Florian Cuminet

Dimitri Berne

Sébastien Lemouzy

Éric Dantras

Christine Joly-Duhamel

Sylvain Caillol

Éric Leclerc

Vincent Ladmiral

Catalyst-free transesterification vitrimers: activation via αdifluoroesters

Transesterification vitrimers often require high catalyst loadings to achieve 3D networks reprocessable at moderately high temperature. The addition of an activating group close to the ester bonds allows to synthesize catalyst-free transesterification vitrimers. Here, we unveil the effect of the α-difluoromethylene group as a novel activating group for such materials. Fluorine features exceptional properties, in particular a strong electronegativity enabling CF2 groups to activate the epoxy-acid polymerization, and more interestingly also the transesterification reaction on adjacent esters. Consequently, this fluorinated group affords the easy synthesis of a highly crosslinked reprocessable material that do not require any metallic or organic catalyst. This vitrimer is endowed with advantageous reprocessing abilities and underwent 10 consecutive cycles without loss of mechanical properties. In brief, the vitrimer combines durability, recyclability and is catalyst-free. This discovery is one step further towards recyclable greener polymers.

Introduction

For decades, polymer materials have been classified depending on their macromolecular structure. On the one hand are amorphous thermoplastics, composed of entangled linear chains. In these materials, polymer chains are not covalently bonded, and are able to slide on one another. 1 This translates, at the macroscopical scale, in their solubility in suitable solvents and in their ability to be reshaped upon heating. On the other hand are thermosets, composed of a 3D network. Theoretically, a thermoset is a single infinite macromolecule forming a covalently crosslinked network. This permanent structure leads to insolubility in solvents, and prevents recycling by simple thermal methods. 1,[START_REF] Pascault | Thermosetting Polymers[END_REF] In the pursuit of greener, more sustainable polymers, thermosets remain a challenge for recycling. [START_REF] Pickering | Wiley Encyclopedia of Composites[END_REF] Yet, they are assets in many applications for which thermal and chemical resistances are required. A third class of polymer materials consisting in a covalent crosslinked network in which the covalent bonds can be exchanged by reversible chemical reactions was proposed. They were called CANs, short for Covalent Adaptable Networks. In 2005, Bowman et al. [START_REF] Scott | [END_REF] synthesized a covalently bonded 3D network alike thermosets, endowed with covalent bonds able to exchange via an associative mechanism upon irradiation with visible light. This concept was further developed by Leibler et al. 5 in 2011 with a material based on transesterification activated upon heating. This material was insoluble in common organic solvents even at high temperature, but able to flow and to be reshaped upon heating. Interestingly, this new kind of material exhibited an Arrhenian viscosity decrease with increasing temperature, a feature usually observed for inorganic glasses, or more generally for "strong" glasses in accordance with the work of Angell. 6 This analogy inspired the name vitrimers for such materials. Since this discovery, other kinds of exchangeable bonds such as carbonates, 7 boronic esters, 8,9 disulfides, 10,11 silyl ethers, 12,13 urethanes, [14][15][16][17][18] olefins, 19 imines, [20][21][22][START_REF] Duval | Av rous[END_REF] vinylogous urethanes, [24][25][26] diketoenamines, 27 acylated acetals, 28,29 and trialkylsulfonium salts 30,31 have been investigated. Despite the success of vitrimers in materials research, they have yet to be implemented in industrial and commercial applications as they are not adapted to the recycling processes existing for thermoplastics. 32 Moreover, long reprocessing times at high temperatures can trigger premature degradation of the vitrimer after repeated reshaping processes. 33 To accelerate the reshaping process, one solution is the use of catalysts. However, some exchange reactions such as transesterification are so slow that they often require high loadings. 5 Yet, the use of catalysts raise concerns about the materials ageing or the risk of leaching, 34,35 which is undesirable for application such as food-contact materials for example. The first catalyst-free vitrimer was based on vinylogous urethanes exchange. This reaction is fast enough to afford short reprocessing times without catalyst. 25 Since this seminal article, other examples of catalyst-free vitrimers based on fast exchange reactions were proposed, such as hydroxyurethanes, 16 hemiacetal esters, 28 trialkylsulfonium salts, 31 oxime-esters 33 and imines. 36 Lowering the crosslinking density and increasing the number of exchangeable moieties were proven to enhance the reshaping abilities of catalyst-free materials. [37][38][39][40][41] However, since these parameters also influence the materials mechanical properties, this strategy may not be suitable to a wide range of applications. To overcome the risk of catalyst leaching, some vitrimers were also synthesized with a catalyst (amines for instance) embedded in the network.

42-44

In 2015, Guan et al. proposed to implement in vitrimers the concept of neighboring group participation (NGP), well-known in the field of organic chemistry but new to the field of vitrimers.

45 Materials crosslinked using difunctional dioxaborolane and possessing neighboring amino groups relaxed faster than the one deprived of such groups. The potential of internal catalysis or internal activation and neighboring group participation (NGP) to tune CANs has recently been reviewed. 46,47 In their review, Van Lijsebetten et al. 46 made a clear distinction between NGP for which the neighboring group is involved with a covalent bond at some point of the reaction, and the broader concept of internal catalysis including effects such as inductive effects. The present work focuses on transesterification vitrimers, as they usually require a catalyst. To address the concerns raised by the use of external catalysts, some functional groups were reported to be efficient for NGP, such as phthalate monoesters 48 or benzenesulfonic acid groups. 49 Such groups modify the transesterification mechanism and facilitate the exchange. Examples of activation by inductive effect were reported on Meldrum's acids in PDMS 50 or in polyimine vitrimers for instance. 51 In transesterification vitrimers, malonates were used to activate the exchange reaction.

52

Because of its high electronegativity, fluorine has a strong potential to activate bond exchanges. Indeed, the carbonyl group of fluorinated esters is known to be very electrophilic, 53 which facilitates hydrolysis and nucleophilic attack.

54,55

Fluorine substitution thus appear as a good strategy to activate transesterification reactions. Recently, a CF 3 group positionned on the α carbon of esters was demonstrated to have a significant activating effect on transesterification in polyester networks. [START_REF] Berne | [END_REF] In the present work, fluorine atoms were added one atom closer to the ester bond and α-diflluoro esters are implemented for the first time to design a catalyst-free transesterification vitrimer. A trifunctional α-difluoro carboxylic acid was prepared and used in combination with a commercial diepoxide (butanediol diglycidyl ether, BDGE) to prepare an epoxy-acid network which displayed insolubility in organic solvents, but was able to be reshaped under relatively mild conditions without catalyst.

Results and Discussion

α-Difluoro Carboxylic Acid Monomer Synthesis

Transesterification vitrimers require dangling hydroxy groups for the exchange reactions to happen. Fortunately, the epoxy opening by a carboxylic acid produces such groups in beta position to the ester. The strategy developed here therefore relies on the use of a trifunctional carboxylic acid monomer, and a commercial diepoxide to obtain a 3D network. To maximize the inductive effect of fluorine on the ester bond, the strongest fluorinated activating group (CF 2 ) should be positioned as close as possible to the exchangeable bond. A trifunctional monomer bearing three α-difluorocarboxylic acids was thus designed and synthesized in two steps from a commercial triphenol (1,1,1-tris(4-hydroxyphenyl)ethane, "TPE"). TPE was reacted with ethyl bromodifluoroacetate at 70 °C for 40 hours to obtain by nucleophilic substitution a mixture of the di-and trisubstituted esters (15 and 60 mol% respectively). After column chromatography, the resulting αdifluoro triester then underwent facile saponification to yield, after appropriate workup, the desired α-difluoro triacid (TPE-TAF) as a white waxy solid with an overall yield of 50 % (Scheme 1).

Scheme 1. 2-step synthesis of the trifunctional α-difluoro carboxylic acid TPE-TAF

Polymerization and Curing

Butanediol diglycidyl ether (BDGE, ESI Figures S14 and S15 and

Sections A and B) was identified as a promising commercial epoxy resin, as its viscosity is very low and because it can dissolve TPE-TAF before gelation happens. In addition, its structure advantageously adds flexibility to the network, may counterbalance the rigidity of the TPE-TAF structure, and might prevent a high T g which would make the reprocessability of the material more difficult (Scheme 2). TPE-TAF was dissolved in BDGE at room temperature (ca. 20 °C). The mixing time needed to be reduced to a few minutes only, to avoid the gelation caused by the fast reaction between the α-difluoro acids and the epoxides (even at room temperature) and ensure the homogeneity of the mixture. The mixture was then left at room temperature for at least 8 h, resulting into a brittle gel, which was then cured at 150°C. The gel time observed visually was consistent with the gel time determined by rheological analysis (Figure S17). At 20 °C, a gel time of 1.3 h was estimated at the crossover of loss and storage moduli, confirming that the epoxy opening by the α-difluoro carboxylic acid readily happens at room temperature. This behavior is highly unusual for epoxy-acid systems, which usually require catalysis and relatively high temperatures (80-120 °C). 57 For catalyst-free systems, even higher temperatures (around 200 °C) are required. 58 This fast polymerization reaction at low temperature strongly suggests the activating inductive effect of fluorine atoms on the carboxylic acid towards the epoxy-opening reaction.

The polymerization kinetics at room temperature were monitored via FTIR. The disappearance of the epoxy band at 908 cm -1 was clearly observed but the conversion of acid functions at 1758 cm -1 into esters at 1761 cm -1 could not be determined, because the acid and ester bands overlapped too much (Figure S22). The epoxy band decreased significantly for the first 6 h of reaction, and then decreased at a much lower rate until the end of the acquisition after 64 h (Figure 1). This rate change was probably due to the gelation of the mixture which slowed down the reaction between the remaining reactive species.

