Merci À Martine

Isabelle Gaubert

Isabelle Hirsch

Roselyne Je

accepté de rapporter la thèse. Je leur souhaite une très

Introduction

The selection of structural materials for a given application is generally firstly based on their mechanical properties (mechanical engineering) and then on their environmental resistance (surface engineering).

By optimizing these two factors, it is therefore possible to increase the durability of the structural materials since their degradation is unavoidable with time.

Metallic materials are used in a wide field of structural applications including aero and land-based gas turbine engines. For such applications, metallic materials are carefully designed by adjusting their chemical composition and microstructure to meet the highly-demanding mechanical resistance. In the hottest sections of gas turbine engines, the high-temperature resistance of the metallic materials is of utmost importance for safety, economic and environmental considerations. For aeronautical applications, antagonist properties are generally required since the materials must exhibit superior mechanical properties, high surface stability (corrosion and oxidation resistance), lightweight and also be cost effective.

Gas turbine components (blades, vanes and disks) are submitted to highly corrosive environments (O2, SO2, H2O) with pressurized hot gases, combustion impurities and different types of erosive particles (sand, salts, volcanic ashes, etc.). To withstand these extreme conditions, nickel-based superalloys are employed because their improved microstructure based on coherent γ-γ' precipitation and optimized chemistry provide high creep and fatigue resistance at elevated temperatures (mechanical engineering).

In contrast, the surface stability of nickel-based superalloys is not sufficiently adequate over extended exposures at high temperatures under corrosive conditions. Therefore, the application of protective coatings on the surface allows to limit the corrosion and oxidation of the substrates (surface engineering). The application of protective coatings significantly increases the high-temperature surface stability of nickel-based superalloys and are essential for the integrity of gas turbine components in such aggressive environments. Depending on the temperature regime, different modes of degradation can be observed including hot corrosion, oxidation and CMAS (Calcium Magnesium Alumino-Silicates)induced degradation phenomena.

Based on the Brayton thermodynamic cycle, the efficiency of gas turbine engines is a direct consequence of the turbine inlet temperature (TIT). The higher this temperature, the greater the thermodynamic efficiency of the engine. Increasing the TIT by 50°C enhances the specific thrust by 10% and reduces the fuel consumption by 5%. Consequently, this temperature has been continuously risen since the beginning of modern aviation and the materials employed had to be more and more resistant. The increasing TIT also affected the usually non-attacked components (low stages of the turbine) to become sensitive to high-temperature corrosion. This will in future broaden the application of protective coatings to these stages.

For high-temperature oxidation resistance, industrial protective coatings are usually based on CVDrelated techniques to enrich the surface of nickel-based superalloys with Al (i.e. nickel aluminide coatings). The nickel aluminide coatings, based on the -NiAl phase, decrease the oxidation kinetics of the underlying substrate by forming a dense and continuous -Al2O3 scale on the surface. When spallation of the -Al2O3 scale occurs upon thermal cycling conditions (in particular because of thermal stresses induced by the fast cooling from elevated temperatures), the -NiAl phase promotes the outward diffusion of Al for its selective oxidation into -Al2O3. For the hottest sections (i.e. highpressure turbine or HPT), thermal barrier coatings (TBC) combined with internal cooling schemes are employed to decrease the surface temperature of the turbine components. The thermal barrier coatings, with low thermal conductivity, are usually based on YSZ (Yttria-Stabilized Zirconia). The combination of the metallic bond coat (usually a Pt-modified aluminide coating) and the ceramic top coat constitutes a so-called complete thermal barrier system. When hot corrosion resistance is required, Si and Cr enrichments offer the best protection forming respectively SiO2 and Cr2O3 at intermediate temperatures (between 600 and 1000°C). The corrosion resistance is usually conferred with thermal spray coatings (MCrAlY and MCoCrAlY coatings) or Si-modified aluminide coatings.

The increasing TIT implies that the turbine components of the low stages will need to be coated to fight against hot corrosion and that the thermal barrier systems will be extrapolated to the first stages of the low-pressure turbine (LPT). It is worth to notice that the conventional industrial processes handle hazardous chemicals (halide activators, chromic acid, etc.) and can produce lots of waste in the particular case of pack-cementation. Moreover, these coating processes are time-consuming as they involve many different steps. Therefore, to comply with environmental and safety regulations (e.g. REACh, Registration Evaluation Authorization and Restriction of Chemicals), new approaches have been developed in the last decades for the high-temperature protection of metallic materials.

-3-As a promising alternative approach, different studies have been carried out in the LaSIE laboratory under the European FP7 project "PARTICOAT". This approach was based on the formation of a complete thermal barrier system (metallic bond coat + oxide scale + ceramic barrier coating) from water-based Al microparticles-containing slurries. The complete thermal barrier system was formed in a single-step process with a nickel aluminide bond coat and a ceramic top coat composed of hollow Al2O3 spheres. The thermal insulation performances of these ceramic top coats appear greater than that of industrial YSZ thermal spray coatings, which lets expect their potential applications as thermal barrier systems. However, these thermal barrier systems present a low resistance to erosion and meet some limits because of the formation of inward grown coatings with high-activity microstructures.

Therefore, the main objectives of the present study were to control the activity of Al upon the slurry aluminizing process of nickel-based materials by introducing tailored amounts of Cr microparticles in the composition of the slurry or by changing the periodicity of the Cr and Al deposited layers. We focused our studies on the influence of the coating process parameters on the formation of the metallic bond coats and the complete thermal barrier systems. The present study focused on the addition of Cr microparticles in the Al slurry compositions as a potential candidate to overcome the abovementioned limitations. This manuscript has been structured in 6 main chapters:

Chapter I presents the state of the art and the industrial context for the high-temperature corrosion and protection of metallic materials in gas turbine engines. A general overview of the technological outcomes is given. The different modes of degradation encountered in gas turbine engines and their solution in the form of protective coatings are also introduced.

Chapter II describes the experimental methods of the present study including the materials of study, the elaboration of the slurry coatings and the environmental degradation tests performed. The characterization techniques carried out in this work are also presented.

In Chapter III, the influence of Cr additions on the exothermic reactions occurring between Al microparticles and nickel substrates is investigated. Different coating architectures were tested on pure nickel for a better comprehension of the mechanisms of formation before extrapolation to a Ni-Cr alloy and a nickel-based superalloy in Chapter IV.

The influence of the atmosphere (Ar, air) on the formation of complete thermal barrier systems is studied in Chapter V as well as the role of Cr on the involved aluminothermic reactions.

Finally, Chapter VI presents the environmental resistance of the slurry coatings designed in this work.

Hot corrosion and oxidation tests were performed to evaluate the protective behaviour of the slurry coatings.

-11-

I. Degradation and durability of metallic materials

Metallic materials are widely used for structural applications thanks to their unique properties compared to ceramics, polymers and composite materials. However, metals are generally not found in the nature in their native state (except gold, copper…) but under the form of more complex compounds (sulfides, sulfates, oxides, etc.). These compounds are formed over a long period of time due to the aggressive substances present in the atmosphere (oxygen, moisture, sulfur, etc.). Therefore, the production of metals generally requires a step of reduction of corresponding oxides and a step of purification. Two types of degradation can consequently be observed for metallic materials: Although some metals can be used in their pure state, metallic materials generally contain different other elements to increase the metal properties (constituting an alloy). Metallurgy is the science dealing with the elaboration of metallic materials by adjusting the chemical composition and carefully processing the material to get the desired microstructure. The main objective is to confer to the material the optimal properties for a given application (materials by design). In advanced applications, the durability of metallic materials may be shortened. Considering high-temperature structural applications, for energy and aeronautical industries for instance, the degradation of metallic materials usually results from several factors, often with a synergistic effect (Fig. I-2).

Fig. I-1. Schematic representation of (a) the electrochemical corrosion and (b) oxidation mechanisms.

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State of the art: context and motivations of the study -12-

For high-temperature structural applications, the durability of metallic materials is a matter of safety, economic and environmental considerations. In the particular case of gas turbine engines, the performance of metallic materials is of utmost importance to ensure the security of the whole aircraft.

The selection of materials is particularly difficult for these applications requiring antagonist properties: lightweight materials, superior mechanical properties, good resistance to corrosion and oxidation and cost effective.

Concerning the degradation of metallic materials, both thermodynamic and kinetic approaches must be considered. Thermodynamics can predict whether the constitution of a material is impaired or not for a given application whereas kinetics is useful to estimate how quickly the potential degradation of the material is detrimental. In high-temperature aeronautical applications, metallic materials are exposed to harsh environments, at an elevated range of temperature and pressure and must withstand incredibly high mechanical loads [1][2][3]. These factors lead to the unavoidable degradation of the parent material. These structural materials must then be carefully designed to resist at the very best to these extreme conditions (Mechanical engineering). For high-temperature protection of metallic materials, the key point is to limit ionic and electronic exchanges at the surface. Therefore, surface engineering is a critical issue to sustain the durability of metallic materials in aeronautical structural applications for elevated temperatures (Fig. I-2).

Fig. I-2. Schematic representation of the synergy between the environment and the structural conditions leading to the degradation of the metallic material in the high-temperature field.

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State of the art: context and motivations of the study -13-

II. Gas turbine engines: selection of the materials for high-temperature applications

Today, most civil and military aircrafts are provided with gas turbine engines (or turbofans) following the Brayton thermodynamic cycle [4,5]. The aim of a turbofan is to generate a propulsive thrust from the combustion of ingested air over the atmosphere. According to Newton's Third Law "Action and reaction forces pairs", the thrust leads to the propulsion of the aircraft [5,6]. Unlike military engines, only a portion of the air (about 20%) goes through a civil jet engine (primary flow). The remaining 80%

(secondary flow) is used to cool turbine components and to reduce noise pollution. These engines, called "high-bypass" [1,6], are particularly fuel efficient and provide the majority of civil aircrafts. Indeed, a large portion of the thrust is given by the secondary flow. For maximum propulsion (higher velocity), the whole ingested air is usually burned in military aircrafts.

Based on the Brayton thermodynamic cycle [4,5], turbofans are split in several units (Fig. I-3a). First, air is ingested by aspiration through a large fan. The air flow is progressively compressed in the lowpressure compressor then in the high-pressure compressor. As a result, a pressure greater than 30 bars [4] can be achieved increasing the temperature of the air flow (Fig. I-3b). Pressurized air is then injected in the combustion chamber and mixed with kerosene droplets. This mixture is ignited producing exothermic reactions which increases the temperature up to 1700°C [1,6]. Then, pressurized and hot gases are ejected from the combustion chamber (or combustor) to the next unit;

the turbine (highlighted by the red rectangle in Fig. I-3a). [7] and (b) evolution of the temperature and the pressure through the engine [8].

Fig. I-3. (a) Schematic view in cross section of a high-bypass turbofan

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State of the art: context and motivations of the study -14-

A turbine is divided into two sections, the high-pressure and low-pressure turbines (HPT and LPT respectively). The hot gases expand through the turbine and let the cylindrical shafts to rotate by friction on turbine blades (mobile stages). HPT and LPT rotate at two different speeds and they are in charge of driving the fan and the compressor (the high-pressure shaft spins at a higher velocity than the low-pressure one). The greater the velocity of the turbines the greater the air ingestion by the fan.

A turbofan operates therefore as an air flow cycle engine. For a better yield, turbofans are usually formed with alternatively fixed (stator) and mobile (rotor) stages [6]. Finally, the heated and pressurized gases are expelled through the exhaust nozzle and generate the thrust of the aircraft over the ambient air for propulsion (Fig. I-3a).

According to aviation requirements in terms of energy, time and cost savings, jet engines are expected to be more and more efficient and so require-mechanical, aerodynamic and material improvements.

The thermodynamic efficiency of a gas turbine engine is ensured by the Turbine Inlet Temperature (TIT in Fig. I-3b) referring to the Brayton cycle [4,5]. For instance, increasing the TIT by 50°C enhances the specific thrust by 10% and reduces the fuel consumption by 5%. Thus, this temperature has been continuously rising (more than 300°C in 30 years [1,6]) and the materials used to fabricate turbine components need to be carefully selected according to their metallurgy and mechanical properties (chemical composition, solidification process, creep and strain resistance). Obviously, in order to save energy and reduce environmental damages, the use of lightweight materials is a priority for aeronautical applications. However, the conditions are so aggressive in the turbine (high temperatures, strong mechanical stresses and corrosive atmosphere) that, only nickel and cobalt superalloys were found to withstand such extreme conditions [1]. These superalloys also present a relatively good oxidation resistance at intermediate temperatures (between 500°C and 800°C). Nickel-based superalloys are mainly used in the turbine where the temperature and the conditions are the most harmful. For instance, most LPT and HPT blades and vanes are made of nickel-based superalloys. Nickel-based superalloys are the preferred materials for the production of gas turbine components of military and civil aircrafts. A superalloy can be defined as a refractory metallic material designed to maintain superior mechanical properties even at temperatures close to its melting point [1,9]. Nickelbased superalloys rely on the austenitic γ-Ni matrix which is a face-centered cubic crystal structure (FCC) [9,10]. The stability of the γ-Ni phase from room temperature to its melting point allows the addition of other elements [2,11]. Therefore, nickel-based superalloys are complex materials containing usually five to ten alloying elements (Fig. I-4) [9][10][11]. Each alloying element has a potential effect on the mechanical properties, the corrosion/oxidation resistance and the density of the superalloy. Therefore, the alloying elements and their concentration must be carefully selected to produce a component with specific properties.

Solid solution elements like Co, Cr or Fe, with atomic radius close to that of nickel and refractory elements like Mo and W can reinforce and stabilise the FCC γ-Ni phase [1]. Hardening elements allow the formation of coherent precipitates in the matrix with a FCC γ'-Ni3X structure (where X refers to Al, Ti or Ta) [1,9,12]. Thus, nickel-based superalloys display a duplex γ-γ' microstructure (Fig. I-5) which can be cuboidal or spherical according to the superalloy and the post heat treatment [9][10][11][12][13].

With their typical γ-γ' microstructure, nickel-based superalloys present high creep, strain and fatigue resistance by impeding the motions of dislocations in the metallic structure [3].

Fig. I-4. Elements commonly selected for the constitution of nickel-based superalloys and their respective contribution to the microstructure [1].

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State of the art: context and motivations of the study -16-

The mechanical behaviour of the superalloy is a direct consequence of the proportion of the γ' phase.

Today, this phase represents about 70% by volume of a nickel-based superalloy [9,11]. By little addition of carbon in the chemical composition, the precipitation of carbides occurs during the solidification process [1,9,14]. These carbides mainly located at the grain boundaries increase the creep behaviour of nickel-based superalloys and reduce the impact of grain boundaries defects (formed during the solidification process) [2,14]. However, these carbides may serve as cracks initiation sites during fatigue for long term exposure [9,14,15].

Nickel-based superalloys display a good compromise between mechanical behaviour at high temperature and weight reduction in spite of a density between 7.7 to 9 g.cm -3 [2,9]. Even though new classes of lighter metallic materials have recently appeared, i.e. TiAl alloys [16,17], nickel-based superalloys are still the materials of choice for the production of gas turbine components of military and civil aircrafts.

III.2. Microstructural and compositional evolution

With the recent advances in aviation engine technology, the materials have to be more and more mechanically resistant with good strain and creep resistance because of high thermomechanical loads [10,18] and to better withstand the increase of the TIT over the last years. Nickel-based superalloys have been extensively developed through their chemical composition and microstructure to continuously improve their mechanical behaviour at elevated temperatures [1,9]. In the 1970's, turbine blades were made from wrought metallurgy (cast billets were deformed and reheated in a several steps process to give the component its final shape). Then, nickel superalloys containing more and more refractory elements (Mo and W mainly), the components had to be elaborated by powder metallurgy because of lack of ductility [2]. For these two technologies, the solidification process limited the exclusive production of equiaxed microstructures (Fig. I-6a). However, grain boundaries considerably decrease the high temperature creep resistance and thermo-mechanical fatigue (TMF) behaviour of the turbine components [2,19].

Since early 1980's, gas turbine components are mainly produced by investment casting [2,9,20]. With a better control of the solidification process, turbine components manufactured by investment casting present superior mechanical properties. In this process, the alloy is solidified in a preferential crystallographic orientation parallel to the applied-stress direction. This prevents the initiation of cracks through the grain boundaries in operation [9,21] -18-

With the investment casting process, the elaboration of cooling channels within the turbine components was made possible [1,2]. Cooling channels are useful to decrease the surface temperature of the alloy and to lighten the turbine components. Turbine blades are mainly produced by investment casting whereas vanes and disks can be produced by powder metallurgy. HPT blades are usually cast as a single crystal with cooling channels whereas LPT blades are usually produced with DS or polycrystalline superalloys (Fig. I-6). Indeed, single crystal superalloys are particularly expensive to produce and are located in the most critical region of the engine. The selection of the microstructure of the superalloy is a direct consequence of the temperature range (i.e. thermomechanical stresses level) [2].

In the second generation, the addition of Re, at the expense of Mo and W, considerably improved the creep resistance of nickel-based superalloys [12,14,22]. The Re concentration was then increased up to 6 wt.% in the third generation [22]. In the fourth generation, other elements such as Ru were added for the stabilization of the γ-γ' microstructure [23]. However, these strengthening elements increase both the price and the density of the alloys. Therefore, their concentration must be carefully selected to limit the weight of turbine components, which is of utmost importance for aeronautical applications.

IV. Degradation of metallic materials at high-temperature

Depending on the temperature on the alloy surface, three modes of degradation are usually considered (Fig. I-7). Unlike oxidation, hot corrosion phenomena are hardly predictable and lead to the fast degradation of the material at temperatures where the oxidation of the material is, usually, not detrimental.

IV.1. Oxidation

Oxidation is unavoidable in gas turbine engines due to the excess of oxygen in the atmosphere (the 2 is greater than the partial pressure required for the metal to be thermodynamically stable). The overall reaction of oxidation (Eq. I-1) can be written:

n + 2 → (Eq. I-1)

With:

M: metal

MnOy: corresponding oxide n: number of metallic atoms in the stoichiometry of the oxide With a negative standard Gibbs free energy change for the reaction (Eq. I-1) (ΔrG 0 < 0) [3], the oxide is more thermodynamically stable. However, oxidation is not necessarily detrimental for the metallic material and the kinetics of oxidation have to be considered to evaluate the progress of the degradation. Several steps are usually considered in the mechanisms of oxidation (Fig. I-1b). The composition and the morphology of the oxide formed is detrimental for the surface stability of the

Fig. I-7. Schematic illustration of the temperature dependence on the corrosion rate.

Chapter 1

State of the art: context and motivations of the study

-20-metallic material. Therefore, an oxide growing according to a parabolic law is usually preferred since this evolution is characteristic of a slow growing oxide. In the temperature range encountered in the turbines (Fig. I-3b), only a few oxides are able to bring a protection against oxidation to metallic surfaces: chromia (Cr2O3), silica (SiO2) and alumina (Al2O3) [2,3,24]. These oxides present a slow growth rate and a good adherence to the underlying material. On the contrary, fast growing oxides such as nickel oxide (NiO), spinels (NiAl2O4 or NiCr2O4) or cobalt oxide (CoO) are very porous and present a lot of defects. The porosity and the density of point defects is detrimental in oxide scales since they are paths for the diffusion of ionic and molecular species.

Another important parameter to evaluate the protective behaviour of an oxide scale is the Pilling and Bedworth ratio (PBR) [3]. The PBR is defined by the relation (Eq. I-2):

P = = * ⁄ ⁄
With:

M(oxide): molar mass of the oxide (g.mol -1 ) M(metal): molar mass of the metal (g.mol -1 )

n: number of metal atoms in the stoichiometry of the oxide ρ(oxide): density of the oxide (g.cm -3 ) ρ(metal): density of the metal (g.cm -3 )

(Eq. I-2)

If the PBR is less than 1, the oxide scale is exposed to tensile stresses whereas with a PBR superior to 1 the oxide scale undergoes compressive stresses. These stresses may induce cracks within the oxide scale and cause spallation. Therefore, an oxide scale with a PBR close to unity will be more protective in most cases. This parameter is of utmost importance under thermal cycling conditions since the CTE (Coefficient of Thermal Expansion) mismatch between the oxide and the underlying metal is large.

With a PBR of 1.28, alumina has the closest value to unity which explains why alumina forming alloys are particularly resistant to spallation in thermal cycling conditions encountered in gas turbine engines for aircraft power generation.

IV.2. Hot corrosion

Hot corrosion of high-temperature materials is an accelerated degradation induced by the condensation of a salt or a solid deposit on the surface of the metallic material. For aircraft gas turbines, the most common salt deposit is Na2SO4 (sodium sulfate) [25], resulting from the combustion of ingested NaCl with the gas flow comprising O2 and SO2 [26,27]. The suggested overall reaction of Na2SO4 deposition (Eq. I-3) is:

Chapter 1

State of the art: context and motivations of the study

-21-2 NaCl (g) + SO2 (g) + O2 (g) → Na2SO4 (s) + Cl2 (g) (Eq. I-3)

In the presence of water vapour, the overall reaction (Eq. I-4) of Na2SO4 deposition is:

4 NaCl (g) + 2 SO2 (g) + 2 H2O (g) + O2 (g) → 2 Na2SO4 (s) + 4 HCl (g) (Eq. I-4)

Two types of hot corrosion are usually considered depending on the range of temperature and of the salt composition (Fig. I-7). Assuming pure Na2SO4, Type-I hot corrosion occurs within the temperature range 850-950°C [18,[28][29][30] (HTHC: High Temperature Hot Corrosion) and Type-II within the range 650-800°C [18,29,31] (LTHC: Low Temperature Hot Corrosion). With a melting point of 884°C [30,31], Na2SO4 is then solid in Type-II conditions and liquid in Type-I conditions, which differentiates the two corrosion processes. However, the presence of other contaminants such as V2O5 (vanadium pentoxide) can lead to the formation of eutectic compounds with Na2SO4 (eutectic point lower than 600°C [26,30]), hence accelerating the corrosion rate [32]. When the temperature is higher than the salt dew point, dry corrosion referred to pure oxidation occurs (Fig. I-7). The different aspects of hot corrosion have been extensively studied for more than 50 years on metallic alloys and coatings for marine and aeronautical applications [33]. It is now well accepted than the mechanisms of hot corrosion consist in an incubation stage and a propagation stage [25]. During the incubation stage, an oxide scale is formed on the substrate and withstands the presence of the deposit at first. In the propagation stage, breakdown of the oxide scale occurs and internal oxidation occurs together with sulfidation, depending on the metallic coating and the temperature range.

