potentiodynamic polarization tests were carried out in 0.5 M NaCl for 720 hours (30 days). Fig. 10 presents the typical chronopotentiometric curves obtained for the samples coated at pH8 and at the solution pH (3.87). At the solution pH, the Eocp shifted strongly towards more cathodic values at the initial stage of immersion. Then, the Eocp moved to anodic direction reaching the steel substrate potential (-0.492 mV/SCE). The fluctuation of Eocp between 50 and 350 h close to the potential of steel can be related to a transient stage in

  addition of low concentrations of acetic acid to 0.1M Ce(NO3)3, 6H2O solutions allows to prevent chemical precipitation of Ce(OH)3 upon cathodic electrodeposition at -0.5 mA/cm 2 for 20 min on a carbon steel substrate. The increase of pH by adding NaOH drops result in quite protective coatings in ambient air for 30 days. The cerium oxide based coatings seem to trap carbonates, nitrates and acetic acid. When immersed in 0.5M NaCl, the coatings require at least 15 days (320h) to stabilize. Then, the Eocp, Icorr and Rp tend to being stable where they come to confirm the good electrochemical behavior of the coating obtained at pH 8 compared to those obtained at the solution pH.

Figure 1 :Figure 2 :Figure 3 :

 123 Figure 1: XRD pattern of the cerium oxide film elaborated in 0.1M Ce(NO3)3 with (0.1M) acetic acid addition at room temperature and 0.5 mA/cm² as applied current density during 20 min.

Figure 4 :Figure 5 :

 45 Figure 4 : SEM Images for deposits obtained from Ce(NO3)3 0,1 M and acetic acid 0.008 M at solution pH adjusted to (a) 5, (b) 6, (c) 7 and (d) 8 with I= 0,5 mA/cm² during 20 min at room temperature.

Figure 6 :

 6 Figure 6: XRD pattern of the cerium oxide film elaborated in 0.1M Ce(NO3)3 with (0.008M) acetic acid addition at room temperature and 0.5 mA/cm² as applied current density during 20 min at different solution pH.

Figure 7 :Figure 8 : 8 0Figure 9 :Figure 10 :Figure 11 :Figure 12 :Figure 13 :

 788910111213 Figure 7: Evolution of the surface of the deposits obtained from 0.1 M Ce(NO3)3 + 0.008 M acetic acid at (a) pH 5 (b) pH 6 (c) pH 7 (d) pH 8 during 20 min and 0.5mA/cm² as applied current density and exposed open to air for (top) 48 hours (second line) 5 days (third line) 15 days (bottom ) 30 days.

  1.68 = 47.86 This results in [Ce(CH3COO)3] = 2.4 10 -4 M and [H + ] = 7.1 10 -4 M. Similarly, the concentrations of cerium acetate and of protons can be calculated for the different concentrations of acetic acid added to the cerium nitrate solution (Table

Table 1 :

 1 pH of the different solutions and of the calculated concentrations of cerium acetate

	and of protons. NB: HA = acetic acid.
	[HA] (M)	0.008	0.01	0.03	0.06	0.1
	pH with 0.1 M Ce + HA	3.17	3.15	3.00	2.87	2.73
	[𝐶𝑒(𝐶𝐻 3 𝐶𝑂𝑂) 3 ] (M) [𝐻 + ] (M)	2.4 10 -4 7.1 10 -4	3.3 10 -3 9.9 10 -3	8.9 10 -3 2.7 10 -2	1.7 10 -2 5.2 10 -2	1.9 10 -2 5.6 10 -2

Table 2 :

 2 Thermal data for the indicated systems

	Transformation	Solution pH		Acetic acid and pH 8
		Tmax (°C) T (°C) H (J/g) Tmax (°C) T (°C) H (J/g)
	Water release	130	[110-141]	5.62	76	[60-93]	9.28
	amorphous phase	145	[150-159]	6.61	127	[120-139]	4.12
	Acid acetic	------	-------	--------	257	[257-274]	29.75
	decomposition
	Ce 3+ to Ce 4+	367	[324-336]	0.097	367	[287-330]	2.13
	Ce(OH)4 to CeO2	438	[418-451]	2.17	445	[427-463]	3.99

Table 3 :

 3 Electrochemical parameters calculated from polarization measurements on uncoated


Claire Boulesteix
email: claire.boulesteix@univ-lr.fr

Vladislav Kolarik

Fernando Pedraza
email: fernando.pedraza@univ-lr.fr

Steam Oxidation of Aluminide Coatings under High Pressure and for Long Exposures

Keywords: Ferritic-martensitic steels, Aluminium diffusion coatings, Oxidation, Steam, Pressure

B2-FeAl-based aluminide coatings were elaborated on P92 ferritic-martensitic steels by slurry. Under atmospheric pressure of steal and 650°C, a thin protective α-Al 2 O 3 oxide layer formed and was maintained until 5000 h. With increasing pressure to 300 bar, the morphology of α-Al 2 O 3 changed markedly. Iron oxide clusters started to appear at the surface of the scale after 5000 h of exposure and kept on growing till 10000 h. It appears that pressure increased the outward diffusion of Fe through the coating and through the oxide scale and fostered the propagation of the tensile cracks of the coating.

INTRODUCTION

Greater temperatures and pressures allow significant increases of efficiency in steam power plants [1,2]. These improvements require metallic materials with adequate mechanical and oxidation resistances [2][START_REF] Holcomb | Ultra-supercritical steam corrosion, 17 th Annual Conference on Fossil Energy Materials[END_REF][START_REF] Viswanathan | [END_REF][5]. In this case, heat resistant ferritic-martensitic steels are commonly used for superheater tubes and pipes in ultrasupercritical boilers [6,7]. They offer a better heat transfer behaviour and lower thermal expansion coefficients [6,8,9] than austenitic stainless steels or the more expensive nickel base alloys Among this steel range, the T/P92 heat resistant steels (ASME grade 92) have been recognized as key structural components for ultrasupercritical power plants [10]. Despite its good properties, the P92 ferritic-martensitic steel has to be employed up to 650°C to maintain its mechanical resistance [6].

The behaviour of this material has been studied in steam environments under different pressures [11,12]. Agüero et al. have shown that after 3000 h under atmospheric pressure and in pure steam, an oxide scale of about 180 µm thick, composed by two major layers constituted by magnetite, was formed. Additional (Fe,Cr) 3 O 4 spinels were found in the inner part of the oxide layer [11]. Similar results were reported by Boulesteix et al. after 2000 h of exposure in steam and under 1 bar [12]. Ehlers et al. found an additional thin outer layer of hematite in water vapour containing atmospheres [13]. By increasing the pressure to 180 bar, the oxide formed in steam was thicker but displayed the same layered structure and oxide phases. In contrast, a higher degree of porosity and delamination from the substrate were observed [11].

Yin et al. have conducted experiments in supercritical water for several hundred h at 600°C and under 250 bar. Again, a bilayered oxide structure was reported. By increasing the pressure, it seems that the porosity increased but by increasing the exposure time, this phenomenon tends to stabilize [14]. Moreover, Zhu et al. have tested the P92 ferritic-martensitic steel in deaerated supercritical water under the same pressure but at 650°C which corresponds to the maximal operating temperature of this alloy. The magnetite/spinel-mixed structure was non adherent and internal oxidation phenomenon of the substrate was noticed [14]. Furthermore, above 600°C, evaporation of volatile CrO 2 (OH) 2 species became significant and the breakdown of the oxide layer formed occurred [6,15]. In all cases, the oxide scale formed is non protective and in power plants, in view of the significant oxide thicknesses encountered, it could cause the clogging of the boiler tubes and their subsequent burst.

Therefore, the elaboration of a protective coating appears to be necessary and, as it happens, the slurry aluminium coating is a promising way. Indeed, this technique is a low cost and easy-to-apply process which allows to coat large components irrespective of its shape. The slurry developed in our research group is "environmentally friendly" because water is used as a solvent [16]. The mechanisms of formation of these coatings on austenitic and ferriticmartensitic steels have been largely described in [17][18][19][START_REF] Boulesteix | Characterisation of Aluminium Diffusion Coatings Elaborated on Austenitic Stainless Steels and on Ferritic-Martensitic Steels[END_REF]. The good behaviour of these coatings in air, argon or steam at intermediate temperatures (650 < T < 700°C) has been demonstrated by previous works conducted in our group [12,19,[START_REF] Pedraza | [END_REF]. As it happens, it was shown that the steam oxidation resistance of ferritic-martensitic and ofr austenitic stainless steels was greatly improved at respectively, 650°C and 700°C and under 1 bar when the materials were coated with a slurry aluminium coating in spite of the presence of tensile cracks. Indeed, the formation of a thin oxide layer exclusively composed by α-Al 2 O 3 ensured the protection of the tested materials even after thousands of h [11,12,[START_REF] Pedraza | [END_REF]22]. However, by increasing the pressure, the degradation of the coating becomes faster and it appears that the scales formed are thicker [11].