The material obtained after 4 days at room temperature was left for 3 h in an oven at 150 °C to ensure complete polymerization. Three consecutive DSC ramps were carried out on the resulting material. The thermograms overlaid perfectly, did not show any residual exotherm and revealed a T g of 47 °C which did not evolve after any of these heating ramps (Figure S18). The curing of the material was thus deemed complete. Solubility tests were performed in acetone, THF, toluene, cyclohexane, DMSO, DCM and acetonitrile (Table S1). The highest values were found for acetone and THF. In particular, after 24 h under agitation, an insoluble fraction of 94 ± 2 % was measured in acetone (the best solvent for BDGE and TPE-TAF), thus proving that the curing process led to the formation of a 3D crosslinked material, as expected.

Vitrimer Characterization

TGA under air showed that no significant mass loss was observed up to 200 °C, with a 2 % and 5 % degradation temperatures T d2% and T d5% of 205 °C and 262 °C respectively (Figure S20 and Table 1). The T g of the TPE-TAF/BDGE material was determined to be 47 °C (Figure S18 and Table 1). TPE-TAF contains three aromatic cycles which bring rigidity to the polymer structure, whereas the linear structure of BDGE adds flexibility. This balance explains the moderate T g value observed.

A few preliminary reshaping trials allowed to set the reshaping temperature value at 100 °C. The material stability to reprocessing cycles was then determined by isothermal TGA experiments performed at 100 °C for 16 h under air, to simulate the oxidative environment during reprocessing (Figure 2 and Table 1). After 4 h, a mass loss of 4.3 % was observed. The value stabilized to 4.4 % after 5 h and did not The kinetics of the flow behavior of the material was studied using stress-relaxation experiments. The relaxation modulus was monitored with time between 170 and 210 °C with 10 °C steps (Figure 3). It is important to state here that analogous non fluorinated polyester epoxy network prepared using 0. The material relaxation observed in the absence of external catalyst demonstrated the activating effect of the two βfluorine atoms on transesterification in the material, as noncatalyzed epoxy-acid networks usually exhibit no relaxation on a measurable time scale. As previously mentioned, tiny amounts of unreacted carboxylic acid, if present, could also add a contribution on this effect. A slight contribution from the phenoxy group is plausible, but would be much weaker than the contributions of the two fluorine atoms given the relative electronegativities of these elements (χ F = 4.2, χ O = 3.6). This is well illustrated by the by the pKas of acetic acid 4.7, glycolic acid 3.8, phenoxyacetic acid 3.2 and difluoroacetic acid 1.2.

63,64

This range stresses out the difference in electron-withdrawing ability of two fluorine vs one oxygen atom. The material behavior in temperature followed an Arrhenius law, as expected for a vitrimer. α-Difluoro esters are thus efficient activated esters for the design of catalyst-free transesterification vitrimers. The discrepancy between the sluggishness in stress-relaxation experiments and the mild reprocessing conditions observed is striking. This difference can be explained by the pressure applied to the material, 65,66 which is an important and often overlooked parameter. The force applied onto the material during reprocessing is 8.8 times the force applied during relaxation experiments. The reprocessing experiments are thus carried out at high pressure value in compression, which explains the difference in the behavior observed.

Reprocessing classical thermosets by compression molding is impossible, in contrast to vitrimers which possess exchangeable bonds. The TPE-TAF/BDGE material was successfully reprocessed using compression-molding further demonstrating its vitrimer character. 1 mm 3 pieces of the material were reassembled into a small ribbon after 1.5 h at 100 °C under a 6 ton load. The required reprocessing conditions were relatively 5,37,40,44,48,52,62,[67][68][69][70][71][72] mild compared to transesterification vitrimers reported in literature, whether they are catalyzed or not (Figure 4). Ten successive reprocessing cycles were successfully performed and each time homogeneous transparent samples were recovered. The color of the material did not significantly change with the successive reshaping cycles (Figure 5). To quantitatively assess the material evolution with reprocessing, thermomechanical analyses were performed. The TPE-TAF/BDGE material evolution with reprocessing was studied using DMA experiments. The values of the storage modulus E' in the glassy plateau and in the rubbery plateau regions, and the evolution of T α allow to detect whether the TPE-TAF/BDGE material mechanical properties decrease after several reprocessing cycles. A ramp from -100 to + 150 °C was performed after each reprocessing cycle to check whether the mechanical properties were fully recovered.

The pristine material exhibited a glassy plateau storage modulus E G ' = 3.7 GPa, a rubbery plateau storage modulus E R ' = 18 MPa and a T α = 39 °C, which is consistent with the values reported in the literature 37,70,[72][73][74] for epoxy-acid and epoxyanhydride vitrimers (Table 2 and Figure S19). Upon the first three reprocessing cycles, the T α increased from 39 °C to 41, 47 and 49 °C respectively, suggesting that the crosslink density slightly increased after each reshaping step, which is consistent with the increase of the rubbery plateau modulus (Table 2). After this 3 rd reshaping process, the T α value did not change with further reprocessing cycles up to the tenth cycle (Table 3). The storage modulus (E') value in the glassy plateau region slightly changed from 3.7 to 3.9 GPa for the pristine network and after the tenth reprocessing respectively. Similarly, the value of E' in the rubbery plateau region increased by a mere 0.6 MPa (Table 2) (from 18.1 to 18.7 MPa for the pristine and 10 th reprocessing cycle). These results show that there is no significant evolution of the network after 10 cycles.

Conclusion

In summary, polyester vitrimers were prepared out of a synthesized trifunctional α-difluoroacid and a commercially available difunctional epoxy resin. Thanks to the activation of the acid by the fluorine atoms, the epoxy-acid polymerization reaction happened readily at room temperature, whereas catalysts and high temperatures are usually needed for nonfluorinated systems. The resulting TPE-TAF/BDGE material was insoluble yet able to be reprocessed under mild conditions (100 °C, 1.5 h, 6 t). As expected, α-difluoro ester underwent significantly accelerated transesterification. This activation was such that no external catalyst was needed, in contrast to most other transesterification vitrimers. The flow activation energy E a was evaluated to be 77 kJ mol -1 which is consistent with the values reported for transesterification vitrimers catalysed by internal amines for instance, but also very close to the value (67-72 kJ mol -1 ) obtained for α-CF 3 activation. [START_REF] Berne | [END_REF] This polymer proved highly stable over repeated reprocessing cycles, with very little degradation of the mechanical properties observed after 10 cycles. The proof-of-concept based on the high electron withdrawing effect of fluorinated groups and exposed here is very promising to tune vitrimers properties and address the concerns related to the use of external catalysts, such as premature ageing, leaching of the catalyst or limited reprocessing abilities.

Synthetic Procedures "TPE-TE" compound:

1,1,1-Tris(4-hydroxyphenyl) ethane (6.13 g, 20 mmol, 1 equiv) was dissolved in dry DMF (120 mL, 0.16 M). 1,8-Diazabicyclo [5.4.0] undec-7-ene (DBU, 15 mL, 100 mmol, 5 equiv) was added in one portion and the reaction was heated to 70 °C. Ethyl bromodifluoroacetate (12.8 mL, 100 mmol, 5 equiv) was then added via a syringe pump at a rate of 5.0 mL h -1 and the reaction was stirred at 70 °C for 40 h. The resulting reaction mixture was very dark due to the probable oxidation of some of the phenol reagents, which could not be prevented even under argon atmosphere or protection from light. The crude mixture was cooled to room temperature, diluted with H 2 O (600 mL), and extracted 5 times with Et 2 O (5 x 100 mL). The combined organic layers were washed 2 x 150 mL with water and with 150 mL of brine, dried with Na 2 SO 4 , filtered, and concentrated under reduced pressure. The dark-brown crude mixture was purified by column chromatography on silica gel using pentane/ethyl acetate mixtures (5/1, 3/1 and finally 1/1) as eluent to afford the pure triester (F1: 8. In a 250 mL round bottom flask, 7 g of triester were dissolved in acetonitrile (90 mL). Then, a 5M aqueous solution of NaOH (12.5 g in 62 mL) was added slowly at room temperature, and the mixture was stirred for 3 h. 200 mL of a saturated NaHCO 3 solution was added to the mixture and the aqueous layer was washed with 100 mL of diethyl ether. The organic layer was extracted with 50 mL of saturated NaHCO 3 solution, and the gathered aqueous layers were acidified to pH=1 using 2M HCl. Finally, the acidified aqueous layer was extracted with 3 x 100 ml of diethyl ether, and the solvent was removed under high vacuum to afford the desired triacid as a white waxy solid (yield over the two steps η = 50 %, purity > 98 % estimated from

Determination of the Epoxy Equivalent Weight (EEW)

The EEW of the BDGE was evaluated by NMR titration using benzophenone as standard in deuterated chloroform (experimental details are given in ESI Section A). This value was confirmed by DSC studies using succinic acid as the curing agent (experimental details are given in ESI Section B). The procedure was described in a previous article. 76 Several BDGE/succinic acid ratios were used to make a series of thermosets. The highest T g was achieved for a 1:1 stoichiometry, from which the EEW was calculated.

"TPE-TAF/BDGE" vitrimer

Typically, 1.156 g (5.90 meq COOH) of TPE-TAF was quickly mixed manually in a 10-mL beaker with 0.682 g (5.93 meq epoxy) of BDGE at room temperature (ca. 20 °C) until the acid was fully dissolved. A clear yellowish viscous liquid was obtained and quickly cast in PTFE molds. The mixture was left at least 8 h at room temperature for gelation. The resulting material (TPE-TAF/BDGE) was then removed from the molds and cured 3 h at 150 °C. F were acquired on a Bruker Avance 400 MHz spectrometer at 23 °C. External reference was tetramethylsilane (TMS) with chemical shifts given in ppm. Samples were diluted in 0.5 mL of CDCl 3 or DMSO-d 6 depending on their solubility. FTIR: FTIR spectra and single wavenumber measurements were acquired on a ThermoScientific Nicolet iS50 FT-IR equipped with an attenuated total reflectance cell (ATR). The data were analyzed using the software OMNIC Series 8.2 from Thermo Scientific. Mechanical characterizations: Gel time experiment was performed at 1 Hz with a 20-mm plane-plane geometry on a ThermoScientific Haake Mars 60 rheometer equipped with a Peltier heating cell. A 30 mL h -1 nitrogen flux was applied.