 Type-I: High Temperature Hot Corrosion

For Type-I hot corrosion, the salt fluxing model is the most widely accepted one. In liquid phase, Na2SO4 decomposes on the metallic surface through the reaction (Eq. I-5):

Na2SO4 (l) → SO3 (g) + Na2O (l) (Eq. I-5)

Oxides can be dissolved in fused Na2SO4 as anionic species (basic fluxing) or cationic species (acidic fluxing) depending on the activities of SO3 (acidic) and Na2O (basic) [33]. As a consequence, a porous oxide scale is formed with an irregular metal-scale interface and an internal attack occurs within the material through the formation of sulfides. The solubility of an oxide can be estimated as a function of Na2O activity and SO3 partial pressure and is commonly represented in a solubility diagram [33] (Fig.

I-8a

). Rapp and Goto proposed that a "negative solubility" gradient in the fused salt film must exist to sustain the hot corrosion attack (Fig. I-8b). Chromia, alumina and silica are the most promising oxides for hot corrosion resistance. Silica is virtually insoluble in fused Na2SO4 but is usually less protective than chromia for Type-I hot corrosion [34].

 Type-II: Low Temperature Hot Corrosion

Low temperature hot corrosion was first observed in marine environment [34]. Since Na2SO4 is a solid in Type-II conditions, a sufficient amount of SO3 is necessary to observe hot corrosion [25,31]. The presence of SO3 converts NiO or CoO into sulfates hence inducing the formation of eutectic phases with the Na2SO4 deposit (Te (Na2SO4-NiSO4) = 671°C and Te (Na2SO4-CoSO4) = 565°C [36]). Type-II hot corrosion is usually characterized by a non-uniform attack of the metallic material in the form of pitting with minimal or no sulfidation.

From an engineering point of view, the key point to get protection against hot corrosion is to extend the incubation stage duration since no degradation of the material is observed at this stage. Alumina and chromia are expected to be stable in the normal conditions (considering the 3 ). However, the presence of Ni or Co within the oxide scale may change the mechanisms by forming eutectic phases with Na2SO4. Therefore, the purer the alumina or chromia scale the greater the hot corrosion resistance of the underlying material.

IV.3. Need for the high-temperature protection of nickel-based superalloys

Gas turbine blades and vanes located just after the combustion chamber are particularly exposed to hot corrosive gases and impurities coming from the kerosene, leading to the chemical degradation of the material. Because of the high concentration of oxygen in the combustion gases, surfaces oxidize in operation and the composition/morphology of the developed oxide scales are critical for the surface stability and the durability of the whole component [3]. Designed for superior mechanical resistance, nickel-based superalloys are vulnerable to aggressive environments forming non-protective oxide scales on the surface [18]. The concentration of beneficial elements (i.e. Al and Cr) is indeed proven to be not high enough to confer them good corrosion and oxidation resistance in the long term. Fig.

I-9 represents the evolution of the Cr and strengthening elements (Mo, Re, Ru) contents along the different generations of nickel-based superalloys. From mechanical considerations, Cr has been progressively replaced by these strengthening elements. Consequently, the new generations of nickelbased superalloys became particularly sensitive to hot corrosion since Cr is known as a beneficial element for hot corrosion resistance [33,37]. However, at temperatures above 900°C, chromia may form volatile species (e.g. CrO3) [18,24] and alumina is consequently the preferred oxide for higher temperatures.

Therefore, in the hottest regions of the engine, nickel-based superalloys need to be coated to confer the required corrosion and oxidation resistance. The high-temperature protection increases the durability of the gas turbine components by slowing down the degradation rate on their surface (surface engineering).

Fig. I-9. Evolution of the Cr and strengthening elements content over the different generations of monocrystalline nickel-based superalloys, from Wu et al. [38].

V.1. High-temperature protection: an overview

Since the oxidation phenomenon is thermodynamically unavoidable, the key point is to reduce at its lowest extent the kinetics of degradation. The driving force is to slow down ionic and electronic exchanges through the oxides on the metallic surface. Surface engineering has become an important research area for aeronautical applications. For the oxidation and corrosion resistance, metallic coatings are generally employed to promote the formation of a thin, continuous and dense oxide scale on the surface [3,18]. When the temperature range is higher, e.g. in the high-pressure turbine, thermal barrier coatings (TBCs) must be added on the underlying metallic coatings for thermal insulation [3].

Due to their low thermal conductivity, thermal barrier coatings can decrease the temperature on the surface of the alloy by up to 170°C [39][40][41].

Metallic coatings must present a mechanical and a chemical compatibility with the underlying substrate [3]. For aeronautical applications, metallic coatings are generally divided in two categories: overlay and diffusion coatings [6,12,18,39]. Overlay coatings, usually referred as MCrAlY, are particularly resistant to both Type-I and Type-II hot corrosion depending on the concentration of Cr [12,18]. These coatings are produced from metallic powders with a selected chemical composition and deposited on the surfaces of the alloy by plasma spray or by high velocity oxy-fuel (HVOF) [3]. Due to the elaboration process, the chemistry and the microstructure of the alloy do not play an important role on the final properties of an overlay coating.

Diffusion coatings are formed by the surface enrichment of the alloy whether with Al, Cr or Si and involve interdiffusion between the donor and the substrate [6,12,18]. Diffusion coatings are generally produced by chemical vapour deposition (CVD) [18], like pack cementation [28,[42][43][44] and different type of gas phase processes [28,[44][45][46]. Slurry coating processes are also classified as diffusion coatings [42,47]. Slurry processes present a high flexibility and they are easily applicable on metallic surfaces.

Historically, these coatings were mainly used for repair purposes [47]. Recently, the European project "PARTICOAT" demonstrated the possibility to elaborate a complete thermal barrier system (aluminide coating + ceramic thermal barrier coating) from environmentally-friendly slurries [48].

The enrichment of the substrate surface with a protective element (i.e. Al, Cr or Si) allows the formation of a thin and adherent oxide scale, sharply decreasing the kinetics of oxidation. The formation of alumina is particularly interesting since this oxide maintains a good stability at elevated temperatures and there is barely any point defects in the scale [3]. Therefore, diffusion coatings developed on nickel-based superalloys for oxidation resistance are generally based on Al (aluminizing).

Depending on the service conditions, these coatings can be modified with one or several elements to confer them specific properties and increase their performance. interdiffusion coefficient between Al and Ni in the phases of the Al-Ni system [56], which is dependent on the temperature and the Al donor source.

 HALT: High-Activity Low-Temperature

When the donor is rich in Al (i.e. with a high thermodynamic activity of Al), the inward diffusion of Al into the substrate occurs at relatively low temperature because the mobility of Al is higher than that of Ni at such temperatures [56]. The fast diffusion of Al drives the formation of Al-rich NixAly phases (i.e. NiAl3 and Ni2Al3) in the substrate. In these phases, the diffusion of Ni is not significant at low temperatures [56] and Al continues to diffuse inward [42]. Consequently, the slowly diffusing refractory elements (e.g. W, Mo, Ta) precipitate in the outer region of the coating. The inert substrate carbides are also entrapped in the inwardly growing layer (Fig. I-11). The initial surface of the substrate therefore remains at the outer surface of the coating upon aluminizing [57]. After the aluminizing step, the diffusion layer is based on Ni2Al3 [42]. Since Ni2Al3 is brittle and not stable at high temperature, a further annealing is required to obtain the desired β-NiAl phase [42,57]. The annealing is usually carried out between 1000°C and 1100°C [18]. Because of the nickel withdrawal from the substrate during the annealing treatment, an interdiffusion zone rich in precipitates formed above the unmodified substrate.

This microstructure is commonly observed for high-activity pack cementation process when Al powder is directly poured in the pack mixture [42,55,58,59] and simple aluminide slurry coatings [42,60].

Fig. I-11. Comparison between high-activity low-temperature (HALT) and low-activity hightemperature (LAHT) nickel aluminide coatings [42].

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State of the art: context and motivations of the study

-27-

 LAHT: Low-Activity High-Temperature

The microstructure of low-activity diffusion coatings is different from high-activity one with an interdiffusion zone and an additive layer (pure β-NiAl layer) with approximately the same thickness (Fig. I-11). By alloying Al with Ni [42] or Cr [17,45] (i.e. masteralloys), the thermodynamic activity of Al is decreased. Thus, Al is not able to diffuse inward at low temperature. At the aluminizing temperature (between 1000°C and 1100°C generally), Ni is able to diffuse outwardly and directly forms the Ni-rich β-NiAl phase by interdiffusion with the Al supply at the surface [42,61]. Since the diffusion coefficient of Ni is higher than that of Al in the stoichiometric and Ni-rich β-NiAl phase [56,62], the coating mainly grows by outward diffusion of Ni. As a consequence, the β-NiAl additive layer is free of precipitates from the slowly diffusing elements, present in the substrate, whose remain entrapped in the interdiffusion zone with the carbides [42]. Therefore, low-activity diffusion coatings are preferred for high-temperature protection as the external β-NiAl layer is cleaner (Fig. I-11) and

show better oxidation resistance than high-activity coatings [63]. This microstructure is commonly observed with out-of-pack processes [45,55,59,64] and pack cementation using masteralloys as Al donors [42]. Recently, Montero et al. achieved low-activity slurry coatings on nickel-based superalloys using pre-alloyed Al and Cr microparticles [61,65].

b) Simple and modified aluminide coatings

Nickel aluminide coatings are particularly resistant to oxidation at temperatures up to 1100°C [28].

When additional resistance is required to ensure the durability of the coating, nickel aluminide coatings are modified with beneficial elements. By decreasing the critical amount of Al to form an exclusive alumina scale [66,67], the addition of Cr is interesting for high-temperature protection. Thus, with the addition of 5 at.% Cr in the Al-Ni system, the critical level is reduced from 35 at.% Al to about 12 at.% Al [57]. This is commonly referred as the "third element effect" and was first proposed by Wagner [68]. The presence of Cr is also beneficial to accelerate the transformation of θ-Al2O3 to the stable α-Al2O3 [50]. For the hot corrosion protection, nickel aluminide coatings are generally modified with Cr and Si. Numerous studies investigated the beneficial effect of Cr and Si on the Type-I and Type-II hot corrosion behaviour of aluminide coatings [37,47,59,[69][70][71].

Reactive elements (RE) like Ce, Hf or Y were found to promote the adherence of alumina, especially in cyclic oxidation conditions [72,73]. Their addition was found to decrease the growth rate of alumina preventing the formation of voids underneath the oxide scale by forming mechanical bonding through oxide pegs at the surface of the diffusion coating [74]. RE were also assumed to limit the detrimental influence of S segregation at the coating-oxide interface [64].

Chapter 1

State of the art: context and motivations of the study -28-

In the hottest region of the engine, nickel aluminide coatings are modified with Pt to enhance the oxidation and hot corrosion resistance and can be used as environmental coatings or as bond coats for thermal barrier coatings [45,46,51]. The role of Pt on the oxidation resistance has been intensively studied over the last years [51,58,72,75]. However, there is no widely accepted effect on the beneficial role of Pt and different mechanisms were proposed. One beneficial effect commonly reported is the selective oxidation of Al with the addition of Pt, preventing the transient formation of Ni-rich oxides within the oxide scales [51,72]. Pt also improves the adherence of the alumina scale and prevents spallation in cyclic conditions [58,76].

For Pt-modified aluminide coatings, Pt is first electroplated (about 5-10µm thick) and is subsequently annealed to promote interdiffusion between Pt and the substrate. After the aluminizing process, the coating usually exhibits a single β-(Ni,Pt)Al phase [12,45,46,51]. However, a dual PtAl2 + β-(Ni,Pt)Al layer can be formed in the outer part of the coating depending on the initial Pt concentration and the Al activity during aluminizing [12,45,46,51,[77][78][79]. When Pt-modified aluminide coatings are used as bond coats for TBCs, the developed alumina scale acts as a TGO (Thermally Grown Oxide) ensuring a chemical and mechanical bonding between the metallic bond coat and the ceramic top coat in operation [51,[80][81][82]. The stability of the TGO is detrimental for the adherence of the ceramic top coat and is usually the critical life-limiting factor for TBC systems [80,[83][84][85].

V.3. Thermal Barrier Coatings (TBCs)

Thermal barrier coatings have been applied on gas turbine components for thermal insulation due to the continuously increasing TIT. These advanced coatings enhance the temperature capability of the underlying material by decreasing the temperature at the surface. In the HPT section of current gas turbine engines, these coatings are applied together with the elaboration of cooling channels within the material [3].

The conventional TBC systems are applied on gas turbine components located in the hottest region of the turbine (i.e. for HPT blades and vanes). They are composed of two distinct layers (Fig.

I-12):

─ A metallic bond coat generally developing an alumina scale at elevated temperatures acting as a TGO (Thermally Grown Oxide) to ensure a chemical and mechanical bond with the ceramic top coat. ─ A ceramic top coat with a low thermal conductivity to decrease the temperature at the surface of the metallic material.

Chapter 1

State of the art: context and motivations of the study -29-

In current gas turbine engines, the top coat is usually deposited by air plasma spray (APS) or electronbeam physical vapour deposition (EB-PVD) [28]. TBCs are usually made of yttria-stabilized zirconia (YSZ). With a low thermal conductivity, the ceramic top coat can decrease the temperature at the surface of the metallic material by up to 170°C [39][40][41]. These coatings present a lot of porosity and pathways allowing the diffusion of O2 toward the bond coat to grow and maintain the TGO at the interface bond coat-top coat. TBCs are also strain compliant and are able to accommodate thermal stresses in cyclic conditions due to porosity. The bond coats are generally made of Pt-modified aluminide coatings with the β-(Ni,Pt)Al phase (low-activity coatings) or MCrAlY coatings by plasma spray. The aim of these bond coats is to promote the formation of a pure α-Al2O3 scale (TGO) at the interface with the ceramic top coat. Since there is no chemical and mechanical continuity between the bond coat and the top coat, the TGO is a critical point for the durability of the whole TBC system [41,86].

V.4. Degradation of protective coatings in operation

In addition to the general mechanisms of degradation occurring in high-temperature structural applications (cf. IV.), more peculiar degradation can be observed on protective coatings in operation.

These modes of degradation are a direct consequence of the microstructure and the morphology of the coatings (intrinsic) or due to the environment (extrinsic).

Fig. I-12. SEM cross section image of a conventional TBC system (Pt-modified aluminide bond coat + YSZ EB-PVD top coat) showing the thermal insulation effect for a HPT blade [41].

Chapter 1

State of the art: context and motivations of the study

-30-

 Al depletion of nickel aluminide coatings:

Diffusion mechanisms are thermally activated and are very important at the elevated temperatures encountered in gas turbine engines. Under thermal cyclic conditions, the protective oxide scale developed on aluminide coatings is subjected to spallation. As a result, Al diffuses outwardly to maintain a protective alumina scale and Al is thus progressively consumed from the coating. In addition, Al diffuses into the substrate in operation due to the positive Al gradient between the coating and the substrate. Consequently, the progressive Al depletion of the coating causes the β-NiAl to γ'-Ni3Al transformation, which is usually observed in the grain boundaries of the coating at first [19,75,87]. As long as the activity of Al is sufficient in the coating, a protective alumina scale will be maintained on the surface. However, when the Al content reaches a critical value, non-protective oxides start to form and increase the oxidation kinetics. According to Giggins and Pettit, the critical amount of Al to maintain an exclusive alumina scale can be lowered with the addition of Cr [66]. The inward diffusion of Al into the substrate in operation is also responsible for the formation of a secondary reaction zone (SRZ) below the coating [88]. The progressive enrichment in Al of the substrate underneath the coating decreases the solubility of solid solution strengthening elements like W, Re or Mo [88]. Thus, the formation of needle-like brittle phases or topologically-closed packed phases (TCPs) is observed in the SRZ hence decreasing the mechanical properties of the alloy.

 TBCs failure:

TBCs failure is critical in operation since TBCs do not provide self-rejuvenating protection. In other words, when the TBCs spall off, the thermal insulation is lost [41]. The failure of TBC systems can result from different intrinsic or extrinsic features:

─ Rumpling of the metallic bond coat and the TGO due to extensive thermal cycling (formation of periodic undulations at the interface with the ceramic top coat) [51] ─ Formation of non-protective oxides rather than pure α-Al2O3 in the TGO due to the progressive Al depletion of the bond coat ─ Delamination across the interface bond coat-TGO or in the ceramic top coat due to erosion or FOD (Foreign Object Damage) ─ Environmental attack of the TBCs through the deposition and penetration of CMAS (Calcium-Magnesium-Alumino-Silicates) Therefore, TBCs are exposed to a combination of chemical and mechanical degradations in operation usually with a synergistic effect. The particular case of CMAS attack has been extensively studied over the last decades as it appeared as a critical failure mode for TBC systems [89][90][91]. CMAS results from the ingestion of siliceous mineral debris in the engine. When the temperature is high enough (usually

Chapter 1

State of the art: context and motivations of the study -31-over 1200°C), the CMAS melt and infiltrate the TBCs down to the TGO. As a result, the strain compliance of the ceramic top coat decreases and becomes brittle upon cooling [89,90]. The sintering of the ceramic top coat by the CMAS induces extensive delamination in cyclic conditions. Erosion of TBCs is also an important mode of degradation especially when large solid particles hit the surface of the component (FOD) [92].

The degradation of protective coatings in operation is therefore a combination of several features with, most of time, a synergistic effect. In order to propose the adequate or the most appropriate protection, the comprehension of the degradation mechanisms is of utmost importance. At laboratory scale, surface engineering tends to identify the influence of each mode of degradation on the durability of materials. Surface engineering also proposes to investigate the potential improvements (microstructural, compositional or architectural) to enhance the performance of metallic materials.

VI. Motivations of the present study: from "PARTICOAT" to "FONBAT"

With the overall increase of the temperature in gas turbine engines, for performance and socioeconomic considerations, the use of thermal barrier coatings need to be extended from the HPT to the LPT section (Fig. I-3). From an economic point of view, the extrapolation of current thermal barrier systems (cf. paragraph V.3) is not possible at an industrial scale. CVD-related techniques are particularly expensive, involve hazardous chemicals (halides) and produce lots of wastes in the case of pack cementation processes. EB-PVD and plasma spray coatings are also very expensive to produce and require high electrical supply. Moreover, some turbine components which used to be unaffected by high-temperature corrosion are now susceptible to low temperature hot corrosion (Type-II) in the low-pressure turbine. Therefore, many studies and projects came up recently to propose viable and cost effective alternative solutions. Among them, the slurry coatings appeared as a convincing choice to answer these energy issues, especially for the aircraft industry. In this way, the University of La Rochelle was involved in the European FP7 project "PARTICOAT" (2008-2012) [48].

VI.1. Presentation of the "PARTICOAT" concept

The main objective of the European project "PARTICOAT" was to propose a complete thermal barrier system (diffusion coating + ceramic top coat) in a single step process from Al microparticles in waterbased slurries (Fig. I-13). The mechanisms of formation of the coating were studied in argon on pure Ni [93], in air on nickel-based superalloys [94,95] and iron-based alloys [96].

-32-

The formation of the coatings on different nickel-based substrates was found to involve self-propagating high-temperature synthesis (SHS) or combustion synthesis [95] [93] was found to present a thermal conductivity close to conventional APS thermal barrier coatings [START_REF] Pedraza | Synthesis and related properties of low conductivity TBCs with hollow alumina microspheres, 43rd ICMCTF[END_REF]. The oxidation resistance of these coatings was investigated under isothermal [START_REF] Rannou | [END_REF]100] and cyclic conditions [63] in air and in water vapour [START_REF] Brossard | Influence de l'eau (vapeur, liquide) et du régime d'oxydation sur la dégradation de revêtements alumino-formeurs sur superalliage à base nickel[END_REF], in comparison with conventional diffusion coatings produced by pack cementation or by out-of-pack. The slurry coatings elaborated on pure nickel and nickel-based superalloys have shown faster oxidation kinetics than the conventional coatings [START_REF] Mollard | Elaboration de systèmes barrière thermique par barbotine. Comportement du nickel et de ses superalliages revêtus en oxydation cyclique à haute température[END_REF][START_REF] Rannou | [END_REF]100]. However, the ceramic top coat containing alumina hollow spheres was found to contribute the most to the oxidation kinetics. The presence of the ceramic top coat was found to limit rumpling of the oxide scale [START_REF] Rannou | [END_REF]. In cyclic conditions, the presence of substrate elements in the outer region of the coating due to the high Al-activity upon aluminizing was found to decrease the oxidation resistance [63]. The thickness of the alumina shells was particularly thin and the ceramic top coat presented poor erosion resistance. Therefore, the functionalization of the coatings developed under the PARTICOAT project was required for both military and civilian applications in aircraft industry.

VI.2. Functionalization of aeronautical thermal barrier systems elaborated by slurry (FONBAT)

The present study (Functionalization of aeronautical thermal barrier systems elaborated by slurry:

FONBAT) is funded by the DGA (Direction Générale de l'Armement -French Defense Ministery) and has potential applications in both military and civilian fields. Based on the works of Benoit Rannou [START_REF] Rannou | Slurry coatings from aluminium microparticles on Ni-based superalloys for hightemperature oxidation protection[END_REF] and Maël Mollard [START_REF] Mollard | Elaboration de systèmes barrière thermique par barbotine. Comportement du nickel et de ses superalliages revêtus en oxydation cyclique à haute température[END_REF] done at the LaSIE laboratory, the targeted improvements of the coating systems were to:

─ Control the Al activity upon aluminizing.