Based on these observations, it is difficult to ascertain whether the detrimental effect derives from cyclic conditions in the plant or from the pressure itself.

Thus, this paper investigates the effect of pressure (1 vs. 300 bar) on a slurry aluminised P92 steel exposed to steam, at 650°C, for 5000 h. A second part of the paper shows the evolution of the slurry aluminised P92 steel exposed to 300 bar till 10000 h.

EXPERIMENTAL PROCEDURE

Steel substrate

P92 coupons (20x10x2 mm) were ground with SiC #180 paper, grit blasted with a 220 mesh Al 2 O 3 particles, then rinsed in deionised water and finally cleaned with ethyl alcohol under ultrasonic agitation before drying with hot air. The compositions (nominal and determined by Energy Dispersive Spectroscopy (EDS) analysis) of this alloy are given in Table 1.

Aluminisation

The aluminium micro-particles employed to elaborate the slurry were the same used in our previous works [12,19,[START_REF] Boulesteix | Characterisation of Aluminium Diffusion Coatings Elaborated on Austenitic Stainless Steels and on Ferritic-Martensitic Steels[END_REF][START_REF] Pedraza | [END_REF]23] (4-5 µm average size, Hermillon, France). The P92 samples were coated by spraying a slurry containing 43wt% Al microparticles dispersed in 57wt% of binder (1/10 polyvynil alcohol/deionized water). After drying in ambient air for at least 1 hour [16], the coated samples were annealed in argon (400°C/3h + 650°C/3h + 1050°C/1h) [START_REF] Boulesteix | Characterisation of Aluminium Diffusion Coatings Elaborated on Austenitic Stainless Steels and on Ferritic-Martensitic Steels[END_REF]. The O 2 partial pressure in the furnace where the thermal treatment takes place is lower than 10 -6 mbar and by referring to the Ellingham diagram [START_REF] Mitchell | An introduction to materials engineering and science for chemical and materials engineers[END_REF], it appears that Al 2 O 3 can be formed. It was shown that in these conditions, α-Al 2 O 3 was formed after annealing of the aluminised samples [17,[START_REF] Mollard | Elaboration de systèmes barrière thermique par barbotine: comportement du nickel et de ses superalliages revêtus en oxidation cyclique[END_REF]. Once the thermal treatment done, the samples were grit blasted with a mesh Al 2 O 3 in order to eliminate the residual empty microparticles on the coated material ("bisque"). Nevertheless, the coating appeared preoxidised after the annealing [F. Pedraza, M.

Steam tests

The isothermal oxidation tests were carried out at 650°C and under 1 bar. For these experiments, the aluminised P92 samples were introduced in a tubular furnace and heated at 20°C/min until 650°C and were subsequently maintained for 5000 h with a constant steam flow (7 x 10 -2 m.s - 1 ). The steam was produced from deionized water using silica filtered and content of O 2 dissolved was lower than 0.05 mg.L -1 .

The high pressure tests were conducted at the Fraunhofer Institut, Pfinztal (Germany).

Laboratory tests autoclaves from Alloy 602 were closed by welding, inserted into a muffle furnace and connected to the pressured gas supply system. The samples were oxidised for up to 10000 h at 650°C under a pressure of 300 bar [START_REF] Juez-Lorenzo | 9 th International Sumposium of High-Temperature Corrosion and Protection of Materials HTCPM 9, Oxidation of Metals[END_REF].

Characterisation of the samples

Before and after oxidation, the crystal structures of the samples were determined by X-Ray Diffraction (XRD) using a Bruker AXS D8 Advance (CuK α radiation), in θ-2θ mode. Morevoer, the surfaces and the cross-sections of the aluminised P92 samples were characterised by scanning-electron microscopy (SEM) coupled to energy-dispersive spectrometry (EDS) using a FEI Quanta 200F Field Emission Gun Scanning Electron Microscope (FEG-SEM) at 20 kV with an EDAX detector operating under low vacuum. In order to observe the spatial distribution of the elements in a wider area and to observe the oxide layers formed, X-ray maps were realized after oxidation tests. The cross-sections of the oxidised samples were also observed using a Leica DMRM microscope and a LEICA MC170HD camera coupled with the software LAS v4.9 for the images acquisition. Raman microspectrometry analyses were also performed with a Horiba Jobin Yvon LabRam HR apparatus using a He-Ne laser (632.8 nm) in order to obtain more information about the oxide layers. The analyses were performed every 500 µm from the centre to the edges to map the likely evolutions in composition at local scale.

RESULTS

Aluminide coatings elaborated on P92

XRD characterisations were performed on the raw P92 steel, on the "as-sprayed" coating (before aluminisation) and after the thermal treatment allowing the formation of the aluminide coating. As shown in Fig. 1, the six characteristic peaks of the B2-FeAl phase were detected (JCPDS 00-033-0020) on the aluminised P92 steel. The blue diffractogram which shows the "as-sprayed" specimen let appeared one peak with a lower intensity (2θ ≈ 37°) which can be attributed to the Al oxyhydroxides contained in the slurry [16]. ). As previously seen in [START_REF] Boulesteix | Characterisation of Aluminium Diffusion Coatings Elaborated on Austenitic Stainless Steels and on Ferritic-Martensitic Steels[END_REF][START_REF] Montero | [END_REF], the coating elaborated on the P92 ferritic-martensitic steel can be divided in two parts. The first part corresponded to the diffusion zone (or additive layer) and was about 60 µm thick. The EDS analyses (Fig. 2. c.)

performed on this outer layer confirmed the XRD findings (Fig. 1). Indeed, by referring to the Fe-Al binary diagram [START_REF]Alloy Phase Digrams[END_REF], with 40 at% Al + 54 at% Fe at the top of the coating and 25 at% Al + 68 at% Fe at the interface between the additive layer and the interdiffusion zone, the major phase contained in the diffusion zone corresponded to B2-FeAl. The inner part of the coating (the interdiffusion zone) was similar in terms of thickness (67 µm thick) but was less homogenous compared to the first one. Needle-like precipitates were observed and were attributed to Al nitrides, due to the presence of N in the substrate, as previously reported in [START_REF] Boulesteix | Characterisation of Aluminium Diffusion Coatings Elaborated on Austenitic Stainless Steels and on Ferritic-Martensitic Steels[END_REF]29,30]. Moreover, the diffusion coefficients of Al and of the compounds of the substrate were different, a Kirkendall effect was induced involving some porosities in the coating [START_REF] Boulesteix | Characterisation of Aluminium Diffusion Coatings Elaborated on Austenitic Stainless Steels and on Ferritic-Martensitic Steels[END_REF]. Cracks were also observed and were correlated to the mismatch of the thermal expansion coefficient existing between the different metallurgical phases [11,[START_REF] Boulesteix | Characterisation of Aluminium Diffusion Coatings Elaborated on Austenitic Stainless Steels and on Ferritic-Martensitic Steels[END_REF][START_REF] Boulesteix | Characterisation of Aluminium Diffusion Coatings Elaborated on Austenitic Stainless Steels and on Ferritic-Martensitic Steels[END_REF]30]. Corresponding EDS profiles.

Influence of pressure -1 vs 300 bar -5000 h of exposure

The aluminised P92 ferritic-martensitic steels were exposed for 5000 h at 650°C in steam and under 1 or 300 bar of pressure.

The surface of the tested steel after 5000 h of exposure under atmospheric pressure (Fig. 3. a.) appeared similar compared to those obtained after aluminisation and grit blasting (Fig. 2. a.).