Stress relaxation experiments were performed on the same apparatus equipped with a textured 8-mm plane-plane geometry. A 10 % torsional strain was applied on 8 mm diameter and 2 mm thickness circular samples, and the rubbery modulus evolution with time was monitored. DMA: Dynamic Mechanical Analyses were carried out on Metravib DMA 25 with Dynatest 6.8 software. Uniaxial stretching of samples (1 × 5 × 12 mm

3 ) was performed while heating at a rate of 3 °C min -1 from -90 °C to 150 °C, keeping the frequency at 1 Hz. TGA: Thermogravimetric thermograms were recorded on a TA TGA G50 instrument using a 40 mL min -1 flux of synthetic air as purge gas. Approximately 10 mg of sample were used for each analysis. Ramps from 20 to 500 °C were applied at a rate of 20 °C min -1

. DSC: Analyses were carried out using a NETZSCH DSC200F3 calorimeter. The calibration was performed using adamantane, biphenyl, indium, tin, bismuth and zinc standards. Nitrogen was used as purge gas. Approximately 10 mg of sample were placed in perforated aluminum pans and the thermal properties were recorded between -100 °C and the temperature of 2 % degradation T d2% at 20 °C min -1 . The reported values are the values measured during the second heating ramp. Reprocessing: the material was cut into 1 x 1 x 1 mm 3 pieces and then pressed in a PTFE mold for 1.5 h at 100 °C under a 6 tons load using a Carver 3960 manual heating press.

Figure 1 .

 1 Figure 1. Evolution of the epoxy FTIR band at 908 cm -1 of the TPE-TAF/BDGE binary mixture versus time at room temperature (ca. 20 °C)

1 a

 1 Scheme 2 a. Network synthesis from TPE-TAF and BDGE and b. Schematic representation of the network

1 , 1 . 42 ,

 1142 1 mol% Zn catalyst by Leibler et al. did not show vitrimer properties. 5 Normalized stress-relaxation experimental curves shown in Figure 3 reveal that, in the 170 °C -210 °C temperature range, the TPE-TAF/BDGE network relaxed the stress applied. This behavior is expected only if the network is able to reorganise. The fact that this material flowed on the rubbery plateau proves that fluorine atoms activated the transesterification. A contribution from minute amounts of unreacted carboxylic acid function might as well favor the transesterification. 60 Furthermore, the relaxation rate depended on the temperature, as expected for a vitrimer. The usual fitting model for such experiments is the exponential decay (Maxwell model). Nevertheless, this model did not fit well the experimental data obtained in the logarithmic time scale (R² between 0.934 and 0.989). A Kohlrausch-Williams-Watts "stretched exponential" equation was found to better fit (see Table S2 for the equation and the fitting parameters) the stress relaxation dataset 1 (R² between 0.995 and 0.99975, depending on the temperature). In KWW model, a stretch parameter β is added to the exponential decay. The closer to 1 is β, the closer to the Maxwell model are the data. From 170 °C to 210 °C, the value of β was about 0.56±0.02, indicating that the flow kinetics was associated with a distribution of relaxations rather than with a single relaxation time kinetics. 1 The relaxation time values obtained from the KWW fitting equations were plotted in an Arrhenius diagram (Figure 3 inset) to determine the value of the flow activation energy E a .This E a was determined to be 77 kJ mol in good agreement with the 29-163 kJ mol -1 range reported so far for transesterification vitrimers,47 especially for catalyst-free vitrimers activated by neighboring groups for which the activation energy values are ranging from 78 to 94 kJ mol -59-61,

Figure 3 .Figure 4 .

 34 Figure 3. Normalized stress-relaxation curves from 170 to 210 °C with 10 °C steps fitted with the Kohlrausch-Williams-Watts equation (KWW) and τKWW relaxation times reported in the Arrhenius diagram (inset, R² = 0.98985)

Figure 5 .

 5 Figure 5. TPE-TAF/BDGE aspect after several reprocessing cycles

Table

Table 2 .

 2 Evolution of DMA characteristic values before and after 10 reprocessing cycles

	Reprocessing Cycle E' Pristine T α (°C) 39	3.7	18.1
	10	49	3.9	18.7

G (GPa) E' R (MPa)

a value at Tα -50 °C b value at Tα + 50 °C

Table 3 .

 3 Evolution of Tα with the successive reprocessing cycles

	Reprocessing Cycle	Pristine	1	2	3-10
	T α (°C)	39	41	47	49

  Hz), 63.8 (O-CH 2 CH 3 ), 51.4 (Ar-C-CH 3 ), 30.9 (Ar-C-CH 3 ), 14.0 (O-CH 2 CH 3 ). Rf (petroleum ether : ethyl acetate 5:1) = 0.30.

			2 g, 61 %, white
	solid) and diester (F2: 1.8 g, 16 %, off-yellow viscous oil).
	Triester TPE-TE characterizations (NMR spectra in Figures S1 to
	S3):	1 H NMR 400MHz CDCl 3 : δ 7.15 -7.10 (m, 6H, aromatic
	protons, m-OCF 2 ), 7.07 -7.02 (m, 6H, aromatic protons, o-OCF 2 ), 4.39 (q, 3 J = 7.2 Hz, 6H, OCH 2 CH 3 ), 2.15 (s, 3H, Ar 3 C-CH 3 ), 1.36 (t, 3 J = 7.1 Hz, 9H, OCH 2 CH 3 ). 19 F NMR 377 MHz, CDCl 3 : δ -76.41. 13 C NMR 101 MHz, CDCl 3 : δ 159.9 (CO 2 Et, t, 2 J C-F = 41.5 Hz), 148.0 (ipso-Ar C-O, t, 3 Ar C-O, t, J C-F = 2.0 Hz), 146.5 (ortho-4 J C-F = 2.0 Hz), 129.9 (ipso-Ar C-C-CH 3 ), 121.2 (ortho-Ar C-C-CH 3 ), 114.1 (OCF 2 ,t, 1 J C-F = 272.5 Hz), 63.8 (O-CH 2 CH 3 ), 51.7 (Ar-C-
	CH 3 ), 30.9 (Ar-C-CH 3 ), 14.0 (O-CH 2 CH 3 ). Rf (petroleum ether : ethyl acetate 5:1) = 0.42. HRMS (ESI+) Calc. for [M+Na] +
	695.1686, found 695.1677.
	Disubstituted ester characterizations (NMR spectra Figures S4
	to S8):	1 H NMR 400MHz CDCl 3 : δ 7.14 -7.08 (m, 4H, aromatic
	protons, m-OCF 2 ), 7.07 -7.03 (m, 4H aromatic protons, o-
	OCF 2 ), 6.95 -6.89 (m, 2H aromatic protons, m-OH), 6.79 -6.72 (m, 2H, aromatic protons, o-OH), 5.06 (s, 1H, OH), 4.39 (q, 3 J =
	3 J = 7.2 Hz, 13 C NMR F NMR 377 MHz, CDCl 3 : δ -76.34. 7.1 Hz, 4H, OCH 2 CH 3 ), 2.12 (s, 3H, Ar 3 C-CH 3 ), 1.36 (t, 6H, OCH 2 CH 3 ). 19 101 MHz, CDCl 3 : δ 160.0 (CO 2 Et, t, 2 J C-F = 41.2 Hz), 154.1, 147.7 (ipso-Ar C-O, t, 3 J C-F = 2.0 Hz), 147.2, 140.6, 129.9, 129.9, 121.1
	(ortho-

Ar C-O t, 4 J C-F = 0.7 Hz), 115.0, 114.1 (OCF 2 , t, 1 J C-F = 272.4 "TPE-TAF" compound:

Acknowledgements

This work was funded by the Institut Carnot Chimie Balard CIRIMAT (16CARN000801) and the French National Research Agency ANR (AFCAN project -ANR-19-CE06-0014).

The authors would like to thank Lise Reymond for her contribution to TPE-TAF synthesis, as well as Dr. Marc Guerre and Dr. François Tournilhac for fruitful discussions.

Notes and references

Author Contributions

Conceptualization: FC, SC, ED, SL, EL, VL. Funding acquisition: SC, ED, EL, VL. Investigation: FC, SL. Methodology: all authors. Project administration: SC, ED, EL, VL. Supervision: FC, SC, ED, EL, VL. Validation: all authors. Visualization: FC. Writingoriginal draft: FC. Writing -review & editing: all authors.