─ Increase the oxidation and hot corrosion resistance of the bond coat.

─ Increase the erosion resistance of the thermal barrier system.

─ Improve the thermal insulating effect of the ceramic top coat.

Since the previous works were focused on the elaboration of the coatings with simple aluminide coatings, the present study aims at elucidate the influence of doping agents into the slurry composition.

Different doping agents were susceptible to increase the performance of the coating systems:

─ Addition of Cr to control the Al activity upon aluminizing.

─ Addition of Cr to improve the hot corrosion resistance.

─ Mixtures of Al + Al2O3 to synthesize composite materials to enhance the erosion resistance.

Chapter 1

State of the art: context and motivations of the study -34-

The present study focused on the influence of Cr since this element has a potential beneficial effect on several features. First, the addition of Cr was expected to decrease the activity of Al upon aluminizing hence to promote the formation of a diffusion coating with a low-activity microstructure [61,65] (Fig.

I-11).

The Cr also has a potential beneficial influence on the hot corrosion and oxidation resistance of slurry aluminide coatings promoting the selective oxidation of Al.

In order to clearly identify the role of Cr addition in the slurry compositions, the reactivity of the microparticles was first investigated without metallic substrate (Fig. I-15). Considering the exothermic reactions usually involved for the formation of the slurry coatings, the reactivity of the microparticles with a metallic substrate was also studied. The influence of the atmosphere (Ar and air) was investigated on the formation of the coatings. The composition of the atmosphere exerts a particular effect on the mechanisms of oxidation of the microparticles and on their sintering in the ceramic top coat [START_REF] Pedraza | [END_REF].

Thanks to the versatility of the slurry process, different architectures were elaborated to tailor the coating with chemical and mechanical gradients. The combination of several doping agents was also investigated to design the coating with the applicative expectation in mind. After the selection of the most promising systems, the different potential improvements were evaluated. ). These (Hf,Ta,W)-rich carbides (according to EDS spot measurements and literature [5]) resulted from the dissociation of MC into M6C with a surrounding γ' envelope [8][9][10]. In the dendritic regions, the coarsening of γ' was observed with a rafted microstructure [6,9,11].

The carbides originally present in the substrate were used as natural markers to identify the initial surface of the substrate after aluminizing and thus to determine the coating growth direction [12,13].

Fig. II-2. (a) Microstructure of the CM-247 LC superalloy in an interdendritic region and (b)

zoom in the γ-γ' duplex microstructure (electrolytically etched with 10% H2C2O4).

-44-

II. Elaboration of the slurry coatings on metallic substrates

The slurry coatings developed in the present study were based on the environmentally-friendly composition developed during the European project "PARTICOAT" [14,15]. The elaboration of a slurry coating follows the sequence:

1. Surface preparation of the metallic substrate to be coated.

2.

Preparation of the slurry (organic binder + deionized water + metallic powders).

3.

Deposition of the slurry by air brush then drying of the slurry deposit.

4.

Heat treatment in Ar or in air to form the diffusion coating or the complete thermal barrier system.

The next paragraphs will present the experimental procedure for the elaboration of the slurry coatings once the metallic powders used in the present study characterized.

II.1. Characterization of the metallic powders

Different metallic powders were selected to investigate their reactivity before the preparation of the slurries (Table II-3). The Al powder supplied by Sibthermochim (Russia) was composed of spherical microparticles with a mean diameter of 8±2µm (Fig. II-3). These microparticles were elaborated by wire electrical explosion (WEE) [16,17] and were passivated to limit their reactivity prior to shipping [18]. Two different Cr powders supplied by Goodfellow were selected in the study. The fine Cr powder was composed of microparticles with a complex geometry and a mean diameter of 5±1µm (Fig. II-3). The microparticles of the coarse Cr powder were approximately ten times larger than that of the fine Cr powder (Fig. II-3). The influence of the size of the Cr microparticles on their reactivity with the Al ones was investigated according to their respective specific surface area (Chapter III). The use of the coarse Cr powder was found to be detrimental upon aluminizing in order to obtain homogeneous diffusion layers and the ceramic top coat. Therefore, the fine Cr powder was selected for the development and the production of the coatings on metallic substrates afterwards. -46-

II.2. Composition and preparation of the slurries

The different slurries used in the present study were composed of:

 Metallic powders (pure Al, pure Cr or Al-Cr mixtures)

 Polyvinyl alcohol (PVA) furnished by Sigma Aldrich as an organic binder soluble in water [19]  Milli-Q water as a solvent

The PVA was dissolved in milli-Q water at approximately 80°C under magnetic stirring. For all the slurry compositions, the PVA/water ratio was fixed at 1/10 in weight. Rannou et al. have shown that there was no influence of the ageing time for this composition up to at least 9 days [14]. For a good reproducibility, the slurry coatings elaborated for testing the reactivity with the substrate and for the development of the coatings (Chapters III, IV and V) were done with as-prepared solutions. To produce a larger number of samples (for oxidation and hot corrosion tests, Chapter VI), a more important amount of solution was prepared and used up to one week afterwards to save time and products. The ageing of the solution with time was found to have no influence neither upon aluminizing nor on the final microstructure of the coatings.

Thereafter, the metallic powders were added to the organic solution in a pill box with the proportions given in Table II-5. The proportions of the metallic powders for each composition were selected empirically after deposition tests to obtain an appropriate "visual" viscosity (not measured in this study) and a good coverage of the metallic substrate by air brush.

II.3. Deposition of the slurries by air brush on metallic substrates

Before applying the slurries, the samples were ground using a SiC #180 paper grade then rinsed with de-ionized water and finally ultrasonically cleaned in ethanol. This surface preparation improved the adherence of the microparticles-containing slurries onto the metallic substrates as well as their good coverage over the whole surface. The deposition of the slurries was realized using an air brush (AZTEK A4709T model) with a nozzle of 0.7 mm in diameter (Fig. II-5a). During the slurry deposition, the sample holder (Fig. II-5b) rotated with a constant speed of 5.25x10 -3 m.s -1 (1 rpm) [20]. The distance between the nozzle and the metallic substrates was fixed at 10 cm for better reproducibility. One can see the progressive drying of the slurry deposit upon the rotation of the sample holder in Fig. II-5b.

For reproducibility and to control the amount of slurry deposited on the metallic surfaces, the samples were weighed before and after spray with a 10 -4 g precision balance. -48-

As discussed in the paragraph VI.2 of Chapter 1, the addition of Cr in the slurry compositions was expected to play a beneficial role on different parameters. For simple Al slurry coatings, the aluminizing of the substrates was found to be fast with the formation of liquid phases because of highly exothermic reactions [21]. This resulted in the formation of high-activity coatings with a significant amount of substrate elements in the outer layer of the coating [22]. In order to decrease the Al activity upon aluminizing [23,24], three different Al-Cr architectures were envisaged (Fig. II-6). The mechanisms of formation of the coatings for these different architectures were investigated using pure Ni as a model material (Chapter III). Depending on the initial architecture, different coatings were successfully obtained with different chemical and mechanical gradients. The elaboration of the coatings was performed keeping in mind the required final application ("materials by design"). Therefore, the optimization of the compositions was carefully processed before extrapolation to industrial alloys. The main objective was to increase the performances of the coatings for high-temperature applications. -49-After deposition of the slurry, the samples were left at least one hour for drying in laboratory air before any other step [14]. The further formation of the coatings was realized by heat treatment in an inert atmosphere (Ar) or in air depending on the design of the coating (environmental coating or complete thermal barrier system).

II.4. Thermal treatments

Two different experimental set-ups were employed to perform the heat treatments.

 Thermogravimetric analysis (TGA):

A SETARAM TGA 92 thermobalance with a 10 -6 g accuracy was employed to study the evolution with time of the different slurry coatings on pure nickel (Chapter III). The evolution of the mass variation during the heat treatment was indeed helpful to detect the different events occurring on the sample (water evaporation, binder decomposition and oxidation of the microparticles). Before the heat treatment, the gas pressure in the reaction chamber was first decreased with a vacuum pump (down to 5x10 -2 mbar) and then argon was progressively introduced. Upon annealing, the reaction chamber was ventilated with an Ar flow (20 mL/min). The heating rate was fixed at 5°C/min and the cooling rate at 50°C/min to keep the microstructure developed at the annealing temperature to room temperature -50-for a better comprehension of the mechanisms of formation. All samples were systematically cured for 3 hours at 400°C in order to expel the organic binder. The different heat treatments performed in the TGA are summarized in Table II According to the supplier, small amount of impurities are present in the gas bottle (5 vpm H2O and 2 vpm O2). Referring to the Ellingham diagram (thermodynamic considerations), the content of these oxidizing species is sufficient to provoke the oxidation of Al and Cr at temperatures above 600°C. To investigate the oxidation and the sintering of the microparticles upon annealing (Chapter V), some heat treatments were also performed partially or completely in synthetic air (80 vol.% N2 + 20 vol.% O2).

Note that the heat treatment selected in previous studies in the group of F. Pedraza [20,25,26] was also performed to investigate the influence of the atmosphere (Ar and synthetic air) on the aluminization of pure nickel from Al microparticles (Chapter V).

 Horizontal furnace:

For the simultaneous production of several samples and the aluminizing of larger samples, a Carbolite horizontal furnace with a quartz tube was employed. In order to remove moisture, the reaction chamber was first purged with a vacuum pump down to approximately 1.0 mbar. Then, the argon flow was alternatively turned on and off for 30 minutes (three cycles in total) while keeping the vacuum pump on to clean thoroughly the furnace. After the procedure, the internal pressure in the furnace decreased down to less than 0.2 mbar. Then, the furnace was filled with argon at a pressure slightly over the atmospheric pressure. During the annealing, the furnace was ventilated with a flow of argon.

The same heat treatments than the TGA experiments were performed depending on the slurry coatings deposited on the samples. However, the cooling rate was slower than in the TGA because of the inertia of the furnace.

Chapter II

Experimental methods -51-

III. Environmental degradation of the slurry coatings

III.1. Isothermal and cyclic oxidation

The oxidation behaviour of the coated and uncoated samples was investigated under isothermal and cyclic conditions under atmospheric pressure. The samples were oxidized for short (3 to 100 hours) and long term exposures (up to 1000 hours). The different thermogravimetric experiments (continuous or discontinuous isothermal testing and cyclic oxidation testing) were expected to give an overview of the high-temperature resistance of the coatings elaborated during the study [27,28].

 Isothermal oxidation:

For short term exposure (up to 100 hours), isothermal oxidation tests were done in a SETARAM TGA 92 thermobalance in synthetic air (80 vol.% N2 + 20 vol.% O2) to continuously monitor the mass variation of the samples (10 -6 g accuracy). As previously described, the reaction chamber was first purged with a vacuum pump down to 5.10 -2 mbar then filled with Ar. The heating ramp was fixed at 50°C/min under flowing Ar (20 mL/min). Once the dwell temperature was reached, the reaction chamber was purged for about 8 seconds before the introduction of synthetic air (16 mL/min). At the end of the oxidation test, the samples were cooled down to ambient temperature at 50°C/min.

The continuous isothermal long term oxidation tests were realized up to 500 hours and 1000 hours using a Pyrox furnace in laboratory air. The mass variation of the samples was recorded with a Sartorius Quintix Secura 10 -5 g precision balance to extrapolate the short term measurements done in TGA.

 Cyclic oxidation:

To simulate thermal cycling conditions observed in aircraft gas turbines [27], cyclic oxidation tests were performed using a Delta Thermique cyclic rig (Fig. II-8a). This rig displays two independent furnaces able to work separately and to execute two different thermal cycling procedures Type-II hot corrosion conditions [29]. Since Na2SO4 is the most common salt deposit in aircraft gas turbines [30], this salt was selected in this study. In Type-II hot corrosion conditions, a sufficient partial pressure of SO3 is required in the gas flow to stabilize liquid sulfates [31,32]. Therefore, four iron zeolite catalysts were placed in the gas flow upstream of the samples to oxidize SO2 following reaction (Eq. II-1):

SO2(g) + O2 g → SO3(g) (Eq. II-1)
Each iron zeolite catalyst was in the form of a 15 cm long tube with a diameter of approximately 3 cm.

To increase the specific surface area in contact with the gas flow and optimize the conversion rate of reaction (Eq. II-1), the catalysts had a honeycomb shape. The calculated equilibrium 3 was 7.5x10 -3 atm. at 700°C from data in Ref. [30]. The gas flow rate was adjusted at 100 mL/min (6 L/h) with an overall pressure in the furnace slightly over 1 atm. This value is higher than the minimum 3 required for the formation of liquid phases at 700°C (3x10 -4 atm. according to Ref. [33]). Even though Na2SO4 is solid at 700°C, liquid phases can be observed in the presence of nickel and cobalt sulfates -54-(Te(Na2SO4-NiSO4)=671°C and Te(Na2SO4-CoSO4)=565°C [34]). According to the thermodynamic stability diagrams calculated by Gheno et al., the 3 of the present study was too low to form aluminium and chromium sulfates [34]. The temperature of the furnace was calibrated with a type-K thermocouple (Ni-Cr/Ni-Alumel). The "hot zone" of the furnace was approximately 14 cm long allowing the simultaneous treatment of up to 24 samples for each batch.

To simulate hot corrosion tests, two different experimental procedures are usually considered in the literature: salt spraying or salt recoating methods [29,[34][35][36][37][38][39] and half or full embedding in the corrosive salt [40][41][42][43]. These experiments progressively replaced the burner rig tests [44,45] at laboratory scale for practical and economic reasons. The half-embedding method (or crucible test method) was selected in this study since this method is highly useful for preliminary screening of materials and coatings [46].

With only half of the sample embedded in the salt, it is possible to characterize three distinct zones for one sample (atmosphere only, salt + atmosphere and the interface between the two zones) [43].

Depending on the initial material or coating, the size of the samples (12.7 mm diameter for pure Ni)

appeared to be a limitation to clearly differentiate these areas after long term exposure (> 100h). With the crucible test method, the analysis of the salt was also possible after the experiments. For reproducibility, the same amount of Na2SO4 was poured in the alumina crucibles (approximately 2.5 g of salt).

The hot corrosion tests were performed for 20h, 100h and 300h of exposure to describe an evolution and investigate the corrosion mechanisms. To emphasize the influence of the atmosphere and the presence of Na2SO4 on the corrosion rate, complementary tests were done and are summarized in The DSC technique is able to detect thermal events occurring in the investigated system by comparison with a reference sample. The DSC analyses were carried out using a SETARAM Labsys Evo 1600 thermal analysis. The measurements were performed in flowing Ar (20 mL/min) to limit oxidation phenomena and prevent secondary reactions in the system. For the different systems investigated, energy changes and temperatures of reaction or transformation were measured with the Calisto software coupled to the DSC instrument. Different heat treatments were performed, using alumina crucibles (90 µL), depending on the system (Table II-8). The alumina crucibles were systematically cured at 1350°C before DSC tests to avoid any contribution from the crucible on the measurements.

The DSC was particularly helpful to analyse the reactivity between Al and Cr powders without and with substrate. Since the thermal barrier coatings are applied to metallic components to decrease the metal temperature, the determination of their thermophysical properties is of utmost importance to ensure the integrity of the underlying metal substrate. Indeed, the thermal conductivity of TBC coatings together with the coating thickness control the temperature drop between the atmosphere and the metal-coating interface [49,50]. The porosity of the coatings has a particularly great influence on the thermal conductivity values of the full thermal barrier system [51][52][53]. Increasing the porosity lowers the thermal conductivity values which attests the attractive design of the thermal barrier system elaborated by slurry coatings [15,20,25,54].

The fundamental relationship to determine the thermal conductivity is given in equation (Eq. II-2):

= • • (Eq. II-2)
With : the thermal conductivity, : the thermal diffusivity, : the specific heat capacity and : the density.

The determination of the thermal diffusivity of the slurry coatings was carried out using a laser flash technique [49]. The laser flash technique is an ASTM standard method (E1461) [55]. A schematic of this technique is given in Fig. II-10a. The test involves:

-subjecting the entire front surface of a sample to a short duration radiant energy pulse from a laser (less than 2 ms usually),

-recording the rear face temperature rise after absorption of the pulse by the front surface.

The thermal diffusivity is calculated from the sample thickness and the time required to the heat from the front surface to reach a certain percentage of the maximum temperature at the rear surface of the sample [49]. The determination of the thermal diffusivity is given by equation (Eq. II-3):

= • ² (Eq.

II-3)

With : the thermal diffusivity, : a constant corresponding to the percentage of the maximum temperature rise, : the sample thickness and : the time required for the rear surface temperature to reach the percentage.

-57-

The laser flash measurements were performed with a LINSEIS LFA 1000 (Fig. II-10b). The thickness of the samples was measured with a 10 -3 mm caliper. The samples with a metallic substrate were generally about 1 mm thick whereas the self-standing coatings measured less than 1 mm thick. To avoid heat losses, both surfaces of the samples were sprayed with graphite (Graphit 33, Kontakt Chemie). For self-standing coatings, a gold layer was deposited by sputtering before graphite spray for a better absorption of the laser pulse [55]. With the experimental set up, up to six samples can be simultaneously analysed. A vacuum of approximately 5.10 -2 mbar was maintained in the reaction chamber during the measurements.

The measurements were done at different temperatures to describe the evolution of the thermal diffusivity of the coating with the working temperature. For repeatability, five laser pulses were done at each temperature for each sample. Between each measuring temperature, the heating ramp was fixed at 10°C/min and a 10 minutes dwell was done before the first laser pulse to stabilize the temperature of the samples (given with a type-S thermocouple placed at the centre of the sample holder). The frequency of the laser pulses was fixed at 1 minute. Each thermogram was saved after the laser pulse and the thermal diffusivity was directly calculated by the software. For validation, the fitting curve must be accurate to at least 95%. The X-ray diffraction (XRD) allows to determine the crystallographic structure of mineral and metallic compounds. Combined with the SEM or the Raman micro-spectrometry, XRD is particularly helpful to differentiate products with the same chemical composition but different crystal structures. The surfaces of as-deposited, as-coated and degraded samples were analysed by X-ray diffraction (XRD) using a Bruker AXS D8 Advance with Cu Kα1 radiation (=0.15406 nm). The XRD patterns were acquired for a 2θ range between 10° and 90° with a step size of 0.02° (DIFFRAC plus XRD Commander software). The accumulation time was fixed at 5s for the 5° grazing incidence procedure and 2s in Bragg-Brentano configuration (θ-2θ mode). The first one allows to collect information from the uppermost surface (e.g. thermally grown oxides) whereas the symmetrical mode makes X-rays penetrate deeper into the bulk [25]. The identification of the crystal structures for the different metallic or intermetallic phases and the oxides was realized using the DIFFRAC plus EVA software by comparison with the JCPDS database (Joint Committee Powder Diffraction files System). Some crystal structures were identified by comparison with data from the literature.

b) Raman micro-spectrometry and Cr 3+ fluorescence

Raman micro-spectrometry is a non-destructive method for local structural analyses (lateral resolution of approximately 1µm [56]). This technique allows the characterization of the molecular composition and of the crystalline structure of non-metallic compounds from their vibrational transitions (oxides, sulfates and sulfides for instance). For this, a monochromatic radiation (usually from a laser) is directed to the sample with a known frequency corresponding to the emission of the photons. By interaction with the probed area, these photons are either reflected, absorbed or diffused. Some of the photons are diffused non-elastically and present therefore a different frequency (Raman scattering). The nonelastic diffusion is characteristic of an energy change between the incident photon and the compound, corresponding to a transition between two vibrational energy levels [57]. Depending on the frequency (or energy) of the diffused photons compared to incident ones, the scattering can be qualified as Stokes or as anti-Stokes [57]. The frequency is usually converted in wave number ̅ (cm -1 ) for the acquisition and the presentation of the results considering the equation (Eq. II-4):

̅ = = 1 (Eq.

II-4)

With : the frequency, : the speed of light and : the wavelength of the emission source

-59-

In this work, a high-resolution micro-spectrometer Jobin Yvon Horiba (LabRam HR8000 model) with a He-Ne monochromatic laser (=632.817 nm) was employed. An Olympus confocal microscope is coupled to the micro-spectrometer to visualize the sample and focus the laser (10x, 50x and 100x magnifications). The data are collected with a CCD detector, cooled by Peltier effect, with a spectral resolution of approximately 0.5 cm -1 . Different filters can be used to adjust the intensity of the laser depending on the products investigated. The depth of the analyses is comprised between 1 and 5µm according to the setup parameters and the sample.

The analyses were performed using the 50x magnification or 100x magnification to characterize thin oxide scales. Since the Raman micro-spectrometry is a local technique, several measurements were realized on different locations of the surface for each sample for repeatability. Depending on the wave number range, different signals can be obtained by Raman spectrometry. The Raman measurements were usually done between 100 and 2000 cm -1 . The intensity of the laser source and the exposure time were adapted for each system investigated to optimize the measurements. In the range 1000-2000 cm -1 , fluorescence signals can be obtained. The fluorescence spectrometry is due to the excitation of a fluorophore compound. This method of detection is particularly helpful to characterize the nature of thermally grown oxide (TGO) scales and their stress rate [58][59][60]. Indeed, the presence of small amounts of Cr 3+ in the Al2O3 matrix favours the appearance of the different polymorphs of alumina by fluorescence [61,62]. For instance, it is possible to differentiate the θ phase from the α phase of alumina In DECHEMA-Forschungsinstitut, the samples were characterized with a LEICA DMRM microscope coupled to a LEICA DFC450 camera. This microscope was used for the metal loss measurements after the hot corrosion tests. Up to twelve micrographs were taken at x100 magnification on different areas of the samples to do the measurements. This probabilistic approach was particularly interesting to estimate the degradation of the coatings in hot corrosion conditions.

b) Scanning Electron Microscopy (SEM) and Energy Dispersive X-ray Spectroscopy (EDS)

With the SEM, a sample is scanned with a focused beam of electrons and the incident energy is transferred to the atoms of the material increasing their potential energy. After being excited, the atoms return to their fundamental state by de-excitation in the form of a secondary radiation.