Indeed, the roughness of the sample was still observable and cracks were detected on its surface but they appeared self-healed as highlighted in the Xray maps shown in Fig. 4. At higher magnification (Fig. 3. b.), whiskers covering the entire oxidised surface were observed. As shown in Fig. 4, in some parts of the coating, the cracks propagated beyond the additive layer/interdiffusion zone interface. However, aluminium and oxygen have been detected into them, as well as on the surface. Indeed, a very thin scale can be observed on the samples and can be compared to the results obtained in [12] after 2000 h in the same conditions. The additive layer thickness was about 5 µm thinner than before oxidation. On the contrary, the interdiffusion zone appeared about 5 µm thicker(∼ 70 µm) after 5000 h of exposure under atmospheric pressure. Moreover, the Cr segregation at the grain boundaries of the coating area between the interdiffusion zone and the substrate was clearly marked in the X-ray maps (Fig. 4). Compared to the initial concentrations, the Al content decreased more markedly (about 10 at%)

with 300 bar than with 1 bar. hConversely, the Fe content increased in the same proportion.

The Cr contents were similar before and after these oxidation tests.

The X-ray maps shown in Fig. 8 highlight the presence of cracks but like with 1 bar, they appeared self-healed. Al and O were also detected on the surface of the sample where a thicker oxide scale (compared with Fig. 4) was formed. The oxide formed looked more compact but two layers appeared. Some traces of Fe can be distinguished on the outer part of the oxide layer.

Concerning the Raman and XRD analyses (Figs. 6), no signal corresponding to the presence of an iron oxide was detected. Indeed, the Raman analyses showed two additional peaks in the Raman domain (Fig. 6. a.) at 650 and 900 cm -1 . Jian et al. [33] have related the first one to the After 5000 h at high pressure, the coating exhibited a thinner additive layer of about 52 µm compared to the fresh one (60 µm). On the other hand, the interdiffusion zone appeared thicker after the oxidation tests (72 µm).

Influence of exposure time -300 bar -5000 and 10000 h of exposure

The high pressure steam tests were extended to 10000 h of exposure. The surface morphology of the oxidised P92 ferritic-martensitic steel is shown in Figs. 9. The platelets were still present and the coverage of the surface was comparable to Fig. 3. c. However, after 10000 h of exposure, clusters were observed in greater extent and their size was bigger compared to those observed after 5000 h. on the analysed areas). Nevertheless, the X-ray maps shown in Fig. 10 were similar compared to those obtained after 5000 h of exposure. Indeed, a dual oxide layer was still observed and Al and O were detected in a large amount in the cracks and on the surface of the oxidised sample.

Fe was also observed in some parts of the oxide but in a smaller extent compared to Al and O.

The Raman analyses (Figs. 11. a. b.) let appear two types of signal depending on the analysed area. Indeed, as shown in Fig. 11. a., the spectrum obtained on the platelets was the same compared to that obtained after 5000 h of exposure under high pressure with the appearance of the α-Al 2 O 3 peaks. On the contrary, the signal obtained on the clusters had 4 additional peaks corresponding to Fe 2 O 3 [START_REF] Maslar | [END_REF][36][37]. The two peaks corresponding to the α-Al 2 O 3 (650 and 900 cm - 1 ) were still observed and two less defined peaks were detected around the alumina peak at 650 cm -1 . Indeed, signals at 625 and 725 cm -1 were recorded and can be correlated to the presence of iron spinels, as reported in [START_REF] Edwards | Raman Spectroscopy in Archaeology and Art History[END_REF].

As regards the coating itself, the Al concentration decreased at the surface of the additive layer to reach about 39 at% (Fig. 5). On the other hand, the Fe concentration was about 54 at% on average on the entire thickness of the diffusion zone. It has been noticed that the variations of Cr with time were negligible in the additive layer with time. However, the Cr content evolved from 6 at% (aluminised samples) to 8 at% after 10000 h of exposure below the coating/substrate interface. The thickness of the additive layer tended to decrease with time whereas the interdiffusion zone evolved in the opposite trend (48 and 72 µm, respectively).

Moreover, the cracks still present after aluminisation seemed to propagate to the end of the interdiffusion but without reaching the steel (Fig. 12). Their propagation under high pressure seemed more significantly than under 1 bar. The Al enrichment zone observed below the interdiffusion zone/substrate interface seemed to increase with time. Indeed, after aluminisation this zone was about 33 µm whereas after 5000 h under high pressure it was 44 µm and reached 53 µm after 10000 h.

DISCUSSION

Aluminised P92 ferritic-martensitic steel

Aluminium slurry coatings were elaborated on P92 ferritic-martensitic by spraying a waterbased slurry. The following annealing step allowed the simultaneous Al inward diffusion and the Fe outward diffusion [17,18,[START_REF] Boulesteix | Characterisation of Aluminium Diffusion Coatings Elaborated on Austenitic Stainless Steels and on Ferritic-Martensitic Steels[END_REF], thus resulting in the formation of the B2-FeAl phase. The thermal treatment carried out in this work allowed the formation of shells around the empty Al microspheres. These shells were oxidised due to the low O 2 partial pressure of the annealing gas (Ar), thus oxidising the surface of the coated material [START_REF] Mollard | Elaboration de systèmes barrière thermique par barbotine: comportement du nickel et de ses superalliages revêtus en oxidation cyclique[END_REF][START_REF] Pedraza | [END_REF]. In these conditions, the aluminised ferritic-martensitic samples appeared to be preoxidised with the appearance of a very thin alumina layer (less than 1 µm) [17]. The coatings elaborated in this study had a good homogeneity but some cracks were detected after the thermal treatment. This phenomenon was correlated in previous studies to the mismatch of the thermal expansion coefficients between the different metallurgical phases [11,[START_REF] Boulesteix | Characterisation of Aluminium Diffusion Coatings Elaborated on Austenitic Stainless Steels and on Ferritic-Martensitic Steels[END_REF]29,30]. Needle-like AlN nitrides were also detected in the interdiffusion zone of coating due to the presence of nitrogen in the substrate [40].

Moreover, Boulesteix et al. have shown the presence of a Cr segregation at the interface between the interdiffusion zone and the substrate [START_REF] Boulesteix | Characterisation of Aluminium Diffusion Coatings Elaborated on Austenitic Stainless Steels and on Ferritic-Martensitic Steels[END_REF], that can act as a diffusion barrier for the compounds contained in the substrate [41].

Influence of a high-pressure steam exposure for 5000 hours of exposure

Steam tests were performed for 5000 h at 650°C and under atmospheric pressure (1 bar). The surface of the oxidised samples (Fig. 3. a.) was similar compared to the aluminised one (Fig. 2.

a.) with the traces of the light grit blasting carried out after aluminisation. At higher magnification, it is clear that the surface was covered by oxide whiskers as shown in Fig. 3. b.

The coating looked similar to the "fresh-ones" in terms of homogeneity but the thickness of the additive layer seemed to decrease from 59 to 54 µm after 5000 h of steam exposure. In contrast, the interdiffusion zone grew thicker. This phenomenon was also reported by Agüero and Muelas after 25000h of exposure of Fe 2 Al 5 -based aluminide to 1 bar of steam [40].

Concerning the Al and Fe content (Table 2), the first one decreased with time (∼ 5 at% Al on the surface of the additive layer), as shown in Fig. 5, whereas the Fe concentration followed the opposite trend (Fig. 6). By comparing with the initial concentrations, it appeared that the Al content was higher at the interface between the coating and the substrate. Bellucci et al. [START_REF] Bellucci | EPRI International Conference on Corrosion in Power Plants[END_REF] and Agüero and Muelas [40] also showed that the Al concentration at the surface of an aluminide coating deposited on different ferritic and ferritic-martensitic steels tended to decrease with time. This phenomenon was directly linked to the Al consumption to form an alumina layer and to Al inward diffusion in the substrate. The X-ray maps (Fig. 4) highlighted a distribution of Al, Fe and Cr comparable to the one obtained after aluminisation with Cr segregating at the grain boundaries of the interdiffusion zone of the coating [START_REF] Boulesteix | Characterisation of Aluminium Diffusion Coatings Elaborated on Austenitic Stainless Steels and on Ferritic-Martensitic Steels[END_REF] that may block the interdiffusion of Al and Fe [41].