Conflicts of interest

There are no conflicts to declare


Ali Kalush

Damien Texier
email: damien.texier@mines-albi.fr

Maxime Ecochard

Quentin Sirvin

Karim Choquet

Thomas Gheno

Nicolas Vanderesse

Walid Jomaa

Philippe Bocher

Size effects on high temperature oxidation of MCrAlY coatings processed via APS and HVOF depositions

Keywords: Oxidation MCrAlY coating, Ultrathin samples, Reservoir effect, Intrinsic chemical failure, Mechanically induced chemical failure 1

INTRODUCTION

Structural materials operating at high temperature, such as Ni-based superalloys, are generally coated with MCrAlY (where M = Ni, Co or Ni/Co) or aluminide coatings to improve their protective resistance in severe environmental conditions [START_REF] Mévrel | State of the art on high-temperature corrosion-resistant coatings[END_REF][START_REF] Rhys-Jones | Coatings for blade and vane applications in gas turbines[END_REF][START_REF] Bose | High temperature coatings[END_REF][START_REF] Nicholls | Designing oxidation-resistant coatings[END_REF]. MCrAlY coatings are widely applied to components exceeding 850 C, such as first and second stage turbine blades and nozzle guide vanes, compressors, turbine shrouds. They are richer in Al (and Cr) than superalloy substrates, which promotes the formation of a continuous, slow-growing, compact and adherent  -Al 2 O 3 scale at temperatures above 1000 C [START_REF] Evans | Oxidation of high-temperature coatings[END_REF][START_REF] Sloof | On the high-temperature oxidation of MCrAlY coatings[END_REF]. The consumption of Al due to both oxide growth and interdiffusion with the substrate [START_REF] Sloof | On the high-temperature oxidation of MCrAlY coatings[END_REF][START_REF] Vialas | Effect of Pt and Al content on the long-term, high temperature oxidation behavior and interdiffusion of a Pt-modified aluminide coating deposited on Ni-base superalloys[END_REF][START_REF] Raffaitin | Cyclic oxidation of coated and uncoated single-crystal nickel-based superalloy MC2 analyzed by continuous thermogravimetry analysis[END_REF][START_REF] Elsaß | Interdiffusion in MCrAlY coated nickel-base superalloys[END_REF] is central to the system lifetime, as it determines the integrity of the protective Al 2 O 3 scale, and thus the delay until faster-growing oxides develop. The -Al reservoir‖ available in the coating can become a critical factor in aeronautical applications, where coating thicknesses are limited due to weight concerns.

MCrAlY coatings can be manufactured using various techniques depending on the geometry of the substrate to be overlaid [START_REF] Tamarin | Protective coatings for turbine blades[END_REF]. On the one hand, electroplating and electro-codeposition routes have the advantage of producing dense and oxide-free coatings on complex shapes because of their non-line-of-sight deposition capability [START_REF] Raffaitin | Cyclic oxidation of coated and uncoated single-crystal nickel-based superalloy MC2 analyzed by continuous thermogravimetry analysis[END_REF][START_REF] Zhang | An Alternative Low-Cost Process for Deposition of MCrAlY Bond Coats for Advanced Syngas/Hydrogen Turbine Applications[END_REF][START_REF] Texier | Tensile properties of a non-line-of-sight processed  - - ' MCrAlY coating at high temperature[END_REF]. For instance, CrAlY-based powders can be codeposited into a (Ni, Co) matrix during the electroplating process.

Appropriate post-plating heat treatments are then applied to allow for element interdiffusion between CrAlY-based particles and the matrix. This operation confers mechanical properties to the particle/matrix structure and forms the final MCrAlY microstructure. However, this deposition quality comes at the expense of the process price and productivity [START_REF] Tamarin | Protective coatings for turbine blades[END_REF][START_REF] Tejero-Martin | Beyond traditional coatings: A review on thermal-sprayed functional and smart coatings[END_REF]. Line-of-sight deposition techniques, such as APS and HVOF deposition processes, are much more affordable and demonstrate a very interesting trade-off between a cost-efficient process and excellent protection for various industrial applications. However, these projected coatings are much more prone to internal defects such as unmelted particles, pores, internal oxides and/or nitrides, intersplats. Oxygen ingress along these open channels causes oxide intrusions within the coating [START_REF]Characterization of plasma sprayed and electron beam-physical vapor deposited thermal barrier coatings[END_REF][START_REF] Léger | Development of a process window for a NiCoCrAlY plasma-sprayed coating[END_REF][START_REF] Niranatlumpong | Failure of protective oxides on plasma-sprayed NiCrAlY overlay coatings[END_REF][START_REF] Belzunce | High temperature oxidation of HFPD thermal-sprayed MCrAlY coatings in simulated gas turbine[END_REF][START_REF] Evans | Diffusion cells and chemical failure of MCrAlY bond coats in thermal-barrier coating systems[END_REF][START_REF] Chen | Oxidation and crack nucleation/growth in an air-plasma-sprayed thermal barrier coating with NiCrAlY bond coat[END_REF][START_REF] Shibata | Comparison of microstructure and oxidation behavior of conicraly bond coatings prepared by different thermal spray processes[END_REF][START_REF] Patterson | Thermal cyclic lifetime and oxidation behavior of air plasma sprayed conicraly bond coats for thermal barrier coatings[END_REF][START_REF] Seo | Comparative study on oxidation behavior of selected MCrAlY coatings by elemental concentration profile analysis[END_REF][START_REF] Richer | Oxidation behaviour of CoNiCrAlY bond coats produced by plasma, HVOF and cold gas dynamic spraying[END_REF][START_REF] Ma | Isothermal oxidation behavior of cryomilled NiCrAlY bond coat : Homogeneity and growth rate of TGO[END_REF][START_REF] Wang | Influence of -island-like‖ oxides in the bond-coat on the stress and failure patterns of the thermal-barrier coatings fabricated by atmospheric plasma spraying during long-term high temperature oxidation[END_REF][START_REF] Jiang | Effect of internal oxidation on the interfacial morphology and residual stress in air plasma sprayed thermal barrier coatings[END_REF][START_REF] Meng | Highly oxidation resistant and cost effective MCrAlY bond coats prepared by controlled atmosphere heat treatment[END_REF][START_REF] Meng | Highly oxidation resistant MCrAlY bond coats prepared by heat treatment under low oxygen content[END_REF][START_REF] Texier | Tensile behavior of air plasma spray MCrAlY coatings: Role of high temperature agings and process defects[END_REF][START_REF] Texier | Screening for al2o3 failure in mcraly coatings using short-term oxidation at high temperature[END_REF][START_REF] Chen | Effects of internal oxide contents on the oxidation and  -phase depletion behaviour in HOVF CoNiCrAlY coatings[END_REF]. This oxide intrusion, concomitant to external oxide scale growth, participates in the Al consumption. Therefore, the capability of the coating to form a continuous Al 2 O 3 scale, which depends on the available Al reservoir, might be affected by the defect morphology (volume fraction, connectivity). Oxide intrusions were also found to participate in the coating swelling due to the Pilling-Bedworth ratio higher than 1 for Al 2 O 3 -forming MCrAlY alloys [START_REF] Jiang | Effect of internal oxidation on the interfacial morphology and residual stress in air plasma sprayed thermal barrier coatings[END_REF].

Oxidation-assisted swelling of the coating generates tensile stress in the external scale, which may lead to scale cracking and spalling [START_REF] Texier | Tensile behavior of air plasma spray MCrAlY coatings: Role of high temperature agings and process defects[END_REF].

Thermal and environmental conditions can vary depending on the application of the coated components. Maximal temperatures of 1100-1150 C are generally encountered for MCrAlY coatings, covered or not with a ceramic top coat for thermal barrier coatings (TBC) applications.

The formation of the Al 2 O 3 scale and intrusions, together with the interdiffusion with the substrate, causes aluminum depletion in the coating. In thin products, this may bring the Al concentration down to a critical level where Al 2 O 3 is no longer stable [START_REF] Texier | Screening for al2o3 failure in mcraly coatings using short-term oxidation at high temperature[END_REF][START_REF] Whittle | Spalling of protective oxide scales[END_REF][START_REF] Ishii | Effect of rare-earth elements on high-temperature oxidation resistance of Fe-20Cr-5Al alloy foils[END_REF][START_REF] Evans | Mechanisms of breakaway oxidation and application to a chromia-forming steel[END_REF][START_REF] Strehl | The effect of aluminium depletion on the oxidation behaviour of FeCrAl foils[END_REF][START_REF] Young | A simple expression for predicting the oxidation limited life of thin components manufactured from FCC high temperature alloys[END_REF][START_REF] Young | Slow transition from protective to breakaway oxidation of haynes 214 foil at high temperature[END_REF][START_REF] Gheno | Alumina failure and post-failure oxidation in the nicocraly alloy system at high temperature[END_REF][START_REF] Romain | Oxidation of thin nickel-based superalloy specimens: Kinetics study and mechanical integrity[END_REF]. Other oxides may then form at the oxide/gas and metal/oxide interfaces on large regions. This evolution of the protective Al 2 O 3 scale is called intrinsic chemical failure (InCF [START_REF] Evans | Mechanisms of breakaway oxidation and application to a chromia-forming steel[END_REF]) and is not necessarily associated with an increase of the mass gain [START_REF] Texier | Screening for al2o3 failure in mcraly coatings using short-term oxidation at high temperature[END_REF][START_REF] Ishii | Effect of rare-earth elements on high-temperature oxidation resistance of Fe-20Cr-5Al alloy foils[END_REF][START_REF] Strehl | The effect of aluminium depletion on the oxidation behaviour of FeCrAl foils[END_REF][START_REF] Gheno | Alumina failure and post-failure oxidation in the nicocraly alloy system at high temperature[END_REF]. Mechanical, thermal cyclic and thermal transient loading promotes mechano-chemical coupling and the onset of early strain-and oxidation-assisted damage and cracking/spallation of the external oxide scale [START_REF] Hindam | Microstructure, adhesion and growth kinetics of protective scales on metals and alloys[END_REF][START_REF] Evans | Stress effects in high temperature oxidation of metals[END_REF][START_REF] Pollock | Oxide-assisted degradation of Ni-base single crystals during cyclic loading: the role of coatings[END_REF]. Repeated through-thickness cracking and/or spallation of the protective Al 2 O 3 scale may bring the Al concentration in the vicinity of the oxide spalled and cracked region to a critical healing level. In comparison with InCF, this change in oxidation behavior is local: .. ie , where the damage of the oxide scale occurs. This phenomenon is called mechanically induced chemical failure (MICF [START_REF] Evans | Mechanisms of breakaway oxidation and application to a chromia-forming steel[END_REF])

and favors the rapid growth of oxides of the base elements (Ni, Co) [START_REF] Sloof | On the high-temperature oxidation of MCrAlY coatings[END_REF][START_REF] Giggins | Oxidation of Ni-Cr-Al Alloys Between 1000° and 1200[END_REF][START_REF] Wilber | Life-time extension of alumina forming FeCrAI-RE alloys: Influence of alloy thickness[END_REF]. Localized and generalized Al 2 O 3 failure are highly detrimental to the integrity of the coating itself, and even more damaging in TBC systems. Indeed, fast growing oxides at the bond-coat/top-coat interface may cause the top-coat to spall, thereby exposing the metal parts to excessive temperatures [START_REF] Evans | Diffusion cells and chemical failure of MCrAlY bond coats in thermal-barrier coating systems[END_REF][START_REF] Busso | Effects of breakaway oxidation on local stresses in thermal barrier coatings[END_REF]. The integrity of the coated system comes from its native protection against environmental stresses through the Al 2 O 3 formation and its long-term chemical stability. Therefore, investigating the Al 2 O 3 failure for coated components is central for high temperature applications.