Depending on the activation energy, different radiations are likely to be observed. The formation of images is due to a contrast difference coming from the topography or a difference of chemical composition. Two imaging modes are usually considered in a SEM:

 Secondary electrons imaging (SEI): secondary electrons are produced by the repulsion of the electrons of an atom by the incident electrons (inelastic interaction). With a low energy (< 50eV), only the secondary electrons extracted from the surface can be detected. Therefore, the SE mode gives a good topography resolution.

 Back-scattered electrons imaging (BSEI): back-scattered electrons result from the elastic interaction of the incident electrons with the nuclei of atoms. The energy released by this interaction is a function of the atomic number. Therefore, this mode mainly produces a contrast of chemical composition. Heavier elements appear brighter whereas lighter elements appear darker in BSE mode.

-61-Surface and cross-section analyses were carried out with an environmental microscope FEI Quanta 200F equipped with a Schottky field emission gun (FEG). For cross-section observations, the samples were first mounted in a polymeric resin and progressively grinded with finer SiC papers up to #4000 grade then polished with a 1µm water based diamond suspension (Struers). The samples were finally polished with a colloidal silica suspension (Struers OP-S 0.04µm) for a better revelation of the microstructures. The energy of the electron beam was fixed at 20 kV with a working distance of 10 mm and a spot size between 3.4 and 3.7. With conductive samples, the analyses were performed in high vacuum (P < 10 -5 bar). To limit the charge effects of low-conductive samples (the ceramic top coat for instance) and avoid further surface preparation, the analyses were performed in low vacuum (0.9 mbar). For elemental analyses, the SEM is coupled to an EDAX detector. Energy dispersive X-ray spectroscopy was carried out to analyse the chemical composition of the products by X-rays microdiffraction. Depending on the energy of the emitted photons X, this technique is able to differentiate an element from another. In this study, spot analyses, window analyses and map analyses were performed depending on the investigated sample. EDS map analyses give the distribution of the elements in the scanned area by counting the probability to find a distinct element for a given position. This is particularly helpful in the case of diffusion coatings and heterogeneous compounds with concentration gradients of the different elements.

In DECHEMA-Forschungsinstitut, a Philips XL40 microscope combined with an EDAX detector was carried out for SEM and EDS characterization. To perform the analyses in high vacuum, the samples were sputtered with graphite and the edges were coated with a silver paste to ensure a good conductivity.

c) Electron Probe Micro-Analysis (EPMA)

The EPMA characterization was done in DECHEMA-Forschungsinstitut by Dr. Gerald Schmidt and Dr.

Melanie Braun during my three months stay in their premises. The EPMA is coupled to a scanning electron microscope JEOL JXA-8100 and gives high resolution (about 1 µm) elemental maps. The energy of the electron beam was fixed at 15 kV and the analyses were performed in high vacuum. To avoid charge effects, the samples were sputtered with graphite and the edges were coated with a silver paste.

- microstructure comparable to that of low-activity coatings produced by CVD related techniques [6,7].

Cr addition was found to control the Al activity, hence to promote the outward diffusion of nickel to form the coating at elevated temperature (i.e. 1100°C). However, this requires a previous mechanical attrition of the Al and Cr powders, hence an additional manufacturing step. In addition, and despite their expensive use in current aluminizing processes, either by CVD or by slurry, little is known about the reactivity between Al and Cr that controls the aluminizing of metallic substrates.

Therefore, the role of introducing Cr in the slurry for aluminizing purposes was studied in this chapter. phase (Al11Cr4) at about 75 at.% Al [12]. The latter would decompose into Al4Cr and Al8Cr5 at 829°C [17]. These Al-Cr phases were often investigated through differential thermal analysis (DTA) and the mechanisms of formation proposed after homogenisation in liquid phase and subsequent cooling for different Al-Cr compositions [10,11]. Other studies by Chen et al. focused on the sintering of Al-Cr alloy targets by different powder metallurgy methods using micro-sized Al and Cr particles [18].

Depending on the method and the temperature reached, different intermetallic compounds were synthesized (Al11Cr2, Al8Cr5 and AlCr2). Therefore, despite the good knowledge of the Al-Cr system,

Fig. III-1. Binary Al-Cr phase diagram [9]

.

-68-

no study has been carried out so far to investigate the reactivity between Al and Cr microparticles, which is of utmost importance to control the self-propagating high-temperature synthesis (SHS)

reactions upon slurry aluminizing. The formation of intermetallic phases by SHS can be predicted by the amount of energy released during the exothermic reaction and the maximum temperature reached [13,19,20].

Table III-1. Selected compositions for DSC tests in weight % and atomic % and corresponding

Al activity in the powder mixtures. For consistency, all compositions will be presented in weight% in the study.

For the DSC tests, the experimental trials were made by weighing (with a 10 -4 g accurate balance) and mixing the selected compositions of Al and Cr powders (Table III-

I.2. Reactivity of Al and Cr powders measured by DSC

The DSC thermograms obtained for the heating ramp from 40°C to 1300°C at 10°C/min for the selected Al-Cr compositions are given in Fig. III-3. For the two systems, an exothermic reaction was identified at a temperature lower than the experimental melting point of Al for all Al-Cr compositions.

This suggests that the transformation should be of the eutectic type as proposed by Mahdouk and Gachon [24]. Therefore, this transformation would occur at a definite chemical composition (as soon as a small percentage of Cr reacts with Al due to its limited solubility in (Al) [24,25]). The onset temperature of this reaction was determined at 637 ± 2°C by mixing Al powder with the fine Cr powder (Fig. III-4a). Using the coarse Cr powder, the onset temperature was shifted to 650 ± 2°C (Fig.

III-4b

). This difference was ascribed to the greater contact surface, hence the reactivity, of the Al microparticles with the fine Cr powder compared to the coarse Cr powder [13]. Nevertheless, the amount of heat released during the exothermic reaction appeared to be dependent on the initial Al-Cr composition both systems (Fig. III-4a and Fig. III-4b). The greatest release of heat arose from the 67Al33Cr composition with an exchanged heat of approximately -370 J/g with the fine Cr powder and -308 J/g with the coarse one. For greater Cr contents (≥ 45 wt.% Cr), the heat quantity progressively decreased as the Cr content increased.

Fig. III-2. DSC thermograms of (a) the pure Al powder and (b) the fine and coarse Cr powders (heating and cooling ramps of 10°C/min) under flowing Ar (20 mL/min).

-70-

Fig. III-3. DSC thermograms (heating branches, 10°C/min) for the different Al-Cr compositions mixing the Al powder with (a) the fine and (b) the coarse Cr powder (20 mL/min Ar flow).

Fig. III-4. Evolution of the heat released for the different Al-Cr compositions measured for the exothermic reaction (20 mL/min Ar flow) for mixtures of Al powder with (a) fine and (b) coarse Cr powder. Table III-2. Experimental data extracted from the DSC tests (heating branches) for the different Al-Cr compositions. The values given refer to the mixtures of Al with the fine ("F") and coarse ("C") Cr powders.

L + Al4Cr ↔ Al11Cr2 895 [24] 902 [25] 67Al33Cr (F) 55Al45Cr (F) 44Al56Cr (F) 1037 ± 1 1041 ± 1 1043 ± 1 L + Al8Cr5 ↔ Al4Cr
1040 [9] 1037 [24] 1035 [25] 67Al33Cr

(C) 55Al45Cr (C) 44Al56Cr (C) 30Al70Cr (C) 1035 ± 1 1041 ± 3 1044 ± 5 30Al70Cr (F) 20Al80Cr (F) 12Al88Cr (F) 882 ± 1 910 ± 1 917 ± 2 Convolution (Cr) ↔ Al8Cr5 + AlCr2 (Cr) ↔ AlCr2 882-910 [9] 874-906 [25] 870-910 [24] 30Al70Cr (C) 20Al80Cr (C) 12Al88Cr (C) 891 ± 6 908 ± 2 914 ± 2
For the 82Al18Cr composition, an endothermic peak was observed between 820-840°C (Fig. occurring in a very narrow range of temperature [9,25].

-72-From these experiments, three Al-Cr compositions were selected for the characterization of the microstructures after the exothermic reaction:

-The 67Al33Cr composition with the highest energy change during the exothermic reaction.

-The 44Al56Cr composition with an intermediate exchanged heat.

-The 20Al80Cr composition with a low exchanged heat.

The main objective was to characterize the intermetallic phases formed during the exothermic reaction and the potential transformation of these phases upon annealing at different temperatures.

For the rest of the study, only the fine Cr powder was considered since the coarse Cr powder was not suited to obtain the desired coating design.

Fig. III-5. (a), (b) DSC heating branches (10°C/min) and (c) evolution of the exchanged heat and the Tonset for the peritectic transformation of the Al4Cr phase for the different Al-Cr compositions.

-73-

II. Synthesis of Al-Cr intermetallic compounds

The three selected compositions (67Al33Cr, 44Al56Cr and 20Al80Cr in wt.%) were heated to 650°C at 5°C/min, held at this temperature for 10 minutes and then quenched at 50°C/min to freeze the microstructure and the crystal phase of the compounds synthesized during the exothermic reaction.

The SEM images of the three Al-Cr mixtures before and after heat treatment in DSC are given on Fig. synthesized phases (Table III-3). In conjunction with the XRD analyses and the literature, the main intermetallic phases were Al11Cr2 for the 67Al33Cr and 44Al56Cr compositions and Al11Cr4 for the 20Al80Cr one. The exothermic reaction identified at 637 ± 2°C by mixing Al and fine Cr powders (Fig. III-3a) was associated with the formation of AlxCry intermetallic compounds with different stoichiometry depending on the maximum temperature reached during the synthesis. Chen et al. reported that the Al11Cr2 phase could form at temperatures (500°C) lower than ours but this required the 55Al45Cr powders to be hot isostatically pressed [18]. In such 55Al45Cr mixture of powders, nearly all the Al had reacted with Cr to form the Al11Cr2 phase at 600°C. The formation of such Al-rich intermetallic

Table III-3. Summary of the EDS spot measurements (Fig. III-6) and XRD analyses (Fig. III-7b) for the different Al-Cr mixtures after annealing (650°C for 10 minutes) and the suggested phase compositions of synthesized products.

Fig. III-7. XRD patterns for the three Al-Cr mixtures (a) before annealing and (b) after annealing at 650°C for 10 minutes (20 mL/min Ar flow).

compounds at temperatures lower than the Al melting point was related to the high exothermic reactions between Al and Cr [14] and Al and Cr2O3 [27]. The local increase in temperature due to the formation of intermetallic phases makes the residual Al to melt and the obtained liquid Al progressively dissolves Cr. Since the adiabatic temperature of the mixtures is not reached, the reaction is selfsustained as far as none of the reagents is exhausted. This is commonly known as self-propagating high temperature synthesis [14] or combustion synthesis if oxygen is present in the reacting media [27]. In

Fig. III-8. Evolution of the heat flow around the exothermic reaction for different Al-Cr compositions (650°C-10 min + quench). The recrystallization of liquid Al was only observed

for the 82Al18Cr composition.

- phases, the Al4Cr compound was also expected to form. For the intermediate composition (44Al56Cr), the maximum temperature should be at least 1000°C according to Naplocha et al. [13]. However, the presence of both Al11Cr2 and Al4Cr in the microstructure suggested that the increase in temperature was local and did not lead to a full transformation of these phases. The mechanism of formation for the Al11Cr4 phase was not fully understood in the present study and the formation of this phase is still under debate in literature [12]. Finally, for the 20Al80Cr composition, Al4Cr was not observed in the microstructure probably because the local increase in temperature was higher than the peritectic transformation of Al4Cr in the 1035-1040°C range. This would explain the detection of the Al8Cr5 phase in the microstructure after the quench. The sequence L + Al7Cr → L + Al11Cr2 → L + Al4Cr for the SHS reaction between Al and Cr microparticles proposed by Naplocha et al. [13], seems to describe very well the mechanisms of formation in the present study.

Fig. III-9. SEM images (BSE mode) of the 82Al18Cr composition (a) before and (b) after annealing at 650°C for 10 minutes (20 mL/min Ar flow).

-77-III. Reactivity of Al, of Cr and of Al-Cr mixtures with pure nickel substrate

III.1. Reactivity of the metallic powders with pure nickel substrate

To investigate the mechanisms of formation of the slurry diffusion coatings on metal substrates, it is necessary to identify the reactions occurring between the metallic powders and the underlying substrate. Therefore, pure nickel substrate was employed as a model material for nickel-based superalloys. The reactivity of the metallic powders with pure nickel substrate was studied by DSC.

Fig. III-10. Principle of the DSC tests to study the reactivity between (a) pure nickel substrate with (b) the metallic powders (without organic binder) or (c) slurry deposits.

As a reference test for high-activity coatings [1], the reactivity between pure nickel and Al powder was first investigated. Since the mechanisms of formation of the slurry diffusion coatings were studied from slurries deposited on metallic substrates (cf. IV), the same conditions were followed for this study (Fig. Nevertheless, the exothermic signal characteristic of the reaction between nickel and Al [1] was not clearly identified. Two other endothermic transformations were observed for this system (Fig. III-12a) at 861 ± 2°C (NiAl3 melting [5,20]) and 1134 ± 2°C (Ni2Al3 melting [5,20]). There was no recrystallization upon cooling suggesting that the whole Al reacted with the nickel substrate. This system (Fig. III-10b) was expected to be less representative of what exactly occurred during slurry aluminizing than the experiment presented in Fig. III-10c.

Fig. III-11. DSC thermograms of a pure nickel sample (heating and cooling ramps of 10°C/min, 20 mL/min Ar flow).

Considering the reactivity between pure nickel and the Al slurry deposit (Fig

III.2. Reactivity of Al-Cr mixtures with pure nickel substrate

The reactivity of Al-Cr mixtures with pure nickel substrate was investigated and compared with the DSC measurements realized for the Al-Cr mixtures with no substrate. The main objective was to determine if the addition of Cr to the slurry limits or supresses the exothermic reaction between Ni and Al. The three Al-Cr slurry compositions (Table III- The addition of Cr in the slurry compositions limited the SHS reactions between Al and the nickel substrate usually observed for high-activity slurry coatings [1].

After the exothermic reaction, a variation of the heat flow was observed with a broad endothermic signal between 650°C and approximately 1000°C (Fig. . This variation was not observed for the nickel reference sample and the different Al-Cr mixtures (with no substrate). Therefore, this variation was characteristic of an interaction between the nickel substrate and the slurry deposit. This broad endothermic signal was more energetic for the 44Al56Cr composition and may be attributed to the interdiffusion between Al-Cr intermetallic compounds present in the slurry deposit and the nickel substrate.

Fig. III-14. DSC thermograms for the pure nickel substrate in contact with (a) the 20Al80Cr and (b) the 44Al56Cr slurry deposits.

III.3. Synthesis of Al-Cr intermetallic compounds from Al-Cr mixed slurries for aluminizing

The reactivity of Al-Cr microparticles in the deposited slurries was investigated for the three selected compositions (Table III-4) in order to investigate the role of Cr addition in slurry compositions upon aluminizing. The slurries were deposited on pure nickel substrate (12.7 mm diameter) by air brush.

The deposited slurry quantity was greater for the Al-Cr slurries than for the pure Al (Table III-4) in order to have a sufficient Al content in the donor for diffusion. The samples were systematically cured for 3 hours at 400°C to remove water and expel the organic binder before any further step [30]. The heat treatments were performed in TGA under flowing Ar and were as follows:

-650°C-10 min: repeats the heat treatment performed in DSC for reproducibility.

-650°C-5h: diffusion step for the homogenization of intermetallic phases and comparison with pure Al slurry (high-activity).

-650°C-5h + 1000°C-3h: complete heat treatment to promote the outward diffusion of Ni at 1000°C and stabilize the -NiAl phase.

-1080°C-6h: conventional "low-activity" heat treatment for the repair of aluminide coatings on LPT blades and vanes made of nickel-based superalloys.

The SEM cross-section images (Fig.

Fig. III-15. BSE cross-section images for the 67Al33Cr, 44Al56Cr and 20Al80Cr slurry compositions deposited on pure Ni after annealing at 650°C for 10 minutes and 5 hours in TGA (20 mL/min Ar flow).

-83-III-6). In the 67Al33Cr mixture, coarse intermetallic phases formed (Fig. with a chemical composition close to Al4Cr according to the EDS measurements (81 ± 1 at.% Al) [9,16]. Only small remnants of Cr were revealed in the microstructure whereas an additional phase (i.e. Al11Cr2) was identified by XRD. With this 67Al33Cr mixture, the microstructure and the composition of the intermetallic phases did not substantially change with increasing the annealing time to 5 hours (Fig.

Fig. III-17. BSE cross-sections images for (a) the 67Al33Cr, (b) 44Al56Cr and (c) 20Al80Cr slurry compositions deposited on pure nickel and (d) corresponding XRD patterns of the surfaces after the complete heat treatment (650°C-5h + 1000°C-3h) under Ar.

 Intermediate summary:

After annealing of the powders at different temperatures (650°C and 1000°C), the intermetallic compounds became progressively richer in Cr. For instance, after 5 hours of annealing at 650°C, the AlxCry intermetallics were Al4Cr and Al11Cr2 for the 67Al33Cr mixture, and Al8Cr5 and Al11Cr4 for both the 44Al56Cr and 20Al80Cr mixtures. Such Cr enrichment was fostered after the final step at 1000°C (i.e. Al8Cr5 and AlCr2) and is believed to result from solid state diffusion (homogenization of the chemical composition between Al and Cr). Depending on the stoichiometry of the Al-Cr intermetallic phases synthesized, the thermodynamic activities of Al and Cr on the substrate surface may be different [21]. The intermetallic compounds synthesized in the slurry deposits can be assimilated to masteralloys in the aluminizing process. The main difference when using the slurries is that the masteralloys are in direct contact with the metallic substrate and not decomposed in a gas phase from halides activators by comparison with CVD-related techniques [21,31,32].

Based on these results, the formation of the identified Al-Cr intermetallic phases in the slurry deposits was therefore expected to decrease the Al activity on the surface of the substrate prior aluminizing. The objective was to limit the inward diffusion of Al in the substrate at low temperature while promoting the outward diffusion of nickel at high temperature to produce a coating with a and the diffusion layers remained thin and discontinuous.

 Conventional low-activity heat treatment (400°C-3h + 1080°C-6h):

Three nickel samples were coated with the 67Al33Cr, 44Al56Cr and 20Al80Cr slurry compositions and heat treated in TGA following the conventional low-activity heat treatment employed in the repair industry. The SEM cross-section images of the samples and corresponding X-ray maps are given in Fig. According to the X-ray maps, the nickel substrate was enriched with Cr for the 44Al56Cr and 20Al80Cr slurry compositions. Oxygen was found to be mainly associated with Cr and Al. This would be explained by the greater affinity for oxygen of Cr and Al than that of Ni.

III-

Fig. III-20. BSE cross-section images and X-ray maps for the (a) 67Al33Cr, (b) 44Al56Cr and (c) 20Al80Cr slurries elaborated on pure nickel after heat treatment in TGA (400°C-3h + 1080°C-6h

).

- system, the 33Al-67Cr and 50Al-50Cr compositions refer respectively, to the system with 5 and 10 g.cm -2 of Al. For the Cr-Al double-layer system, the 67Cr-33Al and 50Cr-50Al compositions refer respectively to the system with 5 and 10 g.cm -2 of Al.

V.2. Mechanisms of formation for the Al-Cr double-layer system  Diffusion heat treatment (400°C-3h + 650°C-5h):

To

Fig. III-22. BSE cross-section image and X-ray maps of Cr, Al and Ni for the 50Al-50Cr doublelayer system elaborated on pure nickel in DSC (20 mL/min Ar flow).

Fig. III-23. BSE cross-section image and EDS spot measurements of the diffusion zone identified in Fig. III-22.

and γ'-Ni3Al layers presented a thickness close to that of the Al slurry coating, the formation of the brittle Ni2Al3 phase was considerably limited for the Al-Cr double-layer system at 650°C. These observations suggested that the Cr top layer was able to limit the diffusion of Al toward the substrate by forming AlxCry intermetallic phases in the slurry deposit. However, the presence of the Ni2Al3 phase suggested that the SHS reactions between Al and Ni were not suppressed.

 Influence of the annealing time at the diffusion temperature (i.e. 650°C):

Complementary experiments were done in TGA in the same conditions. The annealing treatment at

Fig. III-24. BSE cross-sections images of the diffusion layers formed after annealing at 650°C for the Al-Cr double-layer systems on pure nickel (TGA, 20 mL/min Ar flow) and corresponding XRD surface patterns (without top coat).

-94-

The addition of Cr above the Al slurry layer was able to limit the thickness of the diffusion zone and the formation of the Ni2Al3 phase on nickel substrate. However, the SHS reactions between Ni and Al [20,39] were not suppressed for the Al-Cr double-layer system because of the initial contact between Al microparticles and pure nickel substrate (Fig. . Therefore, the diffusion zones were found to be heterogeneous in thicknesses and composition and an additional step at a higher temperature (i.e.

1000°C) was required to promote nickel outward diffusion [22].

Above the diffusion zone in the top coat (Fig was found to cause spallation of the top coat [43]. An excess of oxidized Cr microparticles was present in the top layer after annealing.

 Diffusion heat treatment (400°C-3h + 650°C-5h):

To investigate the mechanisms of formation for the Cr-Al double-layer system on pure nickel, a first sample was submitted to the diffusion heat treatment in TGA. The SEM cross-section image and X-ray

Fig. III-29. BSE cross-section image and corresponding X-ray maps for the 50Cr-50Al doublelayer system elaborated on pure nickel after the diffusion heat treatment in Ar.

Fig. III-31. XRD patterns of the surfaces for the 50Cr-50Al double-layer system elaborated on pure nickel after (a) the diffusion and (b) the complete heat treatment in TGA.

Fig. III-30. BSE cross-sections images and EDS concentration profiles for the Cr-Al doublelayer systems elaborated on pure nickel after the diffusion heat treatment in Ar.