The Raman and XRD analyses (Figs. 6) carried out on the coated P92 ferritic-martensitic steel exposed for 5000 h in steam at 650°C and 1 bar let appear the characteristic signals of the presence of an α-Al 2 O 3 layer. Moreover, the X-ray maps and the EDS analyses realised on these samples confirmed the high amounts of Al and O on the surface of the oxidised samples (Fig. 4). The whisker-like morphology of the oxides shown in Fig. 3. b. demonstrated that the potential transformation from the metastable to the stable alumina occurred [START_REF] Di Girolamo | [END_REF]. Indeed, previous studies [44][45][46] have shown that the evolution of the γ-Al 2 O 3 to θ-Al 2 O 3 and from θ- was observed [49]. The thermal treatment carried out in this study to aluminise the P92 ferriticmartensitic steel allowed the preoxidation of the coated substrate, this probably allowing the faster formation of the stable alumina.

Once the good behaviour of the aluminide coatings checked under atmospheric pressure for 5000 h, the pressure was increased to reach 300 bar. In this case, the oxide morphology was completely different with the appearance of more or less large platelets covering the entire surface of the substrate (Figs. 3. c. d.). The aspect of the as-aluminised surface after grit blasting was not observed anymore after 5000h of exposure to 1 bar of steam. The oxide scale appeared more compact compared to the one obtained under 1 bar (Figs. 3. a. b.). Moreover, some small clusters were observed in some parts of the coating (Fig. 3. D). The chemical homogeneity of the coating was basically kept despite the slight decrease of the thickness of the additive layer and the corresponding small increase of the thickness of the interdiffusion zone. These evolutions can be correlated to the simultaneous Al outward diffusion to form and to maintain an oxide scale and to the Al inward diffusion due to its high coefficient diffusion in Fe [START_REF] Agüero | Overview of steam oxidation behaviour of Al protective oxide precursor coatings on P92[END_REF].

However, the segregation of Cr at the grain boundaries of the interdiffusion zone is believed to block any substantial loss of Al from the coating [41]. The oxide layer formed appeared thicker (Fig. 8) compared to the one obtained with 1 barfor the same time but as with 1 bar, the scale displayed a bilayered structure of oxides mostly composed by Al and O. [54] and Otoguro explained that differences between the metallic ions diffusion rates can be created due to the pressure effect [55]. Holcomb also considers that the solid state diffusion within the Fe and Cr-oxide scales is sufficiently quantifiable but most importantly that the hydrogen injected from steam oxidation increases the oxygen permeability lowering the ability of the alloy to establish a chromia scale and therefore the critical content of chromium to form the chromia scales is also increased [START_REF] Holcomb | [END_REF]. Such reports are very much in line with our observations where Fe incorporated in the scale despite the sufficient amounts of Al in the coating to maintain a pure Al 2 O 3 and to explain the changes in morphology of Al 2 O 3 . However, one shall notice that these effects were quite limited as the iron-containing oxide clusters were limited in size and presence and only the local Raman spectra allowed to reveal them at the surface of the scale.

The influence of pressure on the coatings themselves corresponds also to an increase of solid state diffusion that slightly decreases the thickness of the diffusion zone while that of the interdiffusion is increased. However, the force exerted in the opened cracks make them propagate towards the substrate but they heal simultaneously through the formation of aluminium oxides (Fig. 8) also observed in different works either in steam [29,29] or in air [12].

Influence of a high-pressure steam exposure with time

The high pressure steam tests were finally extended to 10000 h in order to observe possible changes in the coating or in the oxide scale. As shown in Figs. 9, the oxide morphology obtained after 10000 h of exposure was similar to that obtained after 5000 h in the same exposure conditions. The thickness of the diffusion zone of the aluminide coatings continued to decrease thus allowing a slight thickness increase of the interdiffusion zone. The Al content revealed at the surface of the additive layer was about 10 at% lower than that observed after 5000 h of exposure (Figs. 5 and7). This is in agreement with the works of Bellucci et al. because of the outward diffusion of Al to maintain the oxide layer which increased with time (Fig. 10) and the subsequent Al inward diffusion [40,[START_REF] Bellucci | EPRI International Conference on Corrosion in Power Plants[END_REF] . The oxide scale continued to grow to reach few additional microns compared to the 5000 h of exposure. A two-layered structure and the platelets were still observed (Fig. 10) but the proportion of clusters increased (Fig. 9. b.). The EDS analyses carried out on areas of the oxide scale were completely different depending on the analysed areas. Indeed, high amounts of O, Al and Fe were revealed on the clusters (, 56 O, 24 Al and 20 Fe on average). However, 55 O and 43 Al (at%) were observed on the platelets and, therefore, very few amounts of Fe (< 1.5 at%). The Raman analyses carried out on reflected such differences in local composition. Indeed, as shown in Fig. 11. a., the spectrum obtained on the platelets was the same compared to that obtained after 5000 h of exposure under high pressure (α-Al 2 O 3 , Fig. 6. a.). In contrast, six additional peaks appeared on the clusters. Four of them at low wave numbers were characteristic of haematite, as previously demonstrated by different authors [START_REF] Maslar | [END_REF][36][37]. At higher wave numbers, the α-Al 2 O 3 signal was still present but the peak at 650 cm -1 was surrounded by two other signals which were associated with the formation of Al-Fe spinels [START_REF] Edwards | Raman Spectroscopy in Archaeology and Art History[END_REF]. The EDS analyses carried out on the platelets did not show Cr, thus confirming the Raman and XRD analyses (Figs. 11). Moreover, the characteristic doublet of the stable alumina in the fluorescence domain was found for both oxidation times. The enhanced Fe outward diffusion through the oxide scale [START_REF] Holcomb | [END_REF]54,55] seemed to be confirmed by the appearance of Fe 2 O 3 and Fe spinels. Moreover, it has been reported that the pressure involved a higher porosity degree in the oxide scales, thus providing a path of enhanced diffusivity and explaining the pressure effect on the oxidation of the coated P92 ferriticmartensitic steel [START_REF] Angell | 15 th International Conference of the Properties of Water and Steam[END_REF].

In contrast, the was no marked change in the composition of the healed cracks with alumina in spite of their propagation with time at 300 bar did not contribute (Figs. 10 and12).

CONCLUSIONS

Steam oxidation tests were performed on slurry aluminium coated P92 ferritic-martensitic steels.

FIGURE 1 .

 1 FIGURE 1. X-ray patterns of the P92 ferritic-martensitic steel before and after aluminisation.

Figs. 2

 2 Figs. 2. a. and b. show the micrographs obtained on the surface of the aluminised samples and on the corresponding cross section. The EDS analyses performed on the latter are also shown (Fig. 2. c.). The surface of the sample appeared to be rough because of the final step consists of a light grit blasting (Fig.2. a.). As previously seen in[START_REF] Boulesteix | Characterisation of Aluminium Diffusion Coatings Elaborated on Austenitic Stainless Steels and on Ferritic-Martensitic Steels[END_REF][START_REF] Montero | [END_REF], the coating elaborated on the

FIGURE 2 .

 2 FIGURE 2. SEM (a) surface and (b) cross section of and aluminised P92 specimen. (c)

FIGURE 3 .

 3 FIGURE 3. Surfaces of the aluminised P92 ferritic-martensitic steel exposed for 5000 hours in steam, at 650°C and under (a, b) atmospheric pressure and under (c, d) 300 bar.

FIGURE 4 .

 4 FIGURE 4. X-ray maps of the aluminised P92 ferritic-martensitic steel after 5000 hours in steam, at 650°C and under atmospheric pressure.

TABLE 2 .

 2 Fe, Al and Cr contents (± 1 at%) revealed by EDS analyses for the aluminised P92 substrate before and after 2000 h[12] and 5000 h in steam at 650°C under 1 bar. Illustration of a coated sample and the corresponding areas n°1, 2 and 3.

FIGURE 5 .

 5 FIGURE 5. Evolution of the Al content (± 1 at%) at the surface of the coating as a function of time and of pressure. 1 bar (black symboles) vs 300 bar (red symbols).

FIGURE 6 .FIGURE 7 .

 67 FIGURE 6. Raman spectra (a and b) and X-ray patterns (c) of aluminised P92 steel after 5000 h in steam at 650°C under 1 (black curves) and 300 bar (red curves).

FIGURE 8 .

 8 FIGURE 8. X-ray maps of the aluminised P92 ferritic-martensitic steel after 5000 hours in steam, at 650°C and under 300 bar of pressure.

FIGURE 9 .