In the present study, oxidation experiments were conducted at 1150 C on ultrathin freestanding NiCoCrAlY coating specimens extracted from Air Plasma Spray (APS) and High Velocity Oxy Fuel (HVOF)-processed plates. The ultrathin specimens were prepared with a dedicated precision jig [START_REF] Texier | Micromechanical testing of ultrathin layered material specimens at elevated temperature[END_REF] in order to activate InCF for short-term high temperature exposures due to through-thickness Al depletion. Interrupted oxidation tests were performed to document the onset and progression of the generalized breakaway as a function of the exposure time.

Thermogravimetric analyses were conducted to study the mass variation associated with the formation of the oxidation products. The sensitivity to InCF was evaluated for both types of coatings and the macroscopic oxidation behavior was related to the oxide intrusion. Xray micro computed tomography (micro-CT) and cross-sectional characterizations were conducted in order to assess the proportion of oxide intrusion in comparison with external scale. In order to document the intrusive oxidation within APS and HVOF materials, the oxidation behavior and lifetime of both the projected materials were directly compared to the ones of a -defect-free‖ cast material with a similar chemical composition solely developing and external Al 2 O 3 scale (see specimen UT2 in Ref. [START_REF] Gheno | Alumina failure and post-failure oxidation in the nicocraly alloy system at high temperature[END_REF] for additional information on the microstructure, oxidation behavior and lifetime).

EXPERIMENTAL PROCEDURES

Material

NiCoCrAlY coatings were deposited by HVOF and APS processes on 304 stainless steel (SS) plates with dimensions of 2 50250 mm 3 . Both coatings were deposited by PRAXAIR Surface Technologies for this specific investigation. The NiCoCrAlY powder used in the present study was provided by PRAXAIR  (NI-191-4 powder). The powder diameter was about 45  10 μm and its chemical composition is given in Table 1. Coatings with a thickness of 600 μm and 1 mm were deposited on the SS plates with the APS and HVOF processes, respectively. On the one hand, debonding occurred during the thermal spray deposition for the APS coating. Therefore, free-standing coatings samples (FSCS) were naturally obtained with this process. On the other hand, HVOF coatings were adherent even for coating thicknesses of 1 mm. Rectangular samples were then cut with a precision cutting machine and the SS substrate was mechanically ground to roughly obtain 800  50 μm -thick HVOF FSCS.

The FSCS were heat-treated at 1080 C for 8 hours and then at 870 C for 20 hours to reproduce the post-deposition heat treatments that are routinely applied to improve the coating adhesion to the substrate by interdiffusion and consequently to optimize the mechanical properties of Ni-based superalloy substrates [START_REF] Texier | Effect of interdiffusion on mechanical and thermal expansion properties at high temperature of a MCrAlY coated Ni-based superalloy[END_REF]. Neither was required here, but a standard heat-treatment was nonetheless conducted to generate coating microstructures representative of service conditions. The microstructure of the APS and HVOF NiCoCrAlY coatings is illustrated in Fig. 1.It contained  and  phases with surface fractions about 58 % and 42 % , respectively (surface fractions obtained from image analyses). The  and  phases appear in dark and light gray in BSE mode, respectively (see Fig. 1). The microstructure of the APS and HVOF coatings presents features typical of thermal sprayed materials: melted and resolidified powder particles, unmelted powder particle, pores, dispersed alumina oxides and resolidified splats. Both surface and cross-sectional observations evidenced higher occurrence of unmelted powder particles for the HVOF coating (21  2 % using image analysis) for these deposition conditions compared to APS coating (5  2 % ). The presence of unmelted powder particles is not related to the deposition process itself but to the chosen process parameters, .. eg the powder temperature, velocity, projection distance etc. [START_REF] Léger | Development of a process window for a NiCoCrAlY plasma-sprayed coating[END_REF]. While the APS coating depicts a typical microstructure of high temperature projected coatings, the HVOF coating exhibits a rather cold processed microstructure.

Thin sample preparation

APS and HVOF FSCS were thinned down with a home-built precision Jig paired with a BUEHLER  grinding/polishing machine to obtain thin samples with different thicknesses. A particular care was paid on the precision jig stability to achieve low thickness variations over centimeter-size samples. The thinning procedure is fully detailed in Refs [START_REF] Texier | Micromechanical testing of ultrathin layered material specimens at elevated temperature[END_REF][START_REF] Texier | Measurement and Evolution of the gradient of mechanical properties in MCrAlY coated MC2 nickel based superalloy system[END_REF]. Samples with thickness ranging between 400 and 15 μm with less than  3 μm in thickness variation were prepared. A surface finishing of P4000 SiC grit paper was applied on each sample face. In addition, both the projected materials were compared to a cast material with a similar chemical composition but exempt from process-defects (see specimen UT2 in Ref. [START_REF] Gheno | Alumina failure and post-failure oxidation in the nicocraly alloy system at high temperature[END_REF] for additionnal information).

High temperature oxidation

Thicknesses were measured using a micrometer with an uncertainty of  1 m  , samples were photographed using a calibrated binocular magnifying glass (Leica DMS 300) to document the sample surface and measure the actual surface S . After preparation, samples were cleaned with acetone then ethanol using ultrasonic bath. Sample masses were measured with a precision balance (Sartorius MC5), with a precision of  1.0  10 6

Material characterization

Surface and cross-sectional observations of the oxidized samples were conducted using optical microscopy (OM) and scanning electron microscopy (SEM) in BSE mode. SEM observations were conducted using a Nova NanoSEM 450 field-emission gun SEM (FEG-SEM)

from Thermo Fischer Scientific and a JSM 7800F Prime FEG-SEM from JEOL for cross-sectional observations.

Optical observations were performed to quantitatively document oxide intrusion in the different materials using laser scanning confocal microscopy (Olympus LEXT OLS5100).

Automatic stitching in laser intensity mode was conducted to cover large regions of interest with a resolution compatible with the oxide thickness (external oxide scale and oxide intrusion). The whole cross-section was imaged for each sample to analyze sample length compatible with the sparsely distributed intrusive oxide for the APS coating reported in a previous study [START_REF] Texier | Tensile behavior of air plasma spray MCrAlY coatings: Role of high temperature agings and process defects[END_REF]. Images were thus segmented using Fiji [START_REF] Schindelin | Fiji: an open-source platform for biological-image analysis[END_REF] in order to dissociate the oxide intrusion from the external oxide and the remaining metal.

CT-scans were performed on thick samples of APS and HVOF coatings oxidized for 500 h to quantitatively document oxide intrusion in three dimensions. Two slabs with dimensions 1 2 0.2  mm 3 and 1 2 0.4  mm 3 were cut from the APS and HVOF samples, respectively.

They were characterized with a v tome xs 240 scanner (GE) equipped with a DXR250 detector (2024 2024  pixels 2 with a dimension of 200 200  μm 2 ). The scanning parameters were as follows: numerical resolution: 1.75 μm /voxel, voltage: 100 kV, intensity: 180 μA , exposure time: 1 s, filter: 0.1 mm thick Cu plate, 3 frames averaging, 2550 projections for 360 rotation.

The projections were processed with the program Phoenix datos-x reconstruction 2. After reconstruction, the 3D images were analyzed with Dragonfly 2021. 3 [50]. The images were segmented using the Otsu algorithm, enabling a quantitative and localized analysis of the intruded oxides within the matrix.

RESULTS

Mass gain evolution

Mass gain evolution was investigated for both the APS and HVOF coatings up to 250 cumulative hours for all the coating samples, as depicted in Fig. 2. The oxidation behavior of a cast material with a similar chemical composition was added as a green dashed curve on the different graphs for direct comparison with a material without projection defects (results from Ref. [START_REF] Gheno | Alumina failure and post-failure oxidation in the nicocraly alloy system at high temperature[END_REF]).

On the one hand, the mass gain of the APS coating was found nearly insensitive to the sample thickness and matched well with the one of the cast material. For short oxidation duration, the mass gain is slightly higher for the APS material in comparison with the cast material. The oxidation rate follows a power law with an exponent close to 3 similarly to the cast materials, as observed on the log-log representation. More details on the oxidation of the APS coatings are reported in Ref. [START_REF] Texier | Screening for al2o3 failure in mcraly coatings using short-term oxidation at high temperature[END_REF]. On the other hand, the mass gain evolution of the HVOF coating was found particularly sensitive to the specimen thickness. Thick samples demonstrate a greater mass gain for all the oxidation duration. For instance, the mass gain of the 367 μm -thick HVOF specimen after 250 hours is nearly 9 times greater than the one of the cast material. Interestingly, the oxidation rate of the HVOF samples follows also a power law with an exponent close to 3 regardless of the specimen thickness (parallel linear trends in Fig. 2d). In addition, the three thinnest HVOF specimens and the thinnest APS sample have a decrease in oxidation rate after 100 hours. log-log representations. The oxidation behavior of the cast material with a similar chemical composition was added as a green dashed curve on the different graphs for direct comparison (results from Ref. [START_REF] Gheno | Alumina failure and post-failure oxidation in the nicocraly alloy system at high temperature[END_REF]).