-100-

 Complete heat treatment (400°C-3h + 650°C-5h + 1000°C-3h):

After the complete heat treatment, a diffusion zone was formed in the outer region of the nickel substrate (Fig. . This diffusion zone was rich in Ni and Al suggesting the formation of NixAly intermetallics. Cr-rich precipitates within the diffusion layer were also emphasized on the X-ray maps. from the greater thermodynamic stability of the latter, whose enthalpies of formation are much greater than the NixCry ones [44]. Subsequently, a large fraction of Cr remained undissolved in the newly formed NixAly phases because of its low solubility [36][37][38]. The exact composition of the Cr-rich phases was not determined from the SEM-EDS analyses but suggested the presence of α-(Cr).

Fig. III-32. BSE cross-section image and corresponding X-ray maps for a 50Cr-50Al doublelayer system elaborated on pure nickel after the complete heat treatment in TGA.

-101-

In contrast, the Al content of the Al-Cr intermetallic phases formed above the substrate with the 67Cr- By comparing 50Cr-50Al and 67Cr-33Al double-layer systems, the sequence of reaction can be described following equations (Eq. III-1) and (Eq. III-2):

xAl + yCr → AlxCry (Eq. III-1)

Ni + AlxCry → -NiAl + α-(Cr) (Eq. III-2)
The stoichiometry of the AlxCry phases synthesized in the slurry deposit is dependent on the initial contents of Al and of Cr in the reaction (Eq. III-1). The driving force of the reaction (Eq. III-2) is therefore the initial Al content of the slurry deposit. The transformation (Eq. III-2) implies that -NiAl is readily formed and that the outward diffusion of nickel prevails upon the aluminizing process. This is characteristic of low-activity coatings [48,49]. Based on these observations, it appeared therefore that the simplified transformation (Eq. III-2) was controlling the growth of the coating for the 50Cr-50Al system.

Fig. III-33. BSE cross-sections images and EDS concentration profiles for the Cr-Al doublelayer systems elaborated on pure nickel after the complete heat treatment in Ar.

-102-On the opposite, the large excess of α-(Cr) after the exothermic reaction for the 67Cr-33Al compositions considerably decreased the driving force for diffusion and transformation (Eq. III-2). The growth of the coating was therefore considerably hampered. This would require longer time or higher temperatures to foster the interdiffusion between the nickel substrate and the Cr-rich phases (denoted α-(Cr) in the reaction). These observations confirmed that the control of the Al-Cr ratio is of utmost importance to obtain the desired coating design. Considering the Cr-Al double-layer system, it is therefore possible to adjust the deposited mass of Cr and Al to optimize the final microstructure of the coating.

 Mechanisms of formation for the Cr-Al double-layer system:

The sequence for the mechanisms of formation of the Cr-Al double-layer system elaborated on pure nickel is proposed in -104-

 Conventional low-activity heat treatment (400°C-3h + 1080°C-6h):

Pure nickel samples slurry coated with the Cr-Al double-layer system were also submitted to the conventional "low-activity" heat treatment employed in the repair industry before extrapolation to nickel-based superalloys. Two compositions were tested (50Cr-50Al and 67Cr-33Al) to characterize the microstructure and confirm the hypotheses of the outward growth of the coating.

For the 67Cr-33Al composition, a continuous -NiAl layer was formed above the nickel substrate with a thickness of approximately 15-20µm (Fig. III-35a). The delamination between the nickel substrate and the -NiAl layer suggested that the outward diffusion of nickel was more active than the inward diffusion of Al and of Cr. Indeed, nickel from the substrate diffused in the slurry deposit and reached 60 at.% Ni at the extreme surface of the diffusion zone (Fig. III-35c). Some α-Cr precipitates were highlighted by X-ray map in the -NiAl phase because of the low solubility of Cr in the -NiAl phase [36,37]. Above the -NiAl layer, a Cr-rich layer was observed (more than 20 at.% Cr) and confirmed that the formation of NixAly phases was thermodynamically more stable than the formation of NixCry ones. For the 50Cr-50Al composition, the -NiAl layer was thicker with a thickness of approximately was investigated. The addition of Cr in the slurry compositions was found to limit the SHS reactions usually observed between Al and the nickel substrate for high-activity slurry coatings.

Fig. III-35. (a), (b) BSE cross-section images and X-ray maps for the Cr-Al double-layer systems elaborated on pure nickel and (c), (d) corresponding EDS concentration profiles after the conventional low-activity heat treatment in

 Influence of Cr addition upon aluminizing:

Upon aluminizing, Al preferentially reacted with Cr in the slurry rather than with the pure nickel substrate, thereby limiting the SHS reactions between Al and Ni at low temperature (i.e. 650°C).

Therefore, unlike simple Al slurry coatings, Al-rich nickel aluminides were barely formed on the nickel substrate at 650°C when using Al-Cr mixed slurries. Depending on the initial Al-Cr ratio of the slurry, the AlxCry phases synthesized in the slurry deposit released the Al to grow a diffusion zone in the substrate after annealing at high temperature (i.e. 1000°C and 1080°C). At these temperatures, nickel could diffuse outwardly and react with the Al-Cr intermetallic phases formed above the substrate. The dissolution of Al-Cr phases by nickel resulted in the formation of the -NiAl phase. A fraction of Cr was found to be in solid solution in -NiAl whereas α-Cr precipitates were also observed in the microstructures. With pure nickel substrate, only the 67Al33Cr composition allowed to grow a continuous -NiAl layer with the conventional repair "low-activity" heat treatment (400°C-3h + 1080°C-6h). Richer Cr contents in the Al-Cr mixtures impeded a homogeneous supply of Al to the nickel substrate.

 Mechanisms of formation of the functional coating designs:

The addition of Cr in the slurry compositions was found to successfully decrease the activity of Al

I. Slurry aluminizing of polycrystalline Ni20Cr

Nichrome (nickel with 20 wt.% Cr) was one of the first nickel-based alloys used for high-temperature applications and served as a basis for the development of more complex alloys [1][2][3][4][5]. The addition of Cr considerably increased the high-temperature corrosion resistance of nickel by forming through oxidation a protective chromia scale on the surface [2,4,6,7]. Although Ni-Cr alloys are not present in gas turbine engines anymore, they are still used for structural applications in the industry.

Polycrystalline Ni20Cr (with 1.5 wt.% Si) was employed in this study as a model alloy to investigate the influence of Cr in the parent material upon aluminizing [8][9][10]. This study aimed at elucidating the coating formation and the beneficial influence of Cr addition in slurry compositions upon aluminizing of a Ni-Cr alloy before extrapolation to nickel-based superalloys.

I.1. From Al microparticles

As previously described, aluminizing of Cr-free and Cr-poor substrates from slurries containing Al microparticles occurs very quickly by wetting, dissolution of the parent material (e.g. Ni or Fe) and the resulting self-propagating synthesis and final state diffusion [11]. In contrast, the presence of Cr into the Ni20Cr (1.5 wt.% Si) substrate was found to considerably affect the mechanisms of formation of the diffusion zone from slurries containing Al microparticles [12]. After a complete heat treatment in flowing Ar (400°C-1h + 700°C-2h + 1100°C-2h), B. Rannou observed the partial detachment of the top coat and the formation of random diffusion islands from both dissolution of the substrate into the molten Al and solid-state diffusion of Al [12]. However, other random areas were completely free from diffusion. The island-like diffusion morphology was associated with the segregation of chromium at the interface between the aluminized layer and the substrate [12]. Cr has indeed a limited solubility in Al-rich NixAly phases [13][14][15][16] and is also known to block Al diffusion [17]. In the case of Ni20Cr, the chromium aluminides formed upon aluminizing process even at low temperature (625°C) were responsible for such blockage [10]. Such Al-Cr intermetallic phases are a consequence of the fast inward diffusion of Al into the substrate promoting the formation of NiAl3 and Ni2Al3 phases [11,18].

Since the solubility of Cr is low in these phases (i.e. about 2 at.% in Ni2Al3) [16,19], Cr is expelled out from the Ni2Al3 matrix and combines with further Al ingress to form Al8Cr5 grains [8,9,10,15]. The formation of α-(Cr) and Al8Cr5 grains at the interface between the aluminized layer and the substrate acted as a diffusion barrier for Al ingress. An additional consequence of the blockage is the appearance of Kirkendall porosity at the interface between the additive and the interdiffusion layers that arises from the major outward diffusion of Ni that is not compensated by the Al ingress as it is trapped by Cr.

In this study, Ni20Cr samples were aluminized from Al microparticles in flowing Ar with the same conditions than those of B. Rannou [12]. The optical surface micrographs and SEM cross-section images of the samples after diffusion heat treatment and complete heat treatment are given in Fig

Fig. IV-1. Optical surface micrographs and SEM cross-section images (BSE mode) for Ni20Cr aluminized with Al microparticles after (a), (c) the diffusion heat treatment (400°C-1h + 700°C-2h) and (b), (d) the complete heat treatment (400°C-1h + 700°C-2h + 1100°C-2h) in flowing Ar.

In the work of Le Guevel et al., polycrystalline Ni20Cr (1.5 wt.% Si) was aluminized by out-of-pack process (APVS, Snecma process) at 1080°C for 6 hours [20]. This low-activity high-temperature (LAHT) process fostered the formation of an outward grown coating based on the β-NiAl phase [21,22]. The coating thickness was even with an outer β-NiAl matrix free of Cr-containing precipitates [20,23]. Chromium was mainly located at the inner interdiffusion zone in the form of elongated precipitates [10]. Brossard et al. elaborated thin aluminide coatings on Ni20Cr (1.5 wt.% Si) by first depositing an Al film by using DC cathodic magnetron sputtering technique [7]. The initial deposition of approximately 3 µm Al films resulted in the formation of a Ni2Al3 matrix after annealing at 600°C for 1 hour. Further annealing at 900°C brought about the stabilization of β-NiAl whereas both β-NiAl and γ'-Ni3Al phases were observed after annealing at 1100°C [7]. Tu et al. successfully aluminized several arc-melted Ni-Cr alloys by pack cementation process [8]. The Ni-Cr alloys containing between 5 at.% Cr and 35 at.% Cr were aluminized between 800 and 1150°C using AlF3 activator. The pack composition was adjusted to modify the Al activity upon aluminizing using pure Al microparticles or pre-alloyed Al and Ni powders [8]. Depending on the pack composition and the aluminizing treatment (dwell time and temperature), different microstructures were obtained for the coatings. For these processes, the coatings were homogeneous in thickness and did not exhibit island-like diffusion morphology as observed by Rannou [12] and in the present study (Fig. IV-2).

Fig. IV-2. BSE cross-section images of Ni20Cr substrate slurry coated with Al after (a) the diffusion heat treatment (400°C-1h + 700°C-2h) and (b) the complete heat treatment (400°C-1h + 700°C-2h + 1100°C-2h

).

-114-Out-of-pack and pack cementation processes involved the deposition of Al on the substrate from a gaseous phase [24]. Even though a solid Al film was deposited on Ni20Cr in the work of Brossard et al. [7], the diffusion heat treatment was realized at a temperature lower (600°C) than the melting temperature of Al (i.e. 660°C [25]). Therefore, the coating formation involved only solid-state diffusion.

Unlike these processes, the coating formation on a nickel-based material slurry coated with Al microparticles involved self-propagating high-temperature synthesis reactions [11,18]. These high exothermal reactions typical of high Al activity slurry aluminizing, were attributed to provoke the island-like diffusion morphology observed by B. Rannou [12] and in this study. The formation of liquid phases upon high Al activity aluminizing led to the dissolution of the Ni20Cr substrate and fostered the precipitation of Cr at the reaction front between the aluminized layer and the substrate that impeded homogeneous diffusion. This brought us about the idea of limiting the exothermal reactions occurring between the Al microparticles and the Ni20Cr substrate by adding Cr to the slurry composition [26].

I.2. From Al-Cr mixed slurries

The aluminizing of pure nickel substrate with Al-Cr mixed slurries using a typical low-activity heat treatment employed in the repair industry (annealing at 1080°C for 6 hours [20,27]) fostered the formation of the β-NiAl phase by interdiffusion between nickel and AlxCry intermetallic compounds (cf. paragraph IV.2. in Chapter III). Therefore, the same heat treatment was carried out to elaborate the Al-Cr mixed slurry coatings on Ni20Cr substrate. The cross-section micrographs of the aluminized layers for the three Al-Cr mixtures (67Al33Cr, 44Al56Cr and 20Al80Cr in wt.%) are shown in Fig.

IV-3. For the three compositions, the reaction front between the aluminized layers and the underlying substrate was quite even. These observations suggest that the addition of Cr to the Al microparticles was highly beneficial for the slurry aluminizing of Ni20Cr substrate. As observed with pure nickel, the higher the Al concentration in the initial mixture, the thicker the diffusion layer after annealing. For the three compositions, the top coat was composed of intermetallic phases and oxides and remained adherent on the surface (Fig. IV-3).

The diffusion layers were based on the β-NiAl phase for the three compositions (Fig. IV-4) and the thickness of the continuous aluminized layer reached about 50µm, 20µm and 10µm for respectively; the 67Al33Cr, 44Al56Cr and 20Al80Cr slurry mixtures. Some voids were present at the interface between the aluminized layer and the underlying substrate (Fig. IV-4a andb). These voids resulted from counter-current vacancy flow to the motion of Ni, Cr and Al [28] and strongly suggested that the

Fig. IV-3. Optical cross-section micrographs (low magnification) of Ni20Cr substrate slurry coated with Al-Cr mixtures after annealing in flowing Ar (400°C-3h + 1080°C-6h

).

-116-outward diffusion of Ni was faster than the inward diffusion of Al upon aluminizing. A lateral decohesion of the coating layers seemed to occur by coalescence of these voids when using the richest Cr composition (20Al80Cr) (Fig. IV-4c). Some Cr-rich precipitates (identified as α-Cr) were found in the β-NiAl matrix because of the low solubility of Cr in β-NiAl [9,13].

For a better comprehension of the mechanisms of formation of the coatings, X-ray maps were Al8Cr5 (respectively about 1 at.% and 3 at.% [16]). In addition, since the solubility of Cr in the β-NiAl is lower than 8 at.% at 1080°C [2,13], undissolved Cr was trapped in the β-NiAl matrix. According to the X-ray maps (Fig. IV-5), nickel diffused up to 55µm and 45µm from the substrate surface, in the 44Al56Cr and in the 20Al80Cr deposits, respectively. For both systems, the outermost region of the slurry deposit did not react with nickel and was still composed of AlxCry phases together with α-(Cr) solid solution (Fig. IV-6). Oxygen was found to be mainly associated with Cr and Al, in relation with the higher affinity of these two elements for oxygen than that of nickel.

Fig. IV-4. SEM cross-section images of Ni20Cr substrate slurry coated with different Al-Cr mixtures after annealing in flowing Ar (400°C-3h + 1080°C-6h).

Fig. IV-5. SEM cross-section images and corresponding elemental distribution by X-ray mapping of Ni20Cr slurry coated with (a) 44Al56Cr and (b) 20Al80Cr composition.

-118-Therefore, the addition of Cr in the slurry composition was found to successfully decrease the Al activity on the substrate surface upon aluminizing and to promote the growth of the coating by predominant outward diffusion of nickel [8]. The highly exothermic reactions occurring between the Al microparticles and the nickel-based substrate were limited by the formation of AlxCry in the slurry deposit at low temperature (around 640°C) [26]. However, the subsequent increase of temperature at 1080°C for 6 hours fostered the direct formation of the -NiAl phase and resulted in homogeneous diffusion layers. For the 44Al56Cr and 20Al80Cr mixtures, the Al content of the slurry deposit was too low to grow a diffusion layer as obtained by conventional low-activity CVD process [20,23]. The threshold to obtain suitable low-activity coatings on Ni20Cr substrate was therefore ranging from 67Al33Cr to 44Al56Cr composition (mixtures between 80 at.% Al and 60 at.% Al). These Al concentrations were higher than the threshold considered with pack cementation process where the outward diffusion of nickel predominates for packs with less than 50 at.% Al [8]. This difference might be explained by the absence of gas phases in the study with respect of the CVD reactions ruling pack cementation, out-of-pack and CVD itself [29]. analyses of the sample were not carried out to identify the crystal structure(s) of these phases, the EDS spot measurements in the cross-section revealed a composition (27Al-1Si-64Cr-7Ni in at.%) alike to a three-phase range (Cr)-β-γ2 when considering a temperature of 1080°C [16]. The γ2 phase is the low-temperature phase of the Al8Cr5 compound [16,30]. Since a slow cooling ramp was used for the heat treatment (5°C/min), the congruent transformation (Cr) → AlCr2 occurred at about 910°C upon cooling [31]. Indeed, Huang et al. found that AlCr2 can precipitate in both α-(Cr) and β-NiAl phases upon cooling [10,15].

Fig. IV-7. (a) SEM cross-section image and (b) EDS concentration profile for the 45Cr-55Al double-layer system elaborated on Ni20Cr after annealing in Ar (400°C-3h + 1080°C-6h

).

-120-Elemental X-ray maps of the cross-section were performed to depict the elemental distribution of nickel, chromium, aluminium and oxygen (Fig. IV-8). The interdiffusion zone (IDZ) was easily identified by both the depletion of nickel and the enrichment in chromium. Some Al-rich spots were observed just above the IDZ (Fig. IV-8b). These areas did not coincide with oxygen nor silicon distribution and therefore could not be attributed to the formation of alumina or AlxSiy phases. Probably, such Al-rich precipitates can be associated with carbides but the whole surface is contaminated with C and we cannot confirm this hypothesis. If a small quantity of Cr was found to be dissolved in the β-NiAl coating matrix, fine Cr-rich precipitates were detected right above the IDZ (Fig. IV-8c). Similar precipitates were also observed with a conventional low-activity process on Ni20Cr [23]. Some Cr was found to be trapped in the outer Alrich β-NiAl phase. The large particles observed suggested that these Cr enrichments were not caused by precipitation [10,13,15] -122-

The Ni20Cr substrate with the Cr-Al double-layer system was also submitted to the complete heat treatment (400°C-1h + 700°C-2h + 1100°C-2h) for comparison with the simple aluminide coating (Fig.

IV Chapter II, the directionally solidified CM-247 LC superalloy displays a duplex γ-γ' microstructure [34][35][36]. Aluminizing of CM-247 LC was realized on as-received samples without either solutionizing or ageing treatments, which are required for mechanical testing like creep [36] or tensile tests [34,35] but not for coating manufacture [37][38][39]. Some carbides were originally present in the interdendritic regions of the superalloy (cf. Fig. II-2a). These inert carbides were particularly helpful to investigate the mechanisms of formation of the diffusion coatings as they acted as natural markers to determine the initial substrate surface location before aluminizing [40,41]. The aluminizing of CM-247 LC from high Al activity slurries has already been reported in Refs. [12,11,39,42].

The three Al-Cr mixed slurries selected in this study (67Al33Cr, 44Al56Cr and 20Al80Cr in wt.%)

were deposited on the CM-247 LC superalloy. The resulting coatings produced after aluminizing treatment in flowing Ar (400°C-3h + 1080°C-6h) are presented in Fig. IV-11.

Fig. IV-11. Optical cross-section micrographs (low magnification) of CM-247 LC superalloy slurry coated with Al-Cr mixtures after annealing in flowing Ar (400°C-3h + 1080°C-6h

).

-124- For the 44Al56Cr slurry mixture, the tetragonal AlCr2 phase [15,31] was clearly identified together with low reflections of α-Al2O3 structure. This phase was probably formed during the slow cooling ramp (5°C/min) following the (Cr) → AlCr2 invariant reaction around 910°C [31]. For the 20Al80Cr one, three broad peaks were observed at 2θ positions typical of a cubic structure and the EDS analyses revealed average compositions of 24Al-75Cr-1Ni in at.%. This suggests that the cubic structure probably refers to (Cr) containing dissolved Al and Ni, which is consistent with different reports whereby Cr dissolves about 25 at.% of Al between 0 and 1000°C [31,43] and even up to 43 at.% at 1000°C [16,43]. The incorporation of Al and Ni in the Cr lattice modified its parameters which could ). The IDZ resulted from the precipitation of substrate elements having low solubility in the β-NiAl phase (Cr, Ta and W) [40,41]. Some Ta and Wrich precipitates were also found in the outer region of the coating as observed in high Al activity slurry aluminizing [39]. This suggests that the formation of the diffusion zone with the 67Al33Cr slurry

Fig. IV-13. XRD patterns of the surfaces for the 67Al33Cr, 44Al56Cr and 20Al80Cr slurry compositions elaborated on CM-247 LC superalloy after annealing in flowing Ar (400°C-3h + 1080°C-6h).

mixture involved both inward diffusion of Al at low temperature and outward diffusion of Ni at high temperature. It therefore appeared that the Al content of the 67Al33Cr slurry mixture was too high to produce a coating with a low-activity microstructure.

For the samples aluminized with the 44Al56Cr slurry mixture (Fig. IV-14c), two layers were identified:

an outer additive layer with a rough surface and an inner IDZ with a thickness of approximately 20µm.

The additive layer was composed of the Ni-rich β-NiAl phase with an Al concentration ranging from 36 to 40 at.% (Fig. IV-14d). This phase was found to dissolve more Cr than the one obtained with the 67Al33Cr slurry mixture (Fig. IV-14b). The outer region of the additive layer was free of precipitates and carbides from the substrate suggesting that outward diffusion of nickel predominantly occurred upon aluminizing [1,40,42]. However, some fine Ta and W-rich precipitates were randomly distributed across the additive layer. These precipitates probably resulted from the initial reaction between Al microparticles and the underlying substrate at low temperature (at about 640°C [11]). This reaction was presumably concomitant with the exothermic reaction occurring at 637±2°C between Al and Cr microparticles in the slurry deposit [26].

Fig. IV-14. BSE cross-section images and EDS concentration profiles for CM-247 LC superalloy slurry coated with (a), (b) 67Al33Cr, (c), (d) 44Al56Cr and (e), (f) 20Al80Cr mixtures after annealing in flowing Ar (400°C-3h + 1080°C-6h

).