 9 FIGURE 9. Surface images of the aluminised P92 ferritic-martensitic steel after 10000 hours of exposure in steam at 650°C and under 300 bar of pressure.

FIGURE 10 .

 10 FIGURE 10. X-ray maps of the aluminised P92 steel exposed for 10000 h in steam, at 650°C, under high pressure.

FIGURE 11 .FIGURE 12 .

 1112 FIGURE 11. (a to c) Raman spectra and (d) X-ray patterns of aluminised P92 steel after 5000 h (black curves) and 10000 h (red curves) in steam at 650°C and under 300 bar.

Al 2 O

 2 3 to α-Al 2 O 3 should take place at higher testing temperatures or for longer exposure time.However, α-Al 2 O 3 is directly formed upon the coating process prior to oxidation[16,[START_REF] Mollard | Elaboration de systèmes barrière thermique par barbotine: comportement du nickel et de ses superalliages revêtus en oxidation cyclique[END_REF] and the Cr contained in the aluminide could play the third element effect, thus sustaining the α-Al 2 O 3 formation [43, 47, 48]. Seraffon et al. have shown that under the same conditions that those used in this work, an aluminised T91 steel was able to form a protective Al 2 O 3 scale after 5000 h of exposure. This observation was explained by the Al diffusion within the Al-Fe intermetallic phases. Simultaneously, the Al inward diffusion from the coating to the substrate

TABLE 1 .

 1 Nominal and experimental (EDS) compositions of P92 (at%).

			P92 ferritic-martensitic steel
	Composition (at%)	Fe	Cr	C	W	Si	Others
							Mn: 0.3 -0.6
	Nominal	Bal 9.2 -10.3	0.03 -0.06 0.5 -0.6 < 1	Mo: 0.2 -0.4
							Ni: < 0.4
							V: 0.15 -0.
							Mn: 0.6
	EDS	87	10	-	0.6	0.8	Mo, Ni: < 0.4
							V, Ni: -

  Holcomb et al. reported very similar results in uncoated austenitic stainless steel and Ni-based alloys at 670°C and 267 bar[START_REF] Holcomb | 8 th International Symposium on Superalloy 718 and Derivates[END_REF][START_REF] Holcomb | [END_REF]. In this study, we have shown that in the thickness of the oxide scale increases with pressure in addition to the changes of scale composition (Fe incorporation) and morphology (platelets and clusters). Muan demonstrated that Fe can be incorporated in alumina scale due to the similar size of Fe 3+ and Al 3+ cations[53]. Wright et al. reported that the pressure could change the metal cation diffusion in oxides

Firstly, atmospheric pressure tests were carried out for 5000 h. It appeared that a thin oxide scale composed of α-Al 2 O 3 was formed. The Al consumption to form the oxide layer coupled with the simultaneous inward diffusion of Al from the coating to the substrate resulted in some decrease of Al and of the thickness of the additive layer.

With increasing pressure to 300 bar oxide morphology was completely different compared to that obtained under atmospheric pressure after 5000h of exposure. Under 1 bar, whiskers covered the entire oxidised surface whereas, under 300 bar platelets and some clusters formed at the surface. It appeared that α-Al 2 O 3 was also formed but some traces of Fe were observed in the clusters. After 10000 h, the same oxide morphology was encountered but the proportion of clusters increased. Therefore, the Fe content contained in the oxide scale increased too in some parts of the oxide. Therefore, the pressure seemed to enhance the outward diffusion of Fe through the oxide scale. Fe 2 O 3 and Fe-Al spinels were revealed by Raman spectroscopy but not by X-ray diffraction.

In addition, the initial tensile cracks of the coatings propagated with increasing pressure and time but healed with Al 2 O 3 . The Al content revealed at the surface of the coatings at end of the 300 bar pressure tests appeared sufficient to maintain the protective alumina scale.

ACKNOLEDGMENTS

The authors gratefully acknowledge European Union for the funding of FP7 project "POEMA: Production of Coatings for New Efficient and Clean Coal Power Plant Materials" (Grant agreement no. 310436). The university of Complutense (Madrid) is also acknowledged for their contribution to the steam tests carried out for 5000 hours under 1 bar.


Claire C Boulesteix
email: claire.boulesteix@univ-lr.fr

Fernando Pedraza
email: fernando.pedraza@univ-lr.fr

Characterisation of Aluminium Diffusion Coatings Elaborated on Austenitic Stainless Steels and on Ferritic-Martensitic Steels

Keywords: Austenitic steels, Ferritic-martensitic steels, Aluminium diffusion coatings, Characterisation

Characterisation of aluminium diffusion coatings elaborated on austenitic stainless steels and on ferritic-martensitic steels

INTRODUCTION

In order to achieve higher efficiencies and to decrease the CO 2 emissions in the power generation industry, the operating temperatures and pressures have to be increased [1,2]. The new generation of ultra-supercritical power plants (USC units) operate with steam temperatures up to 600-650°C and pressures from 25 to 30 MPa [3,[START_REF] Nicol | Status of advanced ultra-supercritical pulverized coal technology[END_REF]. In these plants, austenitic stainless steels are commonly employed in superheaters and reheaters of steam boilers whereas ferritic martensitic steels are preferred for larger components like tubes and pipes [START_REF] Wang | [END_REF][6][7]. Due to their high temperature oxidation and corrosion resistance combined with their high strength, the HR3C and the IN-800HT austenitic stainless steels (ASS) can be employed in USC units [8,9]. With a density of 8 mg/cm 2 , IN-800HT is a modified version of IN-800 containing more Ni, Si, Mn and Ti [10,11]. This material is well-known to resist to stress corrosion and like for the HR3C steel, the presence of chromium in these alloys allows the formation of a protective oxide layer while the presence of nickel improves the stability of the oxide formed [8,10,12]. The nitrogen in the interstitial sites of the matrix contained in the HR3C alloy combined with 0.4Nb, 25Cr and 20Ni (wt%) additions represent an effective measure of the enhancement of creep strength [8,13]. The P92 ferritic martensitic alloy is one of the most commonly employed high Cr creep resistant steels in large components of USC units because of its high creep rupture strength, its excellent mechanical properties, its good weldability, its low thermal expansion, its high thermal conductivity and its adequate corrosion resistance [8,[14][15][16][17]. However, despite the good properties of these three alloys, they need to be coated in order to enhance their lifetime at higher temperatures and pressures. In severe atmospheres encountered in USC units, SiO can dissolve and Cr 2 O 3 can form volatile Cr oxy-hydroxyl species [18,19]. Thus, Al 2 O 3 is the only oxide ensuring the lifetime increase of the components. This external protective Al 2 O scale can grow from Al diffusion coatings [20]. In particular, slurry aluminising offers the possibility of coating large components [21] and is considered as low cost and easy-to-apply process [22]. Slurry coatings can be applied to a metallic substrate by brushing, spraying or by dip-coating [23]. Slurries are suspensions of aluminium micro-particles on a binder-solvent system [22]. In our case, the solvent is water, thus making our coatings environmentally friendly [24]. Once the Al slurry has been sprayed, the subsequent thermal treatment pyrolyses the binder and makes the molten aluminium to react with the substrate to form an aluminide intermetallic phase [21,23,25]. These coatings have shown their excellent resistance even after 2000 hours in steam at 650 or 700°C, for respectively P92 and IN-800HT and for the HR3C alloy [25][26][27]. Experiments conducted by Agüero et al. have shown the good behaviour of another kind of aluminium slurry deposited on P92 ferritic-martensitic steels for 41 000 hours in steam, at 650°C [28]. In addition, they can be repaired [29] and the potential cracks can be sealed [30]. However, although the high temperature behaviour of these aluminides has been discussed, there is no real description of the slurry coatings themselves in the open literature, let alone in what the aluminising treatment may affect the microstructure and microhardness of the parent material, which are the main aims of this work.

EXPERIMENTAL PROCEDURE

1 -Materials of Study

IN-800HT coupons of 16 mm diameter and 2 mm thick were ground with SiC#180 paper. The same procedure was conducted with HR3C and P92 coupons of 20*10*2 mm. The nominal compositions of the three studied substrates are given in Table 1. Aluminium micro-particles (4-5 µm average size, Hermillon, France) were used to aluminise the substrates. In order to obtain a more precise particles size and their composition, the powders were microscopically analysed. Characterisation of the Al powders and of the slurry were performed by DSC measurements, using a SETARAM Labsys Evo 1600 thermal analysis.