The oxidation kinetics were analyzed with a power law:

= n n m kt S      (1)
where m  is the mass gain, S is the external surface area, n and n k and the oxidation exponent and constant, respectively, and t is the time. Values of n and n k fitted on the mass gain data are reported as a function of the specimen thickness in Fig. 3. The average value of the exponent was found to be 3.5  0.6 for APS sample and 3.4  0.4 for the HVOF ones. While n k is insensitive to the thickness for the APS coating with a value of 6.9 10 3   5.710 3  3.5 mg .cm 7  s 1  (with n =3.5), its value increases with the thickness for the HVOF coating from 1.010 3  mg 3.5 .cm 7  s 1  to 4.210 0 mg 3.5 .cm 7  s 1  (with n =3.5).

Evolution of the oxide scale

Macrographs were recorded for both the APS and HVOF materials at each interruption (Fig. 4). After 5 hour oxidation, all the specimens have a homogeneous and continuous light gray aspect, corresponding to the stable Al 2 O 3 external oxide scale. For the APS coating, darker regions appear after 20 and 100 hours on the 24 and 33 μm -thick specimens, respectively. These darker regions, highlighted with green arrows in Fig. 4, correspond to the onset of InCF, as reported in Refs. [START_REF] Texier | Screening for al2o3 failure in mcraly coatings using short-term oxidation at high temperature[END_REF][START_REF] Gheno | Alumina failure and post-failure oxidation in the nicocraly alloy system at high temperature[END_REF]. Specimens with a thickness of 61 μm or greater did not experience any InCF. For the HVOF coating, InCF was found to occur after 20 hours on the 15 μm -thick specimen. Interestingly, Al 2 O 3 failure was observed for all the specimens for oxidation duration lower than 250 hours regardless of the specimen thickness (and even 150 hours for specimens thinner than 200 μm ). Oxide spallation was found to occur for both materials (light gray regions in Figs. 4 and5).

It became significant after 50 hours, and was more important in thicker samples. Thick APS coatings developed an external oxide scale with a homogeneous gray level (except oxide spalled regions) up to 500 hours, corresponding to Al 2 O 3 formation. Interestingly, thick HVOF coatings developed an external oxide scale with blueish islets spreading onto the surface. Blue oxide was mainly observed in regions where Al 2 O 3 spalled at the previous cooling step (see .. eg supplementary materials Fig. S1). This local change in oxidation behavior preceded by Al 2 O 3 spallation is consistent with MICF. Due to the increasing sensitivity to Al 2 O 3 spallation with increasing specimen thickness, more than 10 % of the sample surface is covered by MICF at 250 hours and 150 hours for specimens thinner and thicker than 150 μm , respectively. The blueish oxide scale covers the majority of the specimen surface after 350 hours. While the APS specimens were only prone to InCF, HVOF specimens can be subject to InCF and MICF simultaneously. It is also important to note that specimens thinner than 30 μm did not experience any oxide spallation and were solely sensitive to InCF. In addition, the oxide scale was less prone to spall once InCF or MICF developed.

Figure 5: Surface macrographs of a 238 μm -thick APS and a 367 μm -thick HVOF coating samples, showing the evolution of oxide spallation and the development of the blue oxide.

Cross-sectional observations

Cross-sectional observations were conducted on the APS and HVOF coatings after 250 hour cumulative oxidation exposure for thick and thin samples, as illustrated in Figs. 6 and7 respectively. Oxide intrusions were found in both materials, but extended deeper and developed thicker patterns in the HVOF coating. These intrusive oxides formed in intersplat and unmelted particle regions. For the APS coating, thin oxide stringers are observed and thick oxide intrusions are sparsely distributed on the cross-section views. Both the thick oxide intrusions and thin oxide stringers were characterized as Al 2 O 3 using energy dispersive X-Ray (EDX) analyses. Some AlN precipitates were also observed within the bulk of APS coatings inherited from the deposition process (see Ref. (Ni,Co)(Cr,Al) 2 O 4 spinels within intrusive oxides developed at depths shorter than 60 μm , .. ie at a depth less than 1.5 CrAlYTa particle size.

Quantitative analyses of the intrusive oxidation

Large cross-sectional observations and CT-scans were investigated in order to dissociate the contribution of the external oxide scale and the oxide intrusion after 500 h at 1150 C in two and three dimensions, respectively. While solely the thickest samples were investigated via tomographic observations, different samples thicknesses ranging from 55 to 367 μm were analyzed using cross-sectional observations. Illustrations of the 3D macrographs are depicted in As aforementioned, intrusive oxides were segmented and compared to the external oxide scale, as shown by the int S / ext S ratio in Fig. 9(a). Both surface ratio (solid markers) and volume ratio (blank markers), respectively obtained from cross-sectional views and 3D CT-Scan results, were plotted for direct comparison. For the APS samples, the surface/volume occupied by intrusive oxides was found slightly lower than the one of the external oxide scale ( / int ext SS ratio less than 1). The dark gray horizontal dashed line illustrates the = int ext SS condition. For HVOF samples, the contribution of intrusive oxides is greater and the / int ext SS ratio increased with the sample thickness. Indeed, the surface occupied by intrusive oxides was more than ten times greater than the external oxide surface for the 367 μm -thick HVOF sample after 500 h at 1150 C. Therefore, the contribution of oxide intrusions to the mass gain of the HVOF samples is further demonstrated, and explains the observed effect of the sample thickness on the oxidation kinetics.

In addition, the fraction of intrusive oxide within the metallic material was considered as the cross-sectional surface (or volume from CT-scans) of intrusive oxide Oxide segmentation also aimed to evaluate the distribution of intrusive oxidation within the bulk of the metallic material. The fraction of intrusive oxide within the metallic material was calculated at different locations from the center of the sample, as illustrated in Fig. 10. The profile of the intrusive oxidation fraction across the sample thickness has a -U‖-shape regardless of the sample thickness and deposition process. The external oxide scale is not represented in this graph.

Intrusive oxidation was greater in a shallow region beneath the external oxide scale then the fraction decreased in the sample core. Interestingly, fractions of intrusive oxidation greater than 40 % were reported for both materials. This fraction decreased to a core value of nearly 4 3 %  and 24 3 %  for the APS and HVOF materials, respectively. Cross-sectional and CT-scan results were also consistent.

DISCUSSION

The comparison between the high temperature oxidation of APS and a HVOF coatings with similar nominal composition provides an opportunity to document how processed-inherited features affect the oxidation behavior, .. ie mass gain evolution and oxide formation, and ultimately the lifetime. Both projected materials developed an external oxide scale together with intrusive oxidation along intersplats and around unmelted particles, with a greater sensitivity for the HVOF coating. The greater sensitivity to intrusive oxidation is not attributed to the HVOF process but to the particular process parameters for the HVOF deposition used in the present paper.

The comparison between APS and HVOF materials with the cast material with a similar chemical composition gives the opportunity to document the oxidation behavior when solely external oxidation takes place [START_REF] Gheno | Alumina failure and post-failure oxidation in the nicocraly alloy system at high temperature[END_REF], and provide an opportunity to discuss on the contribution of oxide intrusions on InCF, MICF, and lifetime, .. ie time to Al 2 O 3 failure.

Origin of intrusive oxidation

As aforementioned, APS and HVOF materials developed intrusive oxidation along intersplats and around unmelted particles. Intrusive oxidation was found more pronounced on the HVOF material due to a high fraction of unmelted particles (see .. eg Fig. 1). Intrusive oxidation was reported for different MCrAlY projected coatings, with a sensitivity not clearly associated to a deposition process [START_REF]Characterization of plasma sprayed and electron beam-physical vapor deposited thermal barrier coatings[END_REF][START_REF] Léger | Development of a process window for a NiCoCrAlY plasma-sprayed coating[END_REF][START_REF] Niranatlumpong | Failure of protective oxides on plasma-sprayed NiCrAlY overlay coatings[END_REF][START_REF] Belzunce | High temperature oxidation of HFPD thermal-sprayed MCrAlY coatings in simulated gas turbine[END_REF][START_REF] Evans | Diffusion cells and chemical failure of MCrAlY bond coats in thermal-barrier coating systems[END_REF][START_REF] Chen | Oxidation and crack nucleation/growth in an air-plasma-sprayed thermal barrier coating with NiCrAlY bond coat[END_REF][START_REF] Shibata | Comparison of microstructure and oxidation behavior of conicraly bond coatings prepared by different thermal spray processes[END_REF][START_REF] Patterson | Thermal cyclic lifetime and oxidation behavior of air plasma sprayed conicraly bond coats for thermal barrier coatings[END_REF][START_REF] Seo | Comparative study on oxidation behavior of selected MCrAlY coatings by elemental concentration profile analysis[END_REF][START_REF] Richer | Oxidation behaviour of CoNiCrAlY bond coats produced by plasma, HVOF and cold gas dynamic spraying[END_REF][START_REF] Ma | Isothermal oxidation behavior of cryomilled NiCrAlY bond coat : Homogeneity and growth rate of TGO[END_REF][START_REF] Wang | Influence of -island-like‖ oxides in the bond-coat on the stress and failure patterns of the thermal-barrier coatings fabricated by atmospheric plasma spraying during long-term high temperature oxidation[END_REF][START_REF] Jiang | Effect of internal oxidation on the interfacial morphology and residual stress in air plasma sprayed thermal barrier coatings[END_REF][START_REF] Meng | Highly oxidation resistant and cost effective MCrAlY bond coats prepared by controlled atmosphere heat treatment[END_REF][START_REF] Meng | Highly oxidation resistant MCrAlY bond coats prepared by heat treatment under low oxygen content[END_REF][START_REF] Texier | Tensile behavior of air plasma spray MCrAlY coatings: Role of high temperature agings and process defects[END_REF][START_REF] Texier | Screening for al2o3 failure in mcraly coatings using short-term oxidation at high temperature[END_REF][START_REF] Chen | Effects of internal oxide contents on the oxidation and  -phase depletion behaviour in HOVF CoNiCrAlY coatings[END_REF]. For instance, several authors reported that HVOF MCrAlY coatings were less prone to intrusive oxidation compared to APS (or even low pressure plasma spraying (LPPS)) [START_REF] Shibata | Comparison of microstructure and oxidation behavior of conicraly bond coatings prepared by different thermal spray processes[END_REF][START_REF] Richer | Oxidation behaviour of CoNiCrAlY bond coats produced by plasma, HVOF and cold gas dynamic spraying[END_REF][START_REF] Wang | Influence of -island-like‖ oxides in the bond-coat on the stress and failure patterns of the thermal-barrier coatings fabricated by atmospheric plasma spraying during long-term high temperature oxidation[END_REF]. The sensitivity to intrusive oxidation was found to be related to the presence and interconnection of unmelted powder particles. Process parameters such as powder temperature and velocity play a significant role on the fraction of unmelted powder particles in projected materials [START_REF] Léger | Development of a process window for a NiCoCrAlY plasma-sprayed coating[END_REF]. Hotter process conditions for this specific HVOF material could improve its oxidation resistance and lifetime. Intersplat regions also participate in intrusive oxidation but in a less extensive manner due to morphological aspects. Intersplat regions for the APS material are more likely found orthogonal to the projection direction, .. ie orthogonal to the oxygen ingress direction / inward oxygen flux. The lack of adhesion of unmelted particles in the HVOF material generates continuous channels, whereas intersplat regions appear more discontinuous (see .. eg Fig. 1). Therefore, defect connectivity from particles to particles seems greater for materials having unmelted particles, especially in depth, as depicted in Figs. 1.(c) and(d). This particular defect morphology favors oxide intrusion in depth, thus affecting the core of the projected material due to deep open porosity. With these specific process conditions, the APS material is more likely to have closed porosity while the HVOF materials has a greater open porosity. Oxidation of intersplat regions and/or unmelted particles can be significantly limited using additional surface treatments to close pores and channels [START_REF] Wu | Oxidation behavior of laser remelted plasma sprayed NiCrAlY and NiCrAlY-Al 2 O 3 coatings[END_REF].