-128-

The elemental distributions of Ni, Al, Cr and Ta for the three Al-Cr mixed slurries were also determined after annealing (Fig. . Note that the X-ray maps of W were generally very close to the one of Ta and, for the sake of clarity, only X-ray maps of Ta are presented in the manuscript. Ta and W were found to preferentially precipitate in the IDZ for the three slurry mixtures. For the 44Al56Cr composition, fine Ta and W precipitates were aligned at the interface between the additive layer and the IDZ, which is typical of outwardly grown coatings [40].

Fig. IV-15. BSE cross-section images and X-ray maps of Ni, Al, Cr and Ta for the (a) 67Al33Cr, (b) 44Al56Cr and (c) 20Al80Cr mixed slurries elaborated on CM-247 LC superalloy after annealing in Ar (400°C-3h + 1080°C-6h).

Since the aluminizing treatment must not modify the microstructure of the parent material to maintain the original mechanical properties, an aluminized sample was electrolytically etched with oxalic acid (Fig. . This metallographic attack preferentially etches the γ-(Ni) phase, which allowed to reveal the γ-γ' microstructure of the superalloy underneath the diffusion zone (Fig. IV-16a). Note also that the microstructure of the coating was similar to the one obtained with high Al activity slurry aluminizing [12,39,40,44,47]. This observation confirmed that the Al content of the 67Al33Cr composition was too high to foster the direct formation of β-NiAl. The greater magnification image of the interface between the IDZ and the substrate (Fig. IV-16b) demonstrated that the γ-γ' microstructure of the substrate was maintained after annealing at 1080°C for 6 hours.

As observed with the model materials (pure nickel and Ni20Cr substrate), Al and Cr microparticles reacted in the slurry deposits and formed AlxCry intermetallic compounds. The initial Al content of the Al-Cr slurry mixture was of utmost importance to control the Al activity upon aluminizing [42]. Unlike pack cementation [37,41,48] or out-of-pack processes [21,22], the synthesized Al-Cr masteralloys were in direct contact with the substrate to aluminize with no need of halide activators [42,26]. The stoichiometry of the synthesized AlxCry phases was found to considerably modify the mechanisms of formation of the diffusion coatings. This particularly modified the motion of Ni and of Al upon aluminizing and hence the preferential growth of the coating by inward diffusion of aluminium at low temperature or outward diffusion of nickel at high temperature.

Fig. IV-16. (a) BSE cross-section image of the 67Al33Cr slurry mixture elaborated on CM-247 LC superalloy after annealing in flowing Ar (400°C-3h + 1080°C-6h) and (b) details of the interface between the IDZ and the underlying substrate (10 vol.% H2C2O4 electrolytically etch).

-

Fig. IV-17. (a) Macrograph and (b) surface micrograph of the CM-247 LC sample slurry coated with the 56Cr-44Al double-layer system after annealing in flowing Ar (400°C-3h + 1080°C-6h

).

-132-

The microstructure of the coating was completely different than the one obtained from high Al activity aluminizing [12,39,41,44,47]. The coating structure consisted of three main layers: an interdiffusion

Fig. IV-19. BSE cross-section image and corresponding X-ray maps of (a) Ni, (b) Al, (c) Cr and (d) Ta for the 56Cr-44Al double-layer system elaborated on CM-247 LC after annealing in flowing Ar (400°C-3h + 1080°C-6h

).

-134-

The concentration profile of the coating is given in Fig The low Ti content of CM-247 LC (0.7 at.%) led to a very low concentration gradient with the outward growing -NiAl phase and could explain its low concentration within the additive layer after only 6 hours diffusion at 1080°C [49]. However, one can expect that for longer annealing time (e.g. upon oxidation), the substrate elements with high diffusivity in the -NiAl phase (mainly Ti, Ta and Hf) will diffuse towards the surface [50,52,53].

Fig. IV-20. EDS concentration profile of the 56Cr-44Al double-layer system elaborated on CM-247 LC superalloy after annealing in flowing Ar (400°C-3h + 1080°C-6h).

-

Fig. IV-22. Cross-section optical micrograph of the 56Cr-44Al double layer system elaborated on CM-247 LC superalloy after annealing in flowing Ar (400°C-1h + 700°C-2h + 1100°C-2h).

Fig. IV-23. BSE cross-section image and corresponding EDS maps of (a) Ni, (b) Al, (c) Cr and (d) Ta for the 56Cr-44Al double-layer system elaborated on CM-247 LC after annealing in flowing Ar (400°C-2h + 700°C-2h + 1100°C-2h).

precipitation was confined in the IDZ (same observations for W). The substrate carbides (identified as the bright coarse phases in the BSE image of Fig. were enriched in Hf, Ta and W.

The mean composition of the Cr-rich phases present above the additive layer was approximately 37Al-56Cr-1Co-6Ni in at.%. By comparison with the results obtained with the low-activity heat treatment (i.e. 26Al-66Cr-1Co-7Ni in at.%), it appeared that the Al concentration of these phases was higher after the two-step heat treatment. This could be explained by the longer dwell time at high temperature for the low-activity heat treatment (i.e. 1080°C-6h) than for the two-step heat treatment (i.e. 1100°C-2h). So, the longer the diffusion time, the greater the outward diffusion of nickel to react with the Al-Cr phases of the slurry deposit. Consequently, the amount of Al dissolved from the Al-Cr phases to foster the formation of -NiAl was greater after 6 hours at 1080°C than after only 2 hours at 1100°C.

The longer dwell time at high temperature was also observed to promote the dissolution of the Crrich phases enclosed within the -NiAl coating matrix. This also implied their segregation in the outer region of the coating. All the different observations confirmed that the outward diffusion of nickel was controlling the growth of the coating.

EDS spot measurements in the cross-section of the coating were performed to plot the concentration profile (Fig. . The analyses confirmed that the additive layer was mainly composed of the Al-rich -NiAl phase with local Cr enrichments. This phase extended up to the interface IDZ-additive layer where the concentration of Al sharply decreased. The IDZ was mainly composed of the Ni-rich β-NiAl phase where the substrate elements having low solubility in this phase (Cr, Ta and W) precipitated during the aluminizing process [40,41].

Fig. IV-24. EDS concentration profile of the 56Cr-44Al double-layer system elaborated on CM-247 LC superalloy after annealing in flowing Ar (400°C-1h + 700°C-2h + 1100°C-2h

).

-138-

c) Discussion

 Growth of the coating:

The coating microstructure observed on

 Mechanisms of formation:

The significant Cr enrichments in the additive layers of the slurry coatings suggested that these enrichments were associated with undissolved Cr-rich phases rather than with precipitation in the β-NiAl phase, as also observed by Gleeson et al. [32]. Considering the original system of the present study, AlxCry intermetallic compounds were synthesized in the slurry deposit at low temperature (at about 640°C). The Al activity was thus greatly reduced on the substrate surface, considerably limiting the Al inward diffusion at low temperature (i.e. 650°C) [26,42]. At high temperature (i.e. 1080°C or 1100°C), nickel outwardly diffused and reacted with the AlxCry intermetallic phases synthesized in the slurry deposit (cf. After annealing, undissolved α-(Cr) phases remained in the additive layer suggesting that the outward diffusion of nickel was significantly higher than the inward diffusion of Al at all time during the aluminizing process [32,33,55].

 Influence of the cooling ramp on the formation of Al-rich -NiAl phase:

For the sample annealed at 1100°C for 2 hours (Fig. , the formation of the Al-rich β-NiAl layer should not possible at the diffusion temperature (1100°C) because of the major outward diffusion of nickel. The appearance of this Al-rich β-NiAl zone in the additive layer can be explained by the effect of the cooling ramp on the final microstructure of the coating. For example, a layer of Al-rich β-NiAl was highlighted by EPMA on low-activity coatings produced by out-of-pack CVD process on different nickel-based superalloys [44]. For this process, halide activators are used to attack Cr-Al masteralloys and form a gaseous phase to allow aluminium transport towards the sample surface [22]. During the vapour phase process, the coating box is heated to the aluminizing temperature (typically in the range 1050°C-1150°C) [22,59]. After adsorption of the gaseous molecules on the substrate, solid diffusion of deposited Al in the substrate occurs. At the end of the diffusion dwell time, the coating box is cooled down to room temperature. Because of the inertia of the retort, the temperature profile of the cooling ramp is usually divided in two sections: a first section with a continuous decrease of the temperature down to an intermediate temperature (in the range 800°C-900°C) and a second section with a slow cooling ramp down to room temperature [59]. At temperatures below 900°C, the diffusivity of Al is more significant than that of Ni in β-NiAl phase [46]. Because gaseous aluminium phases are still present in the coating box upon cooling, deposition of Al on the sample surface is maintained. Therefore, Al diffuses into the β-NiAl additive layer, which converts into Al-rich β-NiAl phase thanks to the continuous supply of Al. In Al-rich β-NiAl, aluminium diffuses ten times faster than nickel and this can explain that this layer could reach 10 to 15µm in the outer region of the coating after the aluminizing process [44].

The cooling profile of the horizontal furnace used for the aluminizing tests (cf. paragraph II.4. in Chapter II) was experimentally determined and found to be close to the one given by Wang et al. for a conventional low-activity CVD process [59]. Therefore, it is believed that the Al-rich β-NiAl layer observed in the coating microstructure (Fig. IV-22) also formed upon cooling in this work. Unlike CVD techniques, the Al donors were in direct contact with the underlying substrate here [26] thus giving a continuous Al supply upon cooling until its diffusion becomes negligible at low temperature.

 Influence of the heat treatment parameters on the formation of Al-rich -NiAl phase:

Since the Al-rich -NiAl phase was not formed after the conventional low-activity heat treatment (Fig.

IV-18), the heat treatment parameters were also found to influence its formation. The initial reaction between Al and Cr microparticles to form the intermetallic compounds, occurring at approximately 640°C, was found to be very fast [26]. After further annealing for 5 hours at 650°C, little transformation occurred within the slurry deposits of Al-Cr mixed slurries (cf. in the two-step heat treatment had barely any effect on the diffusion mechanisms between the slurry deposit and the CM-247 LC superalloy. However, at higher temperatures (i.e. 1080°C-1100°C), the predominant outward diffusion of nickel dissolved the AlxCry intermetallic phases and progressively formed the -NiAl coating layer following reaction (Eq. IV-1). The main difference between the two heat treatments was therefore the annealing time (6 hours at 1080°C or 2 hours at 1100°C).

Increasing the annealing time should lead to a greater Al consumption from the AlxCry phases because of the continuous nickel outward diffusion at such high temperatures. Assuming an initial slurry deposit with the same Cr-Al ratio before heat treatment (56Cr-44Al in wt.% in this case), the remaining Al content in the slurry deposit would be higher after two hours of annealing than after six hours (progressive supply of Al to form the β-NiAl phase). Consequently, this implies that the driving force of reaction (Eq. IV-1) (or even (Eq. IV-2)) is much higher after the 1100°C-2h annealing than the 1080°C-6h one. Upon cooling, the driving force (i.e. the Al activity of the Al donors) was high enough after the two-step heat treatment to foster the inward diffusion of Al into the Ni-rich -NiAl layer. This progressively saturated the Ni-rich -NiAl layer and led to the formation of Al-rich -NiAl instead.

After the complete heat treatment, a homogeneous Al-rich -NiAl layer can therefore be observed in the additive layer of the coating (Fig. as previously reported in out-of-pack low-activity coatings [44].

 Comparison with the Al-Cr mixed slurries:

The 44Al56Cr and 20Al80Cr mixed slurries were found to greatly reduce the Al activity on substrate surface and to foster the outward diffusion of nickel at high temperature [26]. It is likely to consider that the reactions (Eq. IV-1) and (Eq. IV-2) were controlling the coating growth as proposed with the 56Cr-44Al double-layer system. After 5 hours annealing at 650°C in Ar (Fig On the other hand, the fine precipitation of refractory metals (i.e. Ta and W) in the outer region of the coating for the 67Al33Cr mixed slurry (Fig. IV-16a) confirmed that the mechanisms of formation were different. This implies that the reactions (Eq. IV-1) and (Eq. IV-2) were not controlling the coating growth when performing the slurry aluminizing with the 67Al33Cr. For a better comprehension of such discrepancies, CM-247 LC samples were slurry coated with a 33Cr-67Al double-layer system and annealed in the same conditions than the 56Cr-44Al double-layer system.

II.3. High-activity high-temperature (HAHT) coatings from the Cr-Al double-layer system

The results of paragraph II.1 demonstrated the importance of the Al/Cr ratio of the slurry deposit on the microstructure of the diffusion coating. The initial Al/Cr ratio has a significant influence on the stoichiometry of the AlxCry phases synthesized upon annealing on the substrate surface, hence, on the Al activity [26,48]. The sample slurry coated with the 67Al33Cr composition (in wt.%) was found to present a final microstructure close to that of high Al activity coatings. Therefore, this composition was selected to produce a Cr-Al double-layer system on CM-247 LC superalloy in order to compare the results to the ones obtained with the low-activity high-temperature coatings (LAHT) of paragraph II.2. The different samples presented in paragraph II.3 were slurry coated with 6±1 mg.cm -2 Cr and 12±1 mg.cm -2 Al to maintain a Cr/Al ratio close to 33Cr/67Al (in wt.%).

a) Low-activity conventional heat treatment (400°C-3h + 1080°C-6h)

The CM-247 LC superalloy slurry coated with the 33Cr-67Al double-layer system was first subjected to the low-activity conventional heat treatment. A macrograph of the annealed sample is given on Fig.

Fig. IV-25. (a) Macrograph and (b) surface micrograph of the top coat structure observed for the CM-247 LC sample slurry coated with the 33Cr-67Al double-layer system after annealing in flowing Ar (400°C-3h + 1080°C-6h

).

-143-

The coating microstructure was however significantly different than the one obtained in Fig. IV-18a with the 56Cr-44Al double-layer system. The aluminized layer was approximately 80µm thick, which is almost 30µm thicker than the outwardly grown coatings presented in paragraph II.2. In addition, the substrate carbides were now observed at the extreme surface of the diffusion zone, which indicated that the initial surface of the substrate before aluminizing remained at the interface between the diffusion zone and the slurry deposit [41]. The microstructure of the coating was very close to the one obtained by Das et al. after a single-step high Al activity pack aluminizing of CM-247 LC superalloy [41].

A schematic representation of the coating formation is given on

Fig. IV-26. (a) Cross-section optical micrograph of the 33Cr-67Al double-layer system elaborated on CM-247 LC superalloy after the conventional heat treatment in flowing Ar and (b) schematic representation of the coating formation proposed by Das et al. [41].

The IDZ/D.L. ratio of the diffusion coating was approximately 0.35, which is 3 times lower than the one obtained on outwardly grown coatings (cf. paragraph II.2). The SEM was carried out to image in BSE mode the cross-section of the 33Cr-67Al double-layer coating and get the X-ray maps (Fig Note that a finer dispersion of bright spots was also observed in the coating matrix, attributed to Ta and W-rich precipitates. These observations confirmed that the initial growth of the coating occurred by inward diffusion of Al into the CM-247 LC superalloy.

Fig. IV-27. BSE cross-section image and corresponding EDS maps of (a) Ni, (b) Al, (c) Cr and (d) Ta for the 33Cr-67Al double-layer system elaborated on CM-247 LC after annealing in flowing Ar (400°C-3h + 1080°C-6h).

-145-

b) Complete heat treatment (400°C-1h + 700°C-2h + 1100°C-2h)

A sample slurry coated with the 33Cr-67Al double-layer system was submitted to the complete heat

Fig. IV-29. (a) Macrograph and (b) surface micrograph of the top coat structure observed for the CM-247 LC sample slurry coated with the 33Cr-67Al double-layer system after annealing in flowing Ar (400°C-1h + 700°C-2h + 1100°C-2h).

Fig. IV-28. (a) EDS concentration profile of the 33Cr-67Al double-layer system elaborated on CM-247 LC after annealing in flowing Ar (400°C-3h + 1080°C-6h) and (b) detailed view of the outer region of the diffusion layer.

-

Fig. IV-30. Cross-section optical micrograph of the 33Cr-67Al double-layer system elaborated on CM-247 LC superalloy after the complete heat treatment in flowing

Ar.

-147-

The EDS concentration profile of the coating layer is given in Fig. IV-31a. The high Al concentration of the diffusion layer confirmed that the Ni-rich -NiAl phase was not directly formed during the aluminizing process and that Al-rich -NiAl was formed instead. Since the Ni2Al3 phase has a very small domain of stability at temperatures greater than 1000°C [16], the formation of Al-rich -NiAl was considered instead [41]. The W concentration in the diffusion layer was much higher than the one measured for the 56Cr-44Al composition after the complete heat treatment (Fig. . This was a consequence of the fast inward diffusion of Al in the superalloy causing the slow-diffusing elements to stay in the outer region of the diffusion layer. The BSE cross-section image highlighted the substrate carbides trapped within the inwardly grown coating (Fig. IV-31b). The formation of a thin IDZ suggested that nickel outward diffusion also occurred during the aluminizing process but to a lesser extent than the inward diffusion of Al. The mean composition of the AlxCry phases was determined by EDS (44Al-49Cr-1Co-6Ni in at.%) and suggested the major formation of the Al8Cr5 phase [14,19].

Fig. IV-31. EDS concentration profile of the 33Cr-67Al double-layer system elaborated on CM-247 LC after annealing in flowing Ar (400°C-1h + 700°C-2h + 1100°C-2h

).

-148-

c) Discussion

The Cr/Al ratio of the Cr-Al double-layer system was found to significantly influence the Such differences were highlighted with the position of the substrate carbides in the coating allowing to determine the initial surface of the substrate before aluminizing (cf. Annex 1).

 Formation of "low-activity" coatings from the Cr-Al double-layer system:

For the 56Cr-44Al composition ( Ni + AlxCry → -NiAl + α-(Cr)

(Eq. IV-3)

Fig. IV-32. Cross-section micrographs of the two Cr-Al double-layer system elaborated on CM-247 LC superalloy after (a), (b) conventional low-activity (400°C-3h + 1080°C-6h) and (c), (d) two-step heat treatment (400°C-1h + 700°C-2h + 1100°C-2h

).

-149-

The predominant outward diffusion of nickel dissolved the AlxCry phases synthesized in the slurry deposit because of its limited solubility in Al-Cr phases (about 1 at.% in Al4Cr and 3 at.% in Al8Cr5) [16]. Since the aluminizing treatment was done between 1080°C and 1100°C, the formation of Ni2Al3 phase was avoided and direct formation of β-NiAl occurred [40,41]. Since Cr has a low solubility in the β-NiAl phase, a large fraction of Cr remained undissolved as α-(Cr) phases through the β-NiAl coating matrix even after 6 hours at 1080°C (Fig. IV-32a).

The continuous outward diffusion of nickel upon aluminizing progressively consumed the AlxCry phases in Al resulting in the formation of Cr-rich phases within the slurry deposit. Since the Cr solid solution dissolves up to 43 at.% of Al [16,43], the transformation of the AlxCry phases into Cr(Al) considerably decrease the driving force of reaction (Eq. IV-3) and the new simplified reaction can be written:

Ni + α-Cr Al → -NiAl + α-(Cr)

(Eq. IV-4)

Considering the Ni-Al-Cr ternary equilibrium [19] and the observations of Huang et al. [10], only β-NiAl and α-(Cr) phases are stable at the aluminizing temperature. Therefore, as soon as α-Cr(Al)

constitutes the main fraction of the slurry deposit (i.e. the Al donor), the kinetics of aluminizing decreased because of the very low driving force of reaction (Eq. IV-4). This would explain why the aluminized layer of the sample heat treated with the 20Al80Cr mixed composition was so thin (Fig.

IV-14e).

As discussed in paragraph II.2, the formation of the Al-rich β-NiAl phase for the sample presented in Note that the coating formed on the CM-247 LC superalloy was composed of an additional layer made of hollow alumina spheres [39]. The overall Al concentration of the diffusion layer remained below 55 at.%. and was not constant in the diffusion layer. The Cr concentration of the outer region of the diffusion layer was higher than that of high activity slurry aluminized CM-247 LC [11,12,44], which indicates that some Cr diffused in the diffusion layer. A deeper characterization of the outer region of the coating highlighted the major composition of both Cr-rich phases and Ni-rich phases (Fig. IV-35a).

The distribution of fine bright precipitates corresponded to substrate elements with low solubility in the β-NiAl phase (mainly Ta and W). On the BSE image of Fig. IV-35a, the coating matrix with a dark contrast refers to the Ni-enriched phase whereas the phases with an intermediate contrast were Crenriched. The exact composition of the phases could not be determined because of the large spot resolution of the EDS but, according to literature, were presumably Al-rich β-NiAl and Al8Cr5 phases [8,10]. This two-phase stability is in good agreement with Ni-Al-Cr ternary diagram at 1000°C [14,16].

These observations confirmed that the Al inward diffusion was significantly faster than the nickel outward diffusion for the 33Cr-67Al double-layer system.

Fig. IV-33. BSE cross-section image of CM-247 LC slurry coated with the 33Cr-67Al doublelayer system after annealing in flowing Ar (400°C-3h + 1000°C-6h

).

-152-EDS spot measurements were performed in the AlxCry phases present above the diffusion layer (Fig.

IV-35a

). Two different phases were observed in the BSE images at higher magnification (not shown)

and gave average composition of 61Al-34Cr-1Co-3Ni (at.%) and 45Al-50Cr-1Co-4Ni (at.%). By

(a) BSE cross-section image and (b) EDS concentration profile of CM-247 LC slurry coated with the 33Cr-67Al double-layer system after annealing in flowing Ar (400°C-3h + 1000°C-6h).

Fig. IV-35. (a) Details of the outer region of the diffusion layer in BSE mode and (b) surface XRD pattern (full system presented on Fig. IV-34b).

will be discussed in Chapter V by investigating the influence of the atmosphere on the peripheral oxidation of Al microparticles.