The heating and cooling rates of these analyses were 10°C.min -1 and the tests were performed under argon (20 mL/min). The Al microparticles were heated up to 1300°C and up to 700°C

for the slurry since after the melting of Al no significant event is recorded for the slurry [31].

2 -Aluminisation

All the samples were coated by spraying aluminium slurry containing 43wt% Al micro-particles dispersed in 57wt% of binder (1/10 polyvynil alcohol/deionized water). After drying in ambient air for at least 1 hour [31], the coated samples were annealed in argon (400°C/3h + 650°C/3h + 1000°C/1h or 1100°C/1h for, respectively the ferritic-martensitic steel and the austenitic stainless steels). These coating conditions were based in previous investigations on nickel-based superalloys [23,32] but were slightly lowered in temperature and shortened in time for these Fe-based alloys.

3 -Vickers microhardness and evalution of the grain size

Microhardness values were measured across the additive layer and the interdiffusion zone of the coatings and in the substrate for each coated material. The acquisitions were performed using a Shimadzu Type M tester with a Vickers indenter, with a load of 15 g and for 10 seconds.

The average values of microhardness were calculated from 10 indentations performed for each part of the coating/substrate system [33].

The average grain sizes of the coatings and of the steel core materials were evaluated using the ASTM E112 standard test method [34]. The intercept procedure was selected and involves an actual count of the number of grain boundary intersections with a test line per unit length of test line. For each system, 6 lines were plotted in 2 spatial dimensions in order to get a better average.

4 -Crystal Structures Before and After Aluminisation

The crystal structure was determined by X-Ray Diffraction (XRD) using a Bruker AXS D8

Advance (CuK α radiation), in θ-2θ mode.

5 -Etching

In order to obtain informations about the distribution and the size of the grains, the samples were mounted in resin, polished and etched before microscopic characterisation. The P92 samples were immersed for few seconds in a solution of Nital 5%. For aluminised IN-800HT, an electrochemical etching was realized by immersion for 3-5 seconds in a 10 vol% H 2 SO 4 aqueous solution. For the uncoated IN-800HT and for the uncoated and aluminised HR3C, chemical etching was realized with a Kalling n°2 solution (5g CuCl 2 + 100mL HCl + 100mL ethyl alcohol).

6 -Microscopic Characterisation

The etched cross-sections of the uncoated/coated samples were analysed using a Leica DMRM microscope and a LEICA MC170HD camera coupled with the software LAS v4.9 for the images acquisition. All the samples were also characterised by scanning-electron microscopy (SEM) in a FEI Qanta 200F FEG with an EDAX detector for chemical analyses.

RESULTS

3. 1 -Aluminium particles and corresponding slurry Preliminary analyses were carried out on the Al micro-particles in order to obtain an average size. Fig. 1. a) highlights the relatively heterogeneous distribution of the particle size. Indeed, some particles can measure up to 12 microns while the smallest ones do not exceed 0.5 microns. The Gaussian plot of Fig. 2. b)

shows that average size of the Al particles used in this study is about 3 microns. Oxygen was detected in small quantities (≈ 5-8 at%) and Ar was also quantified to less than 1 at%, which arise from the wire explosion method employed to produce the particles [START_REF] Mollard | Elaboration de systèmes barrière thermique par barbotine. Comportement du nickel et de ses superalliages revêtus en oxidation cyclique à haute température[END_REF][START_REF] Rannou | Slurry coatings from aluminium microparticles on Ni-based superalloys for high temperature oxidation protection[END_REF]. In contrast, no other elements were detected. In order to clarify this point, DSC measurements were carried out with the Al micro-particles. The corresponding thermogram shown in Fig. 2. a) let appear an endothermic event at 657 ± 1°C which is below the theoretical melting point of Al (660.5°C [START_REF] Okamoto | [END_REF]). This difference of melting temperature is probably due to the presence of impurities in the powders not highlighted with the EDS analyses [38,[START_REF]Physical Metallurgy and Processes, Handbook of Aluminum[END_REF]. Thereafter, the evolution of the transformation of the slurry was characterised by DSC up to 700°C, at which full decomposition of the binder and melting of aluminium are achieved [START_REF] Montero | [END_REF]. The slurry thermogram shown in Fig. 2. b) reveals four phenomena. The first phenomenon occurs below 100°C and was associated with the free water contained in the slurry [31]. The second event taking place at 100°C is endothermic and was ascribed to the evaporation of water.

The third phenomenon appears between 200 and 300°C but is less marked than the previous ones. Rannou et al. ascribed such phenomenon to the PVA, including the free water trapped in the PVA film, its subsequent melting and its final decomposition [31]. A final but important endothermic signal was noticed at 657 ± 1°C similar to the one of Fig. 2. a) for pure Al.

2 -Aluminide coatings

As shown in Fig. 3, the microstructure and the composition of the coatings appeared very homogeneous in the three steels. Despite of a good homogeneity in terms of microstructure, compositions and thicknesses of the coatings, some cracks were detected on austenitic stainless steels. As reported by our group, [27,29] the number of cracks in aluminised HR3C was a little bit greater than in the IN-800HT case due to the most important stiffness of the former [START_REF] Viswanathan | [END_REF]. Previous works have linked these cracks to the mismatch between the thermal expansion coefficients of the different metallurgical phases [28,44,45,46]. Other authors have reported this phenomenon with tensile stress from the substrate occurring upon heating [47].

3 -Metallurgical evolutions induced by the aluminisation process

As shown in Figs. 7. a) and a'), the thermal treatment seemed to reduce the grain size of IN-800HT. Indeed the average grain size of the uncoated and the aluminised substrates were, respectively, 219.5 ± 39.2 µm and 134.2 ± 11.7 µm (Table 2). For the HR3C substrate, it appeared that the grain size had slightly increased after aluminising while different carbides were also observed. Some square-like precipitates were found in the bulk of IN-800HT and were associated with TiC [48]. Moreover, some Cr carbides were detected in both steels as well as Nb-rich carbonitrides [13] and also M 23 C 6 which are typical of the ASS after annealing at high temperatures [9,13]. These precipitates were mainly located within the matrix (grains) of the untreated ASS but appeared with a round shape at the grain boundaries and with a cuboidal morphology within the grains after aluminisation [49] as shown in Fig. 7. a') and b'). By referring to the Schaeffler diagram [50] shown in Fig. 8 the structure of the HR3C evolved with the aluminising step at high temperature while that of In-800HT remained unchanged. The change in HR3C could be related to the greater Al and Cr contents underneath the interface between the substrate and the coating that provokes an evolution from 100% austenite to 95% austenite + 5% ferrite. However, the XRD diffractogram obtained on this sample let not appear any ferrite phase probably because of unsufficient penetration of the X-rays to this interface.

Therefore, the presence of the other minor phases besides austenite cannot be excluded. For the P92 substrate, the untreated material (Fig. 7. c) exhibited a tempered martensitic structure, as reported by Bates et al. for Grade 91 ferritic-martensitic alloy [20]. M 23 C 6 carbides characteristic of this type of alloy [START_REF] Hahn | Vallourec report "Pipe steels for modern high-output power plants[END_REF] were found in the P92 substrate as shown in Fig. 7. c).

By referring to the CCT diagram for the P92 substrate (Fig. 9), it appeared that the structure of the steel was converted to austenite after the aluminisation carried out in this study. The annealing step also resulted in a 10-fold increase of grain size. For the HR3C substrate (Fig. 10. b), the gap between the two values was smaller (252.8 ± 13.4

4 -Microhardness

and 235.8 ± 13.6 HV for, respectively, the uncoated and the aluminised specimens). On the contrary, the difference of Vickers microhardness between the aluminised (401.2 ± 26.3 HV) and the uncoated (217.1 ± 28.4 HV) P92 substrates was quite significant (Fig. 10. c).

The Vickers microhardness of the slurry aluminide coatings on IN-800HT, HR3C and P92 were also evaluated (Figs. 10). For the ASS, it appeared that the additive layers were harder than the interdiffusion zones and than the core of the substrates. In contrast, the opposite occurred in the aluminised P92 since the coating appeared softer than the core of the substrate (Fig. 10. c).