Effect of intrusive oxidation on the Al 2 O 3 failure

The motivation of the present paper is not to discriminate a deposition process but to investigate intrusive oxidation related to process-inherited defects and ultimately its lifetime based on the Al 2 O 3 failure. Both external oxidation and intrusive oxidation participate in the mass gain evolution and in the consumption of metallic elements participating in the oxide formation.

Different kinetics of oxidation and oxides in presence were found for the APS and HVOF materials due to the much more severe intrusive oxidation for the HVOF process. Two different oxidation mechanisms were reported to participate for the Al 2 O 3 failure depending on the specimen thickness: (i) the intrinsic chemical failure (InCF) due to aluminum consumption, and (ii) the mechanically induced chemical failure (MICF) due to oxide spallation. These oxidation mechanisms are dependant on the specimen thickness since thin specimens are sensitive to reservoir effects and thick specimens are more prone to oxide spallation. Therefore, the competition between both the oxidation mechanisms will be discussed below for the APS and HVOF materials and compared to the cast material.

Based on the Pilling-Bedworth ratio of the metal/oxide system, development of oxide intrusions leads to swelling of the material or at least put the metal in a complex stress state. Such volume expansion can contribute to spallation of the external oxide scale, as observed in Ref. [START_REF] Texier | Screening for al2o3 failure in mcraly coatings using short-term oxidation at high temperature[END_REF] and can participate in the development of the intrusive oxidation. However, the coatings developed oxide intrusions before scale spallation started. Therefore, scale spallation was not the primary cause for intrusive oxidation.

APS materials with a thickness greater than 60 μm developed Al 2 O 3 as external and intrusive oxides for samples up to 500 hours. Despite oxide spallation, APS materials did not experience MICF in the time/temperature window. Samples as thin as 60 μm were necessary to activate InCF for the APS coating at 1150 C up to 500 hours oxidation, as depicted in Fig. 4 (time to InCF = 20 and 100 hours for 24 and 33 μm -thick APS specimens). InCF was found associated with the formation of a Cr 2 O 3 layer and CrN precipitates at the Al 2 O 3 /metal interface (See .. eg Fig. 7 and Ref. [30] for more details). Similar results were reported for the cast material experiencing solely external oxidation, as reported in Ref. [START_REF] Gheno | Alumina failure and post-failure oxidation in the nicocraly alloy system at high temperature[END_REF]. of the Al 2 O 3 scale and within intrusive Al 2 O 3 for specimens thicker than 55 μm (Fig. 6). This second stage of oxidation is associated with the blueish coloration of HVOF samples.

(Ni,Co)(Cr,Al) 2 O 4 islets formed and spread from regions subjected to oxide spallation at previous interruptions (see .. eg supplementary materials, Fig. S1). The formation of (Ni,Co)(Cr,Al) 2 O 4 spinels on top of the Al 2 O 3 scale corresponds to MICF and was not associated with a particular increase in the oxidation rate due to the continued presence of the slow-diffusing oxide [START_REF] Ishii | Effect of rare-earth elements on high-temperature oxidation resistance of Fe-20Cr-5Al alloy foils[END_REF][START_REF] Strehl | The effect of aluminium depletion on the oxidation behaviour of FeCrAl foils[END_REF][START_REF] Young | A simple expression for predicting the oxidation limited life of thin components manufactured from FCC high temperature alloys[END_REF][START_REF] Young | Slow transition from protective to breakaway oxidation of haynes 214 foil at high temperature[END_REF][START_REF] Fedorova | Quantification of growth kinetics and adherence of oxide scales formed on Ni-based superalloys at high temperature[END_REF]. While both APS and HVOF materials were subject to [START_REF] Richer | Oxidation behaviour of CoNiCrAlY bond coats produced by plasma, HVOF and cold gas dynamic spraying[END_REF] Al O spallation, solely HVOF formed large regions of (Ni,Co)(Cr,Al) 2 O 4 spinels for specimens with thickness ranging from 55 to 367 μm . Meng et al. reported similar formation of (Ni,Co)(Cr,Al) 2 O 4 spinels on top of the external Al 2 O 3 scale after isothermal oxidation of an APS MCrAlY alloys developing severe oxide intrusion at intersplat regions [START_REF] Meng | Highly oxidation resistant and cost effective MCrAlY bond coats prepared by controlled atmosphere heat treatment[END_REF][START_REF] Meng | Highly oxidation resistant MCrAlY bond coats prepared by heat treatment under low oxygen content[END_REF]. In addition to MICF, InCF occurred for HVOF specimens regardless of the specimen thickness in this temperature-time window. Specimen as thick as 367 μm developed InCF since the fraction of intrusive oxides was found 10 times greater than the one of the external oxide scale (Figs. 2 and9). Both InCF and MICF were observed for HVOF specimens thicker than 55 μm while thinner specimens were only sensitive to InCF due to lack in oxide spallation. Therefore, both oxide spallation and oxide intrusion are central in the Al 2 O 3 failure and lifetime of thin MCrAlY products.

Influence of oxide intrusion on the lifetime of projected MCrAlY materials

In an attend to predict the lifetime of material investigated in the this work, the analytical model previously described in Refs. [START_REF] Texier | Screening for al2o3 failure in mcraly coatings using short-term oxidation at high temperature[END_REF][START_REF] Gheno | Alumina failure and post-failure oxidation in the nicocraly alloy system at high temperature[END_REF] • The alloy molar volume is considered constant;

• The specimen creep due to oxide growth is negligible;

• The kinetics of oxidation for the external and intrusive oxidation are comparable due to the open porosity associated to intersplat/unmelted particles regions.

The time to failure InCF t can thus be expressed as follow (see Refs. [START_REF] Texier | Screening for al2o3 failure in mcraly coatings using short-term oxidation at high temperature[END_REF][START_REF] Gheno | Alumina failure and post-failure oxidation in the nicocraly alloy system at high temperature[END_REF]:

O InCF MCrAlY MCrAlY Al Al 3 1 = 4 n n M tw kM          (2)

CONCLUSION

The oxidation behavior of two projected NiCoCrAlY coatings, .. ie APS and a HVOF coatings, was investigated using thin freestanding specimens exposed at 1150 C up to 500 hours.

A large amount of unmetled particules led to severe oxide intrusion for the HVOF coatings. The main differences in oxidation behavior for both the coatings consist in:

• A greater mass gain for the HVOF coating condition investigated here compared to the APS coating and the model material;

• A mass gain increasing with the coating thickness for the HVOF coating due to severe oxide intrusion via oxygen transport through intersplat channels. The oxide intrusion was found more severe in the extreme surface compared to the specimen bulk due to oxygen transportation and/or diffusion via the intersplat channels. The oxidation rate of the APS coating was found insensitive to the specimen thickness as long as no intrinsic chemical failure (InCF) occurred. For APS coatings, the oxide intrusion was less pronounced and mainly connected to the external oxide scale;

• A greater sensitivity to oxide spallation for the HVOF coating, especially for thicker samples;

• A greater sensitivity to intrinsic chemical failure (InCF) for the HVOF coating. The first stage of InCF, .. ie the Al 2 O 3 failure to form Cr 2 O 3 and/or CrN at the Al 2 O 3 /metal interface, occurred for thin APS specimens (samples thinner than 60 μm ) and for all the HVOF samples (samples as thick as 367 μm ) in the time/temperature window. A second stage of InCF resulting in the formation of (Ni,Co)(Cr,Al) 2 O 4 spinels as an external oxide scale and within oxide intrusion occurred for all the HVOF samples.