Considering the 33Cr-67Al composition, Al4Cr was the main synthesized compound before reaching the dwell temperature of 1000°C (cf.

 Influence of the aluminizing temperature on the outward diffusion of nickel:

The BSE cross-section images of 33Cr-67Al system presented on IV-36a) suggesting that the outward diffusion of nickel was more important than that for the sample annealed at 1000°C (Fig. IV-36b). This also confirms that the formation of such coatings first involved inward diffusion of Al and then the delayed outward diffusion of nickel [41].

Fig. IV-36. BSE cross-section images of the 33Cr-67Al double-layer system elaborated on CM-247 LC after (a) 400°C-3h + 1080°C-6h and (b) 400°C-3h + 1000°C-6h in Ar.

-154-

III. Summary and outlook

Table IV-2 summarizes the different slurry coating systems elaborated on CM-247 LC superalloy. The details of the coating microstructures obtained after annealing treatments in flowing Ar are also given.

Their potential abilities to confer hot corrosion and oxidation resistance are also indicated whereas only two systems were found to confer a potential thermal insulation with an adherent top coat composed of hollow alumina microspheres.

Table IV-2. Summary of the different slurry coating systems elaborated on CM-247 LC superalloy in flowing Ar. HT-1: 400°C-3h +1080°C-6h, HT-2: 400°C-1h + 700°C-2h + 1100°C-2h, HT-3: 400°C-3h + 1000°C-6h. HA for high-activity and LA for low-activity.

Al From these experiments, two systems were selected for the oxidation tests (Chapter VI): the lowactivity Cr-Al double-layer system (with HT-1) as a bond coat and the 33Cr-67Al double-layer system (with HT-3) as a complete thermal barrier system. The 2 hours dwell at 700°C of HT-2 presumably conferred a better adherence of the top coat on the underlying diffusion coating for both 56Cr-44Al and 33Cr-67Al systems. However, the oxide shells of the alumina hollow microspheres remained thin after annealing in flowing Ar (Fig. IV-34a) and the top coat is expected to be unable to provide the required erosion resistance. In this way, Chapter V will present the functionalization of the complete thermal barrier system. The influence of the atmosphere (Ar, air) on the aluminizing and the formation of the complete thermal barrier system will be investigated. Special attention will be given on the thickening and the sintering of the alumina shells forming the ceramic top coat.

- The amorphous alumina shell therefore crystallizes into γ-Al2O3 usually at about 550°C before the melting of Al [2,3,13,14]. Since the density of γ-Al2O3 is greater than that of the amorphous layer [6,15], the newly formed γ-Al2O3 layer is no longer continuous on the aluminium surface, which induces cracks in the alumina shell (Fig. V-1). If the 2 is high enough, healing of the defective alumina shell occurs by oxidation of the exposed metallic Al. When the temperature is greater than that of the melting of Al, molten Al can diffuse through the cracks depending on the pressure build-up in the Al core and the mechanical strength of the alumina shell [6,7,10].

Fig. V-1. Mechanisms of oxidation and of Al diffusion for micro-sized and nano-sized particles proposed by Levitas et al. [7].

Trunov et al. [2,16] and Rufino and collab. [3] studied the oxidation behaviour of Al powders with different particles size in a dynamic dry air atmosphere. These authors identified up to four domains on the oxidation curves depending on the temperature range. Detailed explanations are given in Ref. [16] with the kinetics of aluminium oxidation and the calculation of activation energies for the different identified transformations. identified by thermal analysis (from Ref. [16]). Up to four stages of oxidation were considered from these studies [2,3,16]:

-ambient-450°C (stage I): weight loss due to desorption of water and organic compounds from the surface and to the dehydroxylation of the amorphous oxide layer, followed by a slow increase attributed to the growth of the amorphous oxide layer.

-450°C-650°C (stage II): stepwise oxidation in a narrow temperature range (crystallization of the amorphous alumina into γ-Al2O3). This transformation occurred when the amorphous layer reached a critical thickness of about 4 to 5 nm [3,11,12]. The newly formed γ-Al2O3 is no longer continuous on the aluminium surface exposing metallic Al to the atmosphere [6,15].

This induces an increase of the oxidation rate at the beginning of stage II while the γ-Al2O3 is healing and eventually covers the whole aluminium surface at the end of stage II.

-650°C-950°C (stage III): second acceleration of the oxidation rate attributed to the growth of the γ-Al2O3 layer accompanied by further crystallographic transformations. These transformations probably follow the γ-Al2O3 → -Al2O3 → θ-Al2O3 → α-Al2O3 sequence as reported in other studies [14,15,17,18] although the direct transformation γ-Al2O3 → α-Al2O3

has also been reported [2,15,19]. These transformations were not expected to significantly affect the oxidation rate of γ-Al2O3 according to Trunov et al. [16].

-950°C-1500°C (stage IV): progressive transformation of transition aluminas into the stable α-Al2O3 polymorph by the end of stage III. The phase transition γ-Al2O3 → α-Al2O3 induces a volume shrinkage of about 13.8% because of their different crystal structures [6,20]. This leads to the formation of pores and nano-cracks in the alumina shell [6]. When the oxide shell is completely converted into α-Al2O3, the oxidation rate rapidly decreases. This could be attributed to the slow grain boundary diffusion within the α-Al2O3 scale upon further oxidation [16].

Chapter V

-161-It has to be noted that the temperature at which the events occur mainly depends on the size of the particles [1,3,16]. Thermal events are thus delayed to higher temperatures by increasing the particle size. Kolarik et al. investigated the oxidation of two types of micro-sized Al particles in air by in situ high-temperature X-ray diffraction [20]. The transformation of the amorphous alumina layer into γ-Al2O3 was detected from 425°C with 0.3-0.7 µm particles and from 550°C with 2-5 µm ones.

Therefore, considering that the Al particles of the present study have a mean diameter of 8 µm (cf. Fig.

II-2 in

Chapter II), this transformation is expected to occur to at 550°C or even higher.

The mechanisms of formation of the environmentally friendly slurry coatings (from Al microparticles) have already been presented in inert atmosphere (Ar) [21,22] and air [23] on various nickel-based materials. However, none of such studies investigated the influence of the atmosphere on both the aluminizing of the substrate and the consolidation of the ceramic top coat composed of hollow alumina spheres. Therefore, the present study aimed at elucidating the influence of the atmosphere during the annealing treatment on these two features. Pure nickel was used as a model material before extrapolation to nickel-based superalloys. Note that in Ref. [23], the authors applied an intermediate electrodeposited nickel layer before slurry deposition. The intermediate layer was shown to increase the adherence of the ceramic top coat. If the authors described the mechanisms of formation of the complete thermal barrier system (aluminide coating + alumina top coat) in air, the benefit of the nickel intermediate layer was not fully explained.

Fig. V-2. Mass variation of Al powder oxidizing in a thermal analyser with the identification of the different stages of oxidation. The respective changes in the growing alumina scale are presented schematically (from Ref. [16]).

Chapter V -162-Two different atmospheres were tested: Ar (inert) and synthetic air (oxidizing). Since the heating rate has a strong influence on the peripheral oxidation and the opening of the Al microparticles [1], this parameter was fixed at 5°C/min. Trunov et al. have shown that the slower the heating ramp, the thicker the aluminium oxide shell [2]. Two heat treatments were performed based on previous works done in the group of Pedraza et al. [24][25][26] for a better understanding of the mechanisms (Fig. V-3).

Fig. V-3. Details of the annealing treatments performed to investigate the influence of the atmosphere on the aluminizing of pure nickel with (a) the diffusion heat treatment and (b) the complete heat treatment.

Chapter V -163-II. Influence of the atmosphere (Ar, air) on the aluminizing of pure nickel from Al microparticles

II.1. Diffusion heat treatment (400°C-1h + 700°C-2h)

Pedraza et al. showed by using in situ E-SEM observations that the release of Al due to the shrinkage of the oxide shells developed on the microspheres was very fast [1]. For the different atmospheres investigated (He-4%H2, O2, air and H2O), shrinkage of the oxide shells usually occurred in the 650-700°C range. It was also found that water vapour thickened the oxide shells of the alumina particles more than any other atmosphere. Galetz et al. confirmed that such oxide shells were very thin in inert atmosphere (Ar) allowing Al to react almost completely with the underlying substrate at 700°C [22].

In this case, the composition of the atmosphere is thus expected to play a significant role upon heating from 400°C to 700°C and the 2 hours dwell at 700°C (Fig. V-3a). One objective was to significantly increase the oxide shell thickness and the sintering of microparticles upon annealing to consolidate the ceramic top coat in a single heat treatment, without altering the microstructure of the diffusion layer. 4c). From the SEM observations, it appeared that the larger microparticles burst preferentially whereas the smaller ones were still filled with metallic aluminium. It is known that the solid-liquid transformation of Al is accompanied with a volume expansion of approximately 12% that induces tensile stresses on the peripheral oxide shell [4]. Therefore, it is sound to believe that the smaller microparticles, hence containing less Al but a greater specific surface area, accommodate better the stresses while the Al core is in expansion [6,7,27]. This is in good agreement with data reported only the diffusion layer with the Ni2Al3 crystal structure was probed by the XRD [30]. Raman microspectrometry was carried out to characterize the hollow microspheres (results not shown), since the

Fig. V-4. SEM surface morphology of the top coat structure after the diffusion heat treatment in (a) Ar and (b) synthetic air and (c) corresponding EDS surface analyses.

sensitivity for this technique appears better than that of XRD to detect thin oxide scales [32,33] and shows a better ability to differentiate the alumina polymorphs [34,35]. The signature of the γ-Al2O3 phase could then be identified [34,35]. On several areas of the sample, the energy brought by the laser source was sufficient to transform the oxide shell into the stable α-Al2O3 polymorph. The mechanisms of formation are not yet completely understood but this observation could give an interesting way to promote the peripheral oxidation of Al microparticles and stabilize the stable α-Al2O3 polymorph.

Note that several studies investigated the laser-induced transformation of γ-Al2O3 into α-Al2O3 for plasma sprayed coatings [31,36,37].

For the sample heat treated in synthetic air, additional crystal structures were identified on the XRD phase was also identified by XRD (Fig. V-5c). The composition of the atmosphere did not seem to significantly modify the mechanisms of formation of the diffusion layer previously reported by Galetz et al. [22]. The formation of the diffusion layer is governed by SHS reactions between Ni and Al occurring in both inert and oxidizing atmospheres [22]. The reaction is considered as a combustion

Fig. V-5. Cross-section micrographs of the full thermal barrier system elaborated on pure nickel after the diffusion heat treatment in (a) Ar and (b) synthetic air and (c) corresponding XRD patterns of the surface.

Chapter V -166-synthesis with air. Since a greater amount of metallic Al remained in the top coat after heat treatment in synthetic air, the diffusion layer was thinner than that of the sample annealed in Ar.

The thermogravimetric curves obtained for the diffusion heat treatment of Al slurry coated nickel in Ar and in synthetic air are given on Fig. V-6. Note that the organic binder used to elaborate the slurry coatings may also exert an influence on the oxidation rate of the microparticles. Nevertheless, the influence of the organic binder was not investigated in the present study. On the thermogravimetric curves presented in Fig. V-6, the events occurring below 150°C were mainly attributed to the release of water and to the homogenization of the temperature and of the atmosphere inside the reaction chamber. The strong weight loss recorded after 2 hours of annealing in both atmospheres was attributed to the decomposition of the organic binder [38]. During the dwell at 400°C, a stabilization of the mass was observed for the sample heat treated in synthetic air whereas a slow weight loss was identified for the one heat treated in Ar (stage I in Fig. V-2). This difference might be explained by a faster dehydroxylation of the amorphous oxide layer in synthetic air than in Ar because of the greater 2 in the reaction chamber.

The mass variation for the two systems was completely different upon heating at 5°C/min between 400°C and 700°C (stage II in Fig. V-2). For the sample heat treated in synthetic air, a stepwise oxidation was observed and centred at approximately 615°C [3,16]. This could correspond to the fast crystallization of the amorphous alumina layer into γ-Al2O3 reported by Kolarik et al. [20]. Since the newly formed γ-Al2O3 layer was not continuous on the aluminium surface [2,16], healing of the film occurred by oxidation of exposed metallic Al. For the sample heat treated in Ar, only a small weight

Fig. V-6. Thermogravimetric curves obtained for the diffusion heat treatment of pure nickel slurry coated with Al in Ar and synthetic air.

Chapter V -167-gain was observed maybe indicating that healing of the γ-Al2O3 layer was significantly slower than in synthetic air and that the γ-Al2O3 film remained very thin. Upon annealing at 700°C for 2 hours, both systems presented similar slopes of the mass gain suggesting that the same oxide (i.e. γ-Al2O3) was growing (stage III in Fig. V-2). As a result, the oxide shells developed on the microspheres must be thicker (and more mechanically resistant) for the sample heat treated in synthetic air than the one heat treated in Ar. This would explain why the whole Al reacted with the nickel substrate in Ar whereas metallic Al was still present in the top coat structure after diffusion heat treatment in synthetic air (Fig.

V-5).

II.2. Complete heat treatment (400°C-1h + 700°C-2h + 1100°C-2h)

The ). This evolution is in good agreement with the high resolution gravimetry analyses of Al microparticles done by Rufino et al. [3] and with the results of Trunov et al. [2,16]. The increase of the oxidation rate was attributed to the growth of the γ-Al2O3 layer and further alumina crystallographic transformations from the γ phase (Fig. V-2). When the temperature reached 1100°C, the kinetics of oxidation decreased because of the exclusive growth of the stable α-Al2O3 in the top coat [20] and as a TGO scale [39]. The transformation of γ-Al2O3 in the top coat structure after complete heat treatment was confirmed by XRD (Fig. V-8c) and Raman micro-spectrometry (Fig.

V -10).

Fig. V-8. Cross-section micrographs of the full thermal barrier system elaborated on pure nickel after the complete heat treatment in (a) Ar and (b) synthetic air and (c) corresponding XRD patterns of the surface.

As described in the experimental methods section (Chapter II), several Raman spot analyses were realized on the surface of the samples. After the complete heat treatment, the measurements were proposed by Garriga-Majo et al. [39] and the in situ XRD measurements of Kolarik et al. [20]. However, these studies were carried out in air. According to Pedraza et al. [1], performing the annealing treatment in inert atmosphere must considerably reduce the peripheral oxidation of the microspheres.

The 2 inside the reaction chamber had therefore a critical effect on the oxidation rate of the microparticles and the transition of metastable alumina polymorphs into the stable α-Al2O3 phase. This would explain why the transformation of γ-Al2O3 into α-Al2O3 was not total after the complete heat treatment in Ar (Fig. V-10a). In this respect, the influence of the theoretical 5 vpm of H2O (and of 2 vpm of O2) contained in the Ar bottle on the potential delayed transformation of the metastable alumina into α-Al2O3 [40] has been disregarded since the content is too low to appreciate a marked effect in the conditions studied here. However, as introduced above, the laser source of the Raman was found to transform γ-Al2O3 into α-Al2O3 on some locations. the top coat upon aluminizing [22]. After the complete annealing in Ar, the oxide shells remained (visually) very thin with a weak sintering and the oxidation of the surface of the diffusion zone in contact with the top coat was insignificant. For the sample heat treated in synthetic air, the oxide shells of the microspheres were considerably thicker than the ones obtained in Ar. The microspheres also displayed a better sintering resulting in a more adherent and stronger top coat structure at the naked eye [23].

After the complete heat treatment in synthetic air, a thick α-Al2O3 TGO scale formed on the surface of the diffusion coating (confirmed with local Raman analyses). The compositions of the aluminized layers formed on the nickel substrate were similar in both atmospheres though.

The ceramic top coats were found to be more sintered after heat treatment in synthetic air than in Ar. However, in very local areas (e.g. corners of the samples), the coating did not completely cover the nickel substrate probably because the samples had not been chamfered (Fig. V-12a). This led to the formation of unprotective nickel oxide (NiO) and to some internal oxidation preferentially at the interface between the diffusion layer and the substrate (

Fig. V-12. Details of the complete thermal barrier systems on sample corners after (a) the diffusion heat treatment and (b) the complete heat treatment in synthetic air.

Fig. V-11. BSE cross-section images of the full thermal barrier systems elaborated on pure nickel substrate from Al microparticles. The thermal diffusivity of pure Ni with "PARTICOAT" (complete heat treatment in Ar) was measured to evaluate the thermal insulation behaviour of the alumina hollow spheres (cf. Annex 2).

Chapter was much faster in synthetic air than in Ar. The fast mass gain was attributed to the oxidation of exposed metallic Al. For the "Switch 2" sample, the introduction of synthetic air at the end of the two hours dwell at 700°C was accompanied with a rapid increase of the oxidation rate. However, the oxidation rate of "Switch 1" sample was greater in the range 700°C-1100°C (Stage III in Fig. V-2). This may indicate that the thickness of the γ-Al2O3 layer subjected to further crystallographic transformations was greater for the "Switch 1" sample than for the "Switch 2" one. The same evolution was observed by Trunov et al. by changing the initial size of the microparticles [2,16]: the smaller the microparticles, the faster the oxidation rate. When the temperature reached 1100°C, a stabilization of the mass was observed for both samples (stage IV in Fig. V-2). This suggested that most of the oxide shells had transformed into the stable α-Al2O3. It must be pointed out here that the formation of the TGO scale on the surface of the diffusion layer may also contribute to the mass variation of the samples.

Even though several thermally activated processes overlapped during the annealing, the mechanisms of oxidation proposed by Trunov et al. seems to successfully describe the systems investigated in this study [16]. This value is in good agreement with that of other studies dealing with the oxidation of Al particles, considering the present heating ramp of 5°C/min [2,3,5,6]. For the four samples annealed in Ar, the mass still decreased up to approximately 500°C and this may indicate that the desorption of water and of the organic binder [38] was slower than in synthetic air. Then, a constant mass gain was observed for the samples heat treated in Ar. The slope was in this case considerably smaller than that of the samples heat treated in synthetic air (Fig. V-17a). This mass gain was associated with the growth of the γ-Al2O3 layer. Firmansyah et al. have shown with HR-TEM observations that the thickening of the oxide shell occurred even in flowing Ar before the melting point of Al [11]. The same authors observed that the aluminium core expanded under almost no constraint in flowing Ar because of the inhomogeneous crystallization of the amorphous alumina shell [11]. Consequently, molten Al very rapidly diffuses outwards from the microspheres to react with the underlying nickel substrate upon annealing in Ar [21] (Fig. V-5a). These observations therefore suggest that the thickness of the γ-Al2O3 layer covering the Al particles was presumably much greater after heat treatment in synthetic air than in Ar when melting of Al occurred (Fig. V-17a). In synthetic air, the alumina shells formed on the smaller microparticles were resistant enough to countercurrent the expansion of the molten aluminium core [6,27]. This should explain why metallic Al was still present in the top coat structure in Fig. V-5b even at temperatures higher than the Al melting.

Fig. V-17. Comparison of the thermogravimetric curves upon heating (5°C/min) for the different systems (a) between 400°C and 700°C and (b) between 700°C and 1100°C.

The thermogravimetric curves obtained upon heating (5°C/min) between 700°C and 1100°C also differed significantly between the studied systems (Fig. V-17b). For the sample heat treated in synthetic air, the mass variation was in agreement with data from literature [2,5,6]. A second stepwise oxidation was observed corresponding to the growth of the γ-Al2O3 and its progressive transition into the stable α-Al2O3 polymorph. The slope change after reaching 1000°C was attributable to the predominant growth of α-Al2O3. The same observations were done for the "Switch 1" sample although the overall mass variation was less significant than that of the sample fully heat treated in synthetic air. This suggested that a fraction of the Al present in the top coat structure already reacted with the underlying substrate upon heating to 700°C in Ar. As a result, the diffusion layer was more homogeneous and the oxide shells were thinner (Fig. V-15a). For the "Switch 2" sample, there was barely no stepwise oxidation but a constant mass gain instead (Fig

Fig. V-18. SEM cross-section images of the top coat structures after the complete heat treatment (a) in Ar, (b) in synthetic air, (c) for the "Switch 1" sample and (d) for the "Switch 2" sample.

Chapter V -179-

 Influence of the atmosphere on the oxidation of Al microparticles and the release of Al:

The atmosphere (inert or oxidizing) had a tremendous influence on both the peripheral oxidation of

Fig. V-19. Mechanisms of formation of individual hollow alumina microspheres in Ar and in synthetic air proposed in this work.

Chapter V -180-

 Mechanisms of formation in Ar:

In Ar, the crystallization of the amorphous alumina layer is heterogeneous and the shell is not totally healed before the melting of Al [11]. Upon heating over the melting temperature of Al (about 660°C [41]), the molten Al core can expand under almost no constraint. As a result, liquid Al flows out through the nano-cracks and the microparticle progressively releases its Al content [1]. Simultaneously, the generated molten Al reacts with the underlying nickel substrate following the mechanisms described in Ref. [22] (i.e. dissolution of the metallic substrate and self-propagating high-temperature synthesis). In this work, Al microparticles were completely emptied after 2 hours dwell at 700°C transformation was incomplete at the sample surface after two hours dwell at 1100°C in Ar (Fig.

V-10a). The γ-Al2O3 to α-Al2O3 transformation is accompanied with a volume shrinkage of about 13.8%

because of a change of the crystallographic structure [6,20]. This volume shrinkage provokes the deformation of the thin alumina shell presenting a weak mechanical strength (Fig. V-19). This led to a broken hollow microspheres morphology observed on the surface (Fig. V-7a) and in cross-section (Fig.

V-18a) after the complete heat treatment in Ar.