DISCUSSION

1 -Aluminium particles and corresponding slurry characterisation

The DSC analyses carried out with the Al microparticles and the corresponding aluminium slurry allowed to confirm the choice of the thermal treatment conducted in this study. Indeed, with four phenomena taking place below 400°C, the first ramp of the annealing has to be conducted slowly in order to achieve a gradual decomposition of the PVA. Then, a second heating rate up to 650°C followed by a 3h dwell at this intermediate temperature involves the Al melting and the subsequent sintering of the particles. Also, the final dwell of the annealing has to be carried out at higher temperature in order to promote the Al inward and the simultaneous Ni, Fe outward diffusion to create the aluminide coating [23,32]. The DSC measurements realized in this study confirm thus previous works realized in our group although the Al particles (supplier and size) were different [25][26][27]29].

-Aluminide coatings

The IN-800HT, HR3C and P92 substrates were aluminised using the slurry analyzed in the first part of this study. The EDS profiles (Figs. Therefore, a significant increase of the Cr content was observed below the interface between the additive and the interdiffusion layers in both austenitic stainless steels. Moreover, in the HR3C steel, Cr-rich precipitates were highlighted in the interdiffusion zone and again, Cr was detected while Ni was not. Conversely, Ni was detected in the interdiffusion zone in the Crpoor areas. Carbides were also found within the interdiffusion zone of the IN-800HT substrate. Some elongated precipitates were observed in the HR3C case and were associated to AlN [28,44,55,56].

Overall, the microstructure and the composition of the two coatings on the ASS are very close to those obtained on AISI 304 steel by chemical vapour deposition (CVD) in fluidised bed reactors [33,57]. However, the thermal treatments carried out (Ar (g)) in this work were shorter because of the combustion synthesis occurring between Al and the substrate, thus forming an aluminide, as described by Montero et al. [21] and Galetz and collab. [23]. Indeed, the melting of at least one of two components (Al in this case) is a crucial step because it is the beginning of the intermetallics formation [23,58]. Previous studies have shown the high exothermic reaction occurring between Fe and Al in argon below the melting temperature of Al (i.e. 660°C) [21,59]. Therefore, the intermediate step of the aluminisation is set at 650°C, as also performed in [25][26][27]. A 3h dwell at this temperature was sufficient because of the fast Al enrichment and the simultaneous Fe and Ni dissolution [23]. Such dissolution brings about the formation of the continuous diffusion layer shown in Figs. 3. a) and b). The cracks detected in the additive layers of the austenitic stainless steels were attributed to the difference existing between the different metallurgical phases. Indeed, when a gradient of thermal expansion coefficient exists between the coating and the substrate, the force generated can physically break the system [START_REF] Nicol | High temperature steels in pulverised coal technology[END_REF].

However, these cracks do not propagate into the substrate, thus not jeopardizing the protection of the latter.

The aluminium diffusion coating elaborated on the ferritic-martensitic P92 steel was thicker than those obtained in the austenitic stainless steels. Nevertheless, the additive layer was also mostly composed by the B2-FeAl phase as highlighted by the EDS profiles (Fig. 3. c) and the X-ray patterns (Fig. 5. b). Needle-like precipitates were observed in the interdiffusion zone of the coated P92 and were attributed to AlN [20,28,45]. As reported by Agüero et al. [START_REF] Agüero | [END_REF]62],

these precipitates are due to the N content in the substrate (0.046 wt%). Porosities were also detected at the coating/substrate interface and could be attributed to a Kirkendall effect due to the formation of chromium aluminides impeding locally further Al inward diffusion [54,63].

Otherwise, Al was detected in the substrate until about 200 µm from the coating surface to the substrate because the diffusion through a ferritic structure is relative fast [53,64]. No cracks were detected in the coated P92 because the gap between the substrate and the B2-FeAl phase is not probably significant enough.

3 -Metallurgical evolutions induced by the aluminisation process

A decrease of the grain size of the IN-800HT was noticed after aluminisation (219.5 ± 39.2 µm and 134.2 ± 11.7 µm, respectively for the uncoated and the aluminised substrate). In the HR3C case the evolution was less marked and the grain size was almost comparable. For the two austenitic stainless steels, the morphology observed after the thermal treatment (Figs. 7. a') and b') is typical of an austenitic structure [65]. Some square-like precipitates were found in the bulk of IN-800HT and the EDS analyses carried out let appeared high Ti and C contents and can be thus associated with TiC [45]. Moreover, some Cr carbides were detected and have been reported to confer high temperature strength to this alloy [11]. Segregation of M 23 C 6 at the grain boundaries of the two steels was highlighted and is typical of the ASS after annealing at high temperatures [9,13]. Before aluminisation, these precipitates were revealed in the matrix.

By correlating EDS analyses and the Schaeffler diagram, it appeared that the composition underneath the interdiffusion zone/substrate interface did not evolve in the IN-800HT case. On the contrary, an evolution from the austenitic domain to 95% austenite + 5% ferrite was noticed in HR3C. Al and Cr are ferrite stabilizers and could thus result in a 5-10% ferrite transformation in the substrate. Aluminisation through the ferritic phase is known to occur more readily than through the austenitic phase [53,64]. However, Cr is also a diffusion barrier for Al as CrxAly intermetallic compounds block the inward Al diffusion [29]. Therefore, the greater Cr content substrates are similar to those retrieved by EDS before aluminisation. It appears that the overall thickness of the Al-modified zone is negligible compared to the entire thickness of the system (coatings + substrates: about 2 mm). Therefore, the cracks may appear because of the greater stiffness of the intermetallic compounds (in particular in the diffusion layer) than that of the substrate.

For the P92 substrate, the microstructure evolved in a larger extent during the aluminisation.

Indeed, the initial tempered martensitic structure related by Bates et al. [20] converted to austenite one considering the CCT curves of Fig. 9. In order to recover the initial microstructure of this steel, after a rapid quench in air a subsequent annealing shall be performed at 760°C for 1 hour [66]. Nevertheless, the aluminisation realized in this study was done in a tubular furnace which is not able to cool down quickly. With our thermal treatment, the initial tempered martensitic structure converted -at least partly-to austenite when the samples were heated at 1050°C for 1 hour. Moreover, Wawszczak et al. [66] showed that the microstructure of a ferritic-martensitic steel (compositions close to those obtained for P92) does not change significantly after annealing at temperature below 400°C. In this range of temperatures, the Al coating cannot be formed. The increase of the temperature allows the Al melting and, therefore the wet of the substrate allowing the start of the combustion synthesis process [21] which will enable the Fe enrichment of the coating. By increasing the temperature up to 1050°C, the intermetallic phases created during the heating ramp become homogeneous to form and to maintain the B2-FeAl phase. Moreover, the relatively slow cooling rate to room temperature results in the precipiptation of M 23 C 6 carbides that could have been partly dissolved and the martensitic microstructure recovered upon tempering at 760°C [START_REF] Hahn | Vallourec report "Pipe steels for modern high-output power plants[END_REF]. In contrast, quenching from the high coating temperature (1000 or 1050°C) may introduce mechanical failure of the coating through the ductile-to-brittle transition temperature (DBTT) [START_REF] Ennis | Czyrska-Filemonowicz[END_REF]. The thermal treatment carried out in this study had an other consequence corresponding to the increase of the grain size. Indeed, it appeared that the grains were about 10 times bigger compared to those of the uncoated P92 samples. These observations are in agreement with the findings of Ellis et al. who have reported that the austenite grain size evolves as a function of the austenitization temperature; the higher the austenitization temperature and the bigger the grain size [START_REF] Ennis | Czyrska-Filemonowicz[END_REF]. Such increase of the grain size during the heating of ferritic-martensitic has been also reported by Klueh and Shingledecker [START_REF] Klueh | Investigation of a modified 9Cr-1Mo (P91)[END_REF].

4 -Microhardness

Generally speaking and within uncertainty of the measurements, the slurry aluminising process studied here seemed to soften slightly the matrix of the austenitic stainless steels. Nevertheless, a decrease in term of microhardness has also observed in [33] after a heat treatment due to the dissolution of carbides shown when comparing the microstructures of Fogs 7a) and 7b) (uncoated) with 7a') and 7b') (aluminised). For the P92 substrate, the evolution of the microhardness was non-negligible. Indeed, after aluminisation, the ferritic-martensitic steel appeared two times harder and a significant evolution of microstructure from a martensitic to an austenitic+carbides. As shown in Fig. 9 [51] and as reported by Hahn et al. [START_REF] Hahn | Advanced Heat Resistant Steels for Power Generation[END_REF], the P92 substrate should have been cooled sufficiently fast after the austenitization (1050°C) and subsequently tempered. The significant change of the microstructure of the P92 material before and after the aluminisation explain the more pronounced microhardness evolution compared to the ASS.