It is worth noting that defects inherited from the deposition processes severely impair the oxidation resistance and lifetime of MCrAlY coatings under service. In the present study, the HVOF coating experienced a poor oxidation resistance compared to the APS coating. However, other deposition parameters could inverse this trend. The prediction of InCF occurrence based on the deposition process parameters and/or non-destructive analyses of microstructure is needed to better estimated the lifetime of coating system due to oxidation and interdiffusion. Additional surface treatments preventing oxygen transportation within intersplat channels could improve the reactivity of such projected coatings.

AKNOWLEDGEMENTS

This work was supported by the Agence Nationale de la Recherche (ANR)

[ANR-18-CE08-0003; ANR-JCJC-COMPAACT project funded from the AAPG2018].

7.

Figure 1 :

 1 Figure 1: Microstructure obtained with a SEM in a BSE mode of the APS and HVOF coatings after standard heat treatment: (a) and (b) Top view observations of the samples (normal to the projection direction), (c) to (f) Cross-sectional (CS) observations of the samples showing melted and resolidified powder particles, unmelted powder particle, pores, dispersed alumina oxides and resolidified splats.

g.

  Samples were weighed three times before and after some oxidation period allowing the mass gain () mt  / S to be quantified. Oxidation experiments were conducted at 1150 C using interrupted furnace tests in a Nabertherm N 11/H furnace in laboratory air (no forced convection) for a cumulative time of up to 500 hours. One sample was used for each thickness condition. Interruptions were done to observe local changes in the oxidation behavior, .. eg Al 2 O 3 failure and oxide spallation. Samples were vertically hung on alumina supports using holders made of platinum wire during oxidation tests. Samples were weighed and observed using optical microscopy (OM) at different interruptions: after 5, 10, 15, 20, 50, 100, 150, 250, 350, and 500 cumulative hours. Optical observations aimed at capturing Al 2 O 3 failure, .. ie new oxidation products either on top of the external oxide scale or at the  -Al 2 O 3 /metal interface owing to the semi-transparency of the  -Al 2 O 3 scale in the visible spectrum.

Figure 2 :

 2 Figure 2: Mass gain evolution at 1150 C for both the APS and HVOF free-standing coating specimens with various thicknesses: (a) and (c)  m/S versus time representations, (b) and (d)

Figure 3 :

 3 Figure 3: Evolution of the (a) oxidation rate exponent n and (b) oxidation constant

Figure 4 :

 4 Figure 4: Macrographs showing surface evolution for both the APS and HVOF coatings at different oxidation cumulative time and specimen thicknesses. Green arrows highlight the onset of InCF. Please see the high-resolution version of the present figure for better details on InCF events. Partly reprinted from Corrosion Science, 184, Texier et al., Screening for Al 2 O 3 failure in

  [START_REF] Texier | Screening for al2o3 failure in mcraly coatings using short-term oxidation at high temperature[END_REF] for additional information). (Ni,Co)(Cr,Al) 2 O 4 spinels developed above the external Al 2 O 3 scale but also in the specimen bulk of some HVOF specimens in the core of sub-surface intrusive oxides (white arrows showing lighter gray regions in the core of Al 2 O 3 intrusions in Fig.6(d)). These spinels were mainly found in a shallow region beneath and connected to the external oxide scale. The spinel formation was not necessarily associated with an advanced consumption of aluminum within the metallic material. These (Ni,Co)(Cr,Al) 2 O 4 spinels were observed in regions of HVOF specimens with a  - microstructure in the specimen core. EDS analyses on the 200-and 367μm thick HVOF specimens revealed a large variability in Al content from regions to regions but Al content as high as 24 at.% was found far from regions developing (Ni,Co)(Cr,Al) 2 O 4 spinels.In addition, a slight undulation of the surface was observed for thick HVOF samples, especially in the vicinity of sub-surface oxide intrusions. This undulation might result from growth stresses associated with the formation of the intense intrusive oxides. The Al 2 O 3 would expand more than the metal it consumes to form the intrusive oxides (Pilling-Bedworth ratio of Al 2 O 3 in Ni-based alloys nearly 1.8[START_REF] Xu | Pilling-Bedworth ratio for oxidation of alloys[END_REF]), thus deforming the material.

Figure 6 :

 6 Figure 6: Cross-sectional observations of thick-samples after 250 hours at 1150 C showing difference in intrusive oxidation for a (a) 520 μm , (c)200 μm , (e)140 μm -thick APS and a (b)367 μm , (d) 367 μm , (f) 367 μm -thick HVOF coatings.

Figure 7 :

 7 Figure 7: Cross-sectional observations of thin-samples after 250 hours at 1150 C showing difference in intrinsic chemical failure (InCF) for 24 μm -thick APS and 25 μm -thick HVOF coatings.

Figure 8 :

 8 Figure 8: Cross-sectional observations of thick samples after 500 hours at 1150 C showing difference in external oxide scale composition between APS and HVOF coatings.

"

  supplementary materials" (Figs. S2 to S6). Intrusive oxides were thus segmented for both 2D and 3D characterizations to quantify their fraction and distribution within the sample and the fraction of intrusive oxidation compared to the external oxide scale. Segmented surfaces of the external oxide scale ( ext S ), of oxide intrusions ( int S ) and of the metallic region ( met S ) were illustrated in "supplementary materials" (Fig. S7).

intSV

  divided by the cross-sectional surface (volume) of the sample except the surface (volume) of the external surface ( met S ). Please refer to "supplementary materials" (Fig.S7) for illustration of are considered using 3D CT-Scan results. The fraction of intrusive oxide within the metallic material for the whole samples was depicted in Fig.9(b). This fraction was roughly constant, regardless of the sample thickness for both the APS and HVOF materials. After 500 h at 1150 C, the intrusive oxide occupied 8 and 35 % of the metallic sample for the APS and HVOF materials, respectively. 2D cross-sectional (CS) and 3D CT-scans (CT-S) results are consistent.

Figure 9 :

 9 Figure 9: Evolution of the surface fraction of intrusive oxidation as a function of the sample thickness for both the APS and HVOF coatings after 500 hours at 1150 C. (a) Comparison with the external oxide, (b) comparison with the metallic material.

Figure 10 :

 10 Figure 10: Distribution of the fraction of intrusive oxidation within the metallic material across the sample thickness for both the APS and HVOF after 500 hours at 1150 C. The external oxide scale is not represented in this graph.

  InCF resulted in the formation of a Cr 2 O 3 layer at the Al 2 O 3 /metal interface without precipitation of CrN. The cast materials experienced InCF after 20 and 200 hours for 11 and 26 μm -thick specimens, respectively. These longer time to Al 2 O 3 failure arise from the absence of intrusive oxidation, thus limiting the aluminum consumption compared to the APS material. For the HVOF material, both InCF and MICF were observed in competition depending on the specimen thickness. Severe intrusive oxidation and oxide spallation for the HVOF material impaired the stability of Al 2 O 3 . A continuous (Ni,Co)(Cr,Al) 2 O 4 spinel layer developed on top

  is used. The assumption of flat diffusion profiles and mass balance equilibrium allows the analytical estimation of the Al concentration in the metallic MCrAlY as a function of the mass gain and thus the time. Time to InCF can thus be considered as the time necessary to consume all the Al content from the MCrAlY alloy due to external and intrusive oxidation. The model is based on the hypotheses detailed below: • Composition profiles are nearly flat due to fast diffusion in both  and  phase, relative to the oxidation rate; • The quantity of non-oxidized elements, .. ie Co, Cr and Ni, remains constant in the metallic alloy during the Al 2 O 3 growth stage;

23 

 23 M the molar mass of O and Al, Al O the density of Al 2 O 3 , MCrAlY  the thickness of the coating sample, and Al w the weight fraction of Al in the MCrAlY alloy. Both model and experimental data were compared in Fig. 11. For the HVOF material, n k was found to increase with the specimen thickness and can be expressed as a power law function of the specimen thickness (see .. eg Fig. 3(b)). Dashed lines correspond to the time to InCF for all the materials and colored domains correspond to time-thickness domain for which InCF is supposed to occur. While the time to InCF rapidly increases for the APS and cast materials for specimens thicker than 30 m  , time to InCF for HVOF specimens was found to linearly and slowly increase with the specimen thickness. For the experimental results, specimens covered by at least 10 % by InCF or MICF were considered as affected by InCF or MICF, respectively. Both experiments and models were found to particularly match for the APS material. For the HVOF materials, the model tends to slightly overestimate the time to InCF for thick specimens. It is worth mentioning that oxide spallation at the different interruption was not considered both in the experimental mass gain measurement and in the model, leading to faster consumption of Al. The distribution of oxide intrusion might also play a role on the local consumption of Al within the metallic MCrAlY alloy due to presence of residual  precipitates when both InCF and MICF occurred for thick HVOF specimens.

Figure 11 :

 11 Figure 11: Lifetime prediction of the MCrAlY alloy as a function of the specimen thickness at 1150 C for the APS, the HVOF and the cast materials.

Figure 4

 4 Figure 4

Figure 5

 5 Figure 5

Figure 6

 6 Figure 6

Figure 7

 7 Figure 7

Figure 8

 8 Figure 8

Figure 9

 9 Figure 9

Figure 10

 10 Figure 10

Figure 11

 11 Figure 11

Table 1 :

 1 Nominal composition of the NI-191-4 powder used to deposit the NiCoCrAlY coating (in

	at. % and wt. %)
	Element	Ni	Co	Cr	Al	Y
	i x (at. %)	Balance	18.8	16.5	23.0	0.3
	w (wt. %)	Balance	22.0	17.0	12.3	0.5
	i

DATA AVAILABILITY

The raw/processed data required to reproduce these findings cannot be shared at this time due to technical or time limitations.