The changes in composition and in microstructure also affected the Vickers microhardness of the slurry aluminide coatings elaborated on IN-800HT, HR3C and P92 (Fig. 10). For the ASS, it appeared that the additive layers were harder compared to the corresponding interdiffusion zone and substrate. The same trend has been reported for another austenitic steel (AISI 304) which was previously aluminised by CVD-FBR where comparable values of about 350 HV were reported [33]. Moreover, the additive layer of IN-800HT was harder (700 HV) than that of HR3C (480HV). By comparing these results with the EDS profiles shown in Figs. 3, it appeared that the Ni content in the additive layer was higher for the IN-800HT (24 at%) compared to the one noticed in the HR3C case (10 at%). It is known that nickel accelerates the precipitation by diminishing the solubility of carbon [START_REF] Mc Guire | ASM International[END_REF] allowing an increase in term of microhardness [13,[START_REF] Dehmolaei | [END_REF]. In both cases, the microhardness noticed in the IDZ was smaller compared to those obtained in the additive layer, probably due to a Ni content decrease. The absence of Ni in the aluminide coating of P92 can thus also explain why the coating is much softer than that of the ASS substrates and of the P92 steel. The aluminisation of the ASS did not significantly change the grain size and the microhardness of the substrate whereas the thermal treatment had a non-negligible impact on the microstructure of the ferritic-martensitic substrate. Indeed, the initial tempered martensitic structure have been converted to austenite + carbides due to the absence of a final quench in air until 760°C. Thus, the grain size appeared 10 times larger after the aluminisation process. The coated P92 specimens also assumed a large increase in term of microhardness because of the precipitation of M 23 C 6 carbides along the grain boundaries during the annealing step.

FIGURE 1 .

 1 FIGURE 1. Micrographs of Al micro-particles used in this study (a) and their size distribution (b).

FIGURE 2 .

 2 FIGURE 2. DSC curves upon heating (10°.min-1) under Ar(g) flow of (a) the Al microparticles and (b) of the Al slurry.

FIGURE 3 .

 3 FIGURE 3. SEM cross-sections of aluminised (a) IN-800HT, (b) HR3C, (c) P92 steels and their corresponding EDS profiles (d, e and f, respectively).

FIGURE 4 .

 4 FIGURE 4. (a) Fe-Ni-Al diagram [41] and (b) Fe-Al diagram [42]. Compositions of the additive layer of IN-800HT (yellow circle), HR3C (green circle) and P92 (orange circle).

FIGURE 5 .

 5 FIGURE 5. X-ray patterns of the (a) IN-800HT and HR3C austenitic stainless steels and (b) of the P92 ferritic-martensitic steel before and after aluminisation.

FIGURE 6 .

 6 FIGURE 6. X-ray maps of the slurry aluminide coatings on (a) IN-800HT, (b) HR3C and (c) P92 steels.

FIGURE 7 .

 7 FIGURE 7. Optical micrographs of uncoated (a) IN-800HT (Kalling n°2), (b) HR3C (Kalling n°2), (c) P92 (Nital 5%) and aluminised (a') IN-800HT (10% vol H 2 SO 4 ), (b') HR3C (Kalling n°2), (c') P92 (Nital 5%).

FIGURE 8 .

 8 FIGURE 8. Schaeffler diagram [50]. Structures of uncoated/aluminised IN-800HT (red points) and HR3C (blue points).

FIGURE 9 .

 9 FIGURE 9. CCT diagram for P92 ferritic-martensitic steel [51].

Figure 10

 10 Figure 10 summarizes the Vickers microhardness values measured before and after aluminisation. It appeared that the microhardness at the core of the ASS before aluminisation was higher compared to those obtained after aluminisation but the gap was negligible in view of the uncertainties of the measures. Before the annealing step, the microhardness of IN-800HT was 210.8 ± 38.7 HV whereas after aluminisation it was about 169.1 ± 11.7 HV (Fig. 10. a).

FIGURE 10 .

 10 FIGURE 10. Vickers microhardness of uncoated and aluminised (a) IN-800HT, (b) HR3C and (c) P92 substrates.

  3. a) and b) and the X-ray patterns (Fig. 5. a) show that the additive layers elaborated on IN-800HT and HR3C are mostly composed by the B2-(Fe, Ni)Al phase, which is indicative of inward Al diffusion into the steel matrix and, simultaneously, outward Fe and Ni diffusion. In both cases, the coatings appeared homogeneous in terms of thicknesses and composition. Although the thickness of the additive layers was the same, the interdiffusion zone of the HR3C substrate was thicker than the one on IN-800HT (respectively 49 and 19 µm). Previous works have related this phenomenon to the greater amount of Cr in the HR3C[25][26][27][28][29]. Indeed, during the aluminisation, Cr is incorporated in the Al coating, thus preventing the Al inward diffusion and the outward diffusion of the compounds contained in the substrate[26,[START_REF] Cheng | [END_REF]. The Al distribution within the additive layers of the ASS appeared homogeneous whereas a Cr and Fe depletion was observed close to the additive/interdiffusion zones interface (Fig.6. a) and b). On the contrary a Ni enrichment was observed in these areas. Cr is known to have a very low solubility in Al-rich and stoichiometric aluminides thus explaining the presence of Cr in the absence of Ni and vice versa[53,54].

  in the interdiffusion zone blocked Al inward diffusion and the additive layer became thus more Al-rich in HR3C than in IN-800HT. Conversely, Fe diffused outwardly through the aluminide layers to stabilise at about 35 at% of the B2-FeAl phase according to the Al-Fe phase diagam of Fig. 4. b). Nevertheless, the compositions at the core of the IN800HT and the HR3C

  Slurry aluminide coatings were elaborated on IN-800HT and HR3C austenitic stainless steels and on P92 ferritic-martensitic steel. All the coatings were composed by an outer layer corresponding to the protective B2-(Fe,Ni)Al and B2-FeAl phase, for, respectively, the ASS and the ferritic-martensitic steel, and an inner interdiffusion layer. The thermal treatments of the three substrates were carried out in Ar and allowed the Al melting involving the combustion synthesis with the substrate. Thereby, the simultaneous Al inward diffusion in the steel matrix and the outward Fe (and Ni for the ASS) occurred forming a continuous and consistent diffusion layer mostly composed by B2.

TABLE 1 .

 1 Nominal and experimental (EDS) compositions of IN-800HT, HR3C and P92 substrates (wt%).

			Fe	Ni	Cr	C	Ti	S	Others
									Si: 1 max
	IN-800HT	Nominal	Bal	30-35	19-23	0.10 max 0.15-0.60 0.015 max	Mn: 1.5 max Cu: 0.75 max
		EDS	45	29	21	< 1	< 1	-	Si: 1-2
									Si: 0.4
	HR3C	Nominal	Bal	17-23	23-27	0.06	-	-	Mn: 1.2 Nb: 0.20-0.60 N: 0.15-0.35
		EDS	53.5	20	25	< 1	-	-	Mo: 0.4 Nb: 0.5
									Si: 0.50
									Mn: 0.30-0.60
		Nominal	Bal	0.4 max 8.50-9.50 0.07-0.13	-	0.01 max	Mo: 0.30-0.60
									W: 1.5-2
	P92								V: 0.15-0.25 Mn: < 1
		EDS (wt%)	87	-	9	-	-	-	Mo: <1
									W: 2

TABLE 2 .

 2 Grain sizes of the uncoated and aluminised substrates studied in this work.

	IN-800HT austenitic stainless steel	HR3C austenitic stainless steel	P92 ferritic-martensitic steel
	Uncoated	Aluminised	Uncoated	Aluminised	Uncoated	Aluminised
	219.5 ± 39.2 µm	134.2 ± 11.7 µm	78.8 ± 11.8 µm 93.5 ± 17.2 µm	4.8 ± 2.3 µm	51.8 ± 8.2 µm