 Mechanisms of formation in synthetic air:

Unlike Ar, the 2 of synthetic air atmosphere could heal the defective γ-Al2O3 layer by fast oxidation of the exposed metallic Al. This was confirmed by the stepwise oxidation occurring before the Al melting (stage II on Fig. V-17a). Upon further heating at a temperature greater than Al melting (i.e. 700°C), the mechanical strength of the healed alumina shell is high enough to countercurrent the expansion of the molten Al core [6,27]. This resulted in the thickening of the alumina shell probably by both outward diffusion of Al 3+ cations and inward diffusion of O 2-anions through the alumina shell [11]. This could explain that metallic Al was still contained in the microparticles after two hours dwell at 700°C (Fig. V-5b). As observed in Ar, the γ-Al2O3 to α-Al2O3 transformation occurring upon heating between 700°C and 1100°C induces a volume shrinkage in the alumina shell of about 13.8%.

Consequently, nano-cracks form in the alumina shell and molten Al can then flow out across these cracks [20,27]. Molten Al is then immediately oxidized because of the high 2 of synthetic air. This The annealing treatments performed in Ar for the Cr-Al double-layer system led to the same conclusions than those from Al microparticles. By performing the annealing treatment in Ar, molten Al flows out from the microparticles through the nano-cracks of the thin alumina shell (Fig. V-19). This resulted in the fast consumption of metallic Al contained in the microspheres and the limited growth of the alumina shells. Molten Al subsequently reacted with the Cr interlayer and formed AlxCry intermetallic phases in the slurry deposit (cf. paragraph V.3. in Chapter III). To increase the sintering between the alumina hollow spheres and consolidate the top coat structure, it is therefore necessary to introduce oxidizing species upon annealing (i.e. synthetic air).

Fig. V-20. BSE cross-section images of the top coat structure after annealing for (a) 400°C-3h + 650°C-5h, (b) 400°C-3h + 650°C-5h + 1000°C-3h and (c) 400°C-3h + 1080°C-6h in flowing

Fig. V-22. Details of the heat treatments carried out to investigate the influence of the atmosphere (Ar, synthetic air) on the formation of the Cr-Al double-layer system on pure nickel.

 Surface characterization of the complete thermal barrier systems after annealing:

The SEM characterization of the surfaces after the different annealing conditions is shown in Fig. V-24.

As expected from the thermogravimetric curves, the morphology of the top coat structure significantly changed. For samples S1 and S2, a dense and adherent top coat structure was formed on the surface. For these two samples, the alumina microspheres were very well sintered whereas a spongy microstructure was also observed on the top coat surface. This microstructure probably resulted from the fast oxidation of Al when shrinkage of the oxide shell occurred upon annealing in synthetic air [6,20]. These observations are in good agreement with the mechanisms proposed on

V-22) for the (a) complete heat treatment, (b) isothermal section at 400°C, (c) heating ramp at 5°C/min between 400°C and 1080°C and (d)

isothermal section at 1080°C.

Fig. V-24. Macrographs and SEM surface images of the samples after the different annealing conditions (pure nickel + 50Cr-50Al double-layer system).

Chapter V -187-

The XRD analyses are given in Fig. V-25. For samples S1 and S2, the XRD patterns were very similar and depicted the major contribution of the α-Al2O3 crystal structure. Some reflections of Cr2O3 were also identified. For samples S3 and S4, additional peaks were observed with the detection of Cr-rich phases formed in the slurry deposit. A cubic structure was identified, probably coming from a Cr solid solution, together with the two main peaks of the AlCr2 tetragonal structure. Additional peaks were labelled with question marks and were presumably related to AlxCry phases enclosed in the top coat structure after annealing. By correlation with SEM surface images, this suggested that the α-Al2O3 top coat was denser for samples S1 and S2. For the sample heat treated in Ar (S6), the detection of the Cr-rich phases (Cr solid solution and AlCr2) was more significant than α-Al2O3. The resulting coating system was similar for samples S1 and S2 with an inner layer composed of oxides ("OL") and a top coat layer ("TC") composed of hollow alumina spheres. Some bright particles were observed in the top coat and were attributed to Al-Cr intermetallic phases. For these two samples, the top coat structure was greatly adherent to the underlying substrate and the alumina hollow spheres were well sintered. For sample S3, the coating microstructure was close to samples S1 and S2.

 Cross-section characterization of the complete thermal barrier systems after annealing:

However, the alumina microspheres were less sintered and there was a greater amount of Al-Cr

Fig. V-25. XRD patterns of the surface for the different samples after annealing (pure nickel + 50Cr-50Al double-layer system).

Chapter V -188-intermetallic phases in the top coat structure. These observations are in good agreement with the XRD surface analyses (Fig. V-25). It must be noted that for these three samples, the formation of NixAly phases was completely hampered on the substrate surface.

The coating system obtained on sample S4 was completely different than the ones of samples S1 to S3. Some diffusion islands were randomly distributed on the substrate surface whereas the oxide layer was significantly thinner and discontinuous for sample S4. Since the heating ramp was performed in Ar for sample S4 (Fig.

Fig. V-27. BSE cross-section image of the sample S1 (50Cr-50Al double-layer system) after annealing in synthetic air and corresponding X-ray maps of (a) Ni, (b), Al, (c), O and (d) Cr.

III.3. Reactivity between Al microparticles and oxide powders (NiO, Cr2O3 and Al2O3)

Considering the aluminizing process of the Cr-Al double-layer system on pure nickel and in synthetic air, one could expect:

-the surface oxidation of pure nickel substrate into NiO, -the peripheral oxidation of the Cr microparticles in the slurry deposit into Cr2O3, -the peripheral oxidation of the Al microparticles and the oxidation of liquid Al in the slurry deposit into Al2O3.

 Overview of the aluminothermic reactions:

For a better comprehension of the mechanisms of formation of the full thermal barrier systems in synthetic air, the reactivity between Al microparticles and these oxides (NiO, Cr2O3 and Al2O3) was studied by DSC. The aluminothermic reactions have been extensively studied in the last decades and are well known for their high exothermicity to produce aluminides [48,49] or intermetallic-ceramic composite materials [50][51][52][53][54]. Since Al has a very high affinity for oxygen, it is able to reduce the most of metal oxides such as NiO or Cr2O3 [48,49,[55][56][57]. The aluminothermic reaction can be written as follows:

+ → + (Eq. V-1)

Depending on the initial stoichiometry of reaction (Eq. V-1), the reduction of the oxide can lead to the formation of corresponding metal or to an aluminide if an excess of Al is employed [48,49,53,58]. As a result, aluminides with different stoichiometries can be obtained by changing the ratio of the reactants, which also depend on the maximal temperature reached during the reaction. High energy ball milling was also found to decrease the ignition temperature and the activation energy of aluminothermic reductions [51,54,59]. The heating branches of the DSC curves obtained for the different mixtures are presented on Fig.

Table V-2. Composition of the different Al-oxide mixtures investigated by DSC and comparison with the stochiometric composition considering reaction (Eq. V-1).

Mixture

V-30. For the three mixtures, the endothermic signal of Al melting was observed at approximately 660°C. For the sake of clarity, the three systems will be investigated separately.

 Al + NiO system:

For the Al + NiO mixture, an exothermic event occurred before the Al melting with a maximal heat flow value at about 620°C. This event had already been reported by Galetz et al. with DTA analyses of Al and Ni microparticles in air and was ascribed to the oxidation of the microparticles [22]. Since it was only observed for the Al + NiO system and not for the powders taken separately, it appeared that this exothermic event was associated with a solid-solid interaction between Al and NiO. This exothermic event occurring before the Al melting was also reported by other authors by mixing Al and NiO powders [52,53]. According to Zhu et al., the initial Al/NiO reaction at about 600°C led to the formation of Ni without any other phases [52]. Therefore, this event was attributed to the solid state reduction of NiO by Al following the reaction [48,52,56,57]:

2 + → + (Eq. V-2)
After the melting of Al, a broad exothermic peak was observed. It initiated at approximately 750°C and ended at about 1020°C. This exothermic peak was associated with the liquid state reaction (Eq.

V-2) [52,59,63]. Note that reaction (Eq. V-2) is highly exothermic with a ∆ value of -968 kJ/mol [59].

A second exothermic event was observed at about 1070°C and likely corresponded to the two-stage reduction of NiO [63]. This suggested the formation of NixAly intermetallic phases because of the excess of Al in the initial powder mixture [48,52,58].

 Al + Cr2O3 system:

Unlike the mixtures of Al and Cr microparticles [42], there was no exothermic reaction at about 640°C for the Al + Cr2O3 system (Fig. V-30). It therefore appeared that the peripheral oxidation of the Cr microparticles hampered the solid state reaction between Al and Cr2O3. When the temperature reached approximately 800°C, the heat flow progressively increased to reach a maximum value at about 950°C. This probably referred to the liquid state reduction of Cr2O3 by Al following the reaction:

2 + → + 2 (Eq. V-3)
The reaction (Eq. V-3) is highly exothermic with a ∆ value of -541 kJ/mol [49]. A strong endothermic peak was observed at approximately 1035°C and convoluted with the broad exothermic peak on Fig.

V-30. This endothermic reaction was attributed to the decomposition of Al4Cr phase [42,[64][65][66]. This confirmed that the excess of Al reacted with Cr to form AlxCry phases upon heating. A second exothermic peak was recorded centred at about 1100°C suggesting a further transformation of the reacting products in the system.

 Reactivity between Al microparticles and the oxides:

The DSC experiments highlighted that the potential oxidation of the nickel substrate and that of the Cr microparticles had a tremendous influence on their reactivity with Al microparticles upon annealing.

In the vicinity of the Al melting, the oxide had a critical effect on its reactivity with Al (Fig. V-30).

Whereas the NiO reduction was initiated at temperatures below 600°C, the Cr2O3 reduction was shifted to temperatures over 800°C with a heating rate of 5°C/min in flowing Ar. This would explain why the Cr interlayer in the Cr-Al double-layer system had a critical effect on the microstructure of the complete thermal barrier system in synthetic air (Fig As expected, NixAly diffusion islands were randomly distributed on the substrate surface of samples S1 and S2 whereas the rest of the surface was covered with a thick oxide layer. Some internal oxidation was observed in the substrate and at the interface substrate-diffusion layer. For sample S3, a continuous diffusion layer was formed and the oxidation of the substrate was confined to a few areas. The top coat was composed of a well sintered intermetallic-ceramic composite. For sample S4, a continuous diffusion layer had formed with a thickness of approximately 90µm. The microstructure obtained for sample S5 was similar to that of sample S6 suggesting that interdiffusion between the nickel substrate and the AlxCry phases was nearly complete after one hour dwell at 1080°C.

Fig. V-31. Cross-section micrographs of the 25Cr-75Al double-layer system elaborated on pure nickel after the different annealing conditions (TC: Top coat, OL: Oxide layer, DL: Diffusion layer and IO: internal oxidation).

Chapter V -198-

These results therefore confirmed that the initial Cr-Al ratio has a critical effect on the microstructure of the coatings hence on their mechanisms of formation. Unlike the 50Cr-50Al double-layer system, Al was in excess in the 25Cr-75Al one. For samples S1 and S2, the excess of Al could diffuse towards the nickel substrate and react by combustion synthesis [22]. As a result, NixAly diffusion layers formed on the substrate surface. The formation of α-Al2O3 at the interface substrate-diffusion layer may indicate that reduction of NiO occurred upon annealing by Al following the reaction (Eq. V-2). Since Al was in excess in reaction (Eq. V-2), formation of NixAly intermetallic phases (i.e. NiAl) occurred [48,53,58] , (b), Al, (c), O and(d

Fig. V-32.BSE cross-section image of the sample S1 (25Cr-75Al double-layer system) after annealing in synthetic air and corresponding X-ray maps of (a) Ni

 Influence of the atmosphere on the coating microstructures:

The composition of the atmosphere did not seem to play any significant role on the formation of the diffusion layers from Al microparticles. The diffusion layers were formed by self-propagating hightemperature synthesis (SHS) by dissolution of nickel in molten Al and subsequent NixAly intermetallic phases formation. The coating matrix was mainly composed of Ni2Al3 phase at 700°C and of β-NiAl phase at 1100°C. However, the microstructure of the top coat layer was significantly changed with the introduction of synthetic air into the reaction chamber. Synthetic air indeed promoted the thickening of the alumina shells and their sintering with extensive formation of α-Al2O3. This is expected to increase the mechanical strength of the top coat structure, in particular its erosion resistance.

Chapter V

-201-

 Influence of the Cr interlayer on the formation of complete thermal barrier systems:

The influence of a Cr interlayer between the nickel substrate and the Al microparticles was also Chapter, two different modes of degradation will be investigated. The Type-II hot corrosion generally occurring within the temperature range 650-800°C [1][2][3] and the high-temperature oxidation becoming predominant at temperatures higher than 1000°C [4]. The main objective of the present chapter was to give an overview of the potential applications of the diffusion coatings designed during this study based on their environmental resistance.

I.1. Simulation of Type-II hot corrosion conditions at laboratory scale

The Type-II hot corrosion conditions are schematically illustrated in Fig. VI-1. This mode of degradation is observed at temperatures lower than the melting point of Na2SO4, which is the most common salt deposit for aircraft gas turbines [4]. At such temperatures (i.e. 650-800°C), Na2SO4 can condensate on the metallic surfaces and form a solid and sticky deposit. The combustion gases produced for the power generation also play an important role on the hot corrosion process. A critical is indeed required to provoke the corrosion attack and initiate the Type-II hot corrosion process.

Therefore, when the 3 is high enough in the gas composition (about 3x10 -4 atm. [5]), the sulfation of NiO can occur following the reaction (Eq. VI-1):

+ → (Eq. VI-1)

Fig. VI-1. Schematic illustration of the conditions to meet Type-II hot corrosion of metals and alloys (also called "under-deposit hot corrosion").

N.B.: in this work, the term "sulfation" refers to the formation of a metal sulfate phase from a metal oxide in presence of SO3 whereas the term "sulfidation" refers to the formation of nickel sulfides from the pure metal.

The formation of NiSO4 following reaction (Eq. VI-1) progressively saturates the salt deposit and leads to a liquid solution of Na2SO4-NiSO4 (eutectic temperature of 671°C [6][7][8]). In dry air, the 3 is not high enough to stabilize Na2SO4-NiSO4 liquid solution and the hot corrosion attack is therefore suppressed. The stability of the different phases can generally be predicted by thermodynamics using stability diagrams where 2 and 3 are used as coordinates (Fig. VI-2a).

At laboratory scale, it is therefore of utmost importance to control the 3 within the gas composition to meet the Type-II hot corrosion conditions [9]. A SO2-containing atmosphere is generally used to perform the hot corrosion tests. At 700°C (the test temperature), the oxidation of SO2 into SO3 following reaction (Eq. VI-2) is very slow. This implies the use of a catalyst to foster the oxidation of SO2. The catalyst is generally placed in the reaction chamber upstream of the samples.

SO2(g) + O2 g → SO3(g)

(Eq. VI-2)
In this work, the hot corrosion tests were performed with the experimental set up presented in Fig. (Eq. VI-3) and of Cr2O3 (Eq. VI-4) should not be observed in the experiments.

Al2O3(s) + 3SO3 g → Al2(SO4)3 (s) (Eq. VI-3)

Cr2O3(s) + 3SO3 g → Cr2(SO4)3 (s) (Eq. VI-4)
These thermodynamic considerations can be used to interpret the mechanisms of corrosion as a function of the 2 and 3 variations. In this work, the samples were half-embedded in Na2SO4 to reproduce the Type-II hot corrosion conditions [10]. This method is generally more aggressive than real hot corrosion conditions but is interesting for the screening of different compositions. Using the half-embedding method, it is also possible to analyse different zones on the sample: the atmosphere side (no contact with Na2SO4), the salt side (embedded in Na2SO4) and the interface between the two areas.

I.2. Oxidation resistance of simple and modified aluminide coatings

The high-temperature oxidation of metals and alloys is another important concern in gas turbine engines. The oxidation or "dry corrosion" becomes predominant at temperatures beyond 1000°C

(above the salt dew point). The oxidation phenomenon is usually more predictable than hot corrosion and is generally simulated at laboratory scale using muffle furnaces or cyclic oxidation rig tests in dry air. Even though both isothermal and cyclic conditions are used to simulate oxidation at laboratory scale [11], cyclic oxidation tests are closer to real conditions by inducing thermal stresses in the growing oxide scales [12].

The oxidation resistance of simple and modified aluminide coatings has been extensively studied in literature. The test temperature is generally selected regarding the application of the coating. If the oxidation resistance of bond coats for thermal barrier coatings is usually performed within the range 1150°C-1200°C [13][14][15][16], the oxidation tests for environmental bond coats are carried out within the range 1000°C-1100°C [17][18][19][20][21][22]. These temperatures are usually well recognized to simulate real conditions even at the atmospheric pressure since many metallurgical transformations observed on gas turbine components are thermally activated [21,23]. However, it is important to consider here that the temperature can vary locally on the surface of a component [24]. This can therefore lead to different modes of degradation on the same component [24,25]. If the simple aluminide coatings present a good oxidation resistance, they are generally very susceptible to both Type-I and Type-II hot corrosion [26][27][28]. The modification of simple aluminide coatings with Cr therefore appeared as a good alternative to confer both low temperature and high temperature resistance to the underlying material [26][27][28][29][30][31].

In this work, the oxidation resistance of Cr-modified coatings elaborated on a nickel-based superalloy (CM-247 LC) was investigated. The potential application of such coatings can be as low as 700°C to confer Type-II hot corrosion resistance and as high as 1100°C for the first stages of the low-pressure turbine (LPT). The oxidation temperature was fixed at 1000°C in this study since the oxidation of Crmodified coatings was generally investigated at this temperature in literature [19,32]. The main objectives of this chapter were to study the surface stability of the coating and the microstructural evolutions that may happen at elevated temperatures. This could therefore inform on the maximal capability of the coating to confer high-temperature protection to the underlying material.

The Type-II hot corrosion resistance of the Cr-modified coating was investigated at 700°C on pure nickel for a better comprehension of the corrosion mechanisms. The use of superalloys with complex chemistry generally gives incomplete results in hot corrosion studies since each alloying element can adversely affect the corrosion mechanisms [33]. Ar atmosphere have already been described on pure nickel [34,35] and nickel-based superalloys [35].

The coating formation was found to be fast involving liquid phases and a strong exothermic reaction between the nickel-based substrate and the Al microparticles [36]. For the simple aluminide coating -212-

The surface of the slurry coatings was also characterized before oxidation tests to obtain the initial state of the surface after the light grit blasting process (#220 mesh alumina). Both coatings presented a relatively rough surface after grit blasting which may increase the reactive surface area upon oxidation [13]. Several EDS window measurements were done on the coating surface to analyse the chemical composition. The mean concentration was 25O-46Al-29Ni (at.%) for the simple aluminide and 22O-23Al-23Cr-32Ni (at.%) for the Cr-modified one. The O concentration on the surface suggested that the coatings were slightly pre-oxidized after elaboration. For a better comprehension of the hot corrosion tests performed on pure nickel, its oxidation behaviour in synthetic air was first investigated. The oxidation of high purity nickel has been extensively studied in literature in both dry air [39] and pure oxygen [40,41]. NiO presents different morphologies depending on the temperature and the oxidation time and can be divided in three families: faceted grains, cellular and platelets [40]. At 700°C, NiO generally presents a cellular morphology with high porosity [8,39,40]. At temperatures lower than 800°C, NiO forms by predominant outward diffusion of Ni cations through the oxide scale [39].

In this work, high purity nickel (99. were dispersed within the nickel sulfide layer and suggested that a Ni-S liquid solution was present at temperature (i.e. 700°C). The darker areas observed above the nickel sulfide layer were attributed to nickel oxide resulting probably to a higher 2 in this region at temperature [46]. These observations are in good agreement with the results obtained by Kofstad and co-workers [43][44][45]. -219-

The EDS spot measurements done in cross-section are summarized in Table VI-2. The nickel sulfide layer was mainly composed of the Ni3S2 phase. The S concentration was very low in the nickel substrate with less than 1 at.% S at 10µm from the interface between the substrate and the sulfide layer (Table VI-

c) Discussion

The oxidation of high purity nickel in synthetic air at 700°C led to the formation of a porous NiO scale with a cellular morphology as reported in literature [8,39]. The growth of the NiO scale presumably occurred by major outward diffusion of Ni cations towards the gas atmosphere through the NiO grain boundaries [39]. When 0.5% SO2 was introduced in synthetic air, the oxidation rate of pure nickel considerably increased (Fig. VI-8). The higher porosity and growth rate of NiO scales in presence of sulfur was unambiguously due to inward transport of SO2-SO3 through the scale [8,47]. The sulfation of NiO in synthetic air + 0.5% SO2 following reaction (Eq. VI-1) was therefore promoted. These that NiO was permeable to sulfur-bearing molecules [8,47]. The formation of a mixed NiO + Ni3S2 scale has also been reported in high SO2-containing environments [6]. In such microstructures, more than 90% of the scale consists of NiO in which the sulfides are present as a continuous network. The inward transport of SO2-SO3 also provided a sufficient S concentration (i.e. 2 ) for sulfidation to occur at the scale/substrate interface as observed by Lillerud and Kofstad [6]. The sulfidation of nickel can be decomposed in two consecutive reactions following equations (Eq. VI-5) and (Eq. VI-6) [6]: aluminide coating was close to the one of γ-Al2O3 at 700°C [48]. For the Cr-modified coating, the oxidation rate was approximately 30 times higher than that of the simple aluminide one. The value for this coating was also higher than that of pure chromium at 700°C [49]. This could arise from the higher reactive surface area of the coating because of its relatively high roughness (Fig. VI-3b). modified coating was probably the cause for the higher mass gain of the sample upon oxidation since the O uptake was also greater (Fig. .

XRD and Raman analyses were conducted to identify the crystal structures and the oxide phases formed on the slurry coatings after oxidation at 700°C. The results are summarized in Table VI-4. For the simple aluminide coating, only the -NiAl phase was detected by XRD using grazing incidence which suggested that the scales formed were very thin [7,50]. From Table VI-

c) Discussion

The application of slurry coatings was found to considerably increase the oxidation-sulfidation resistance of high purity nickel in synthetic air + 0.5% SO2 at 700°C. This was attributed to the formation of a thin and adherent Al2O3 scale on the coating surface upon exposure. The oxidation tests of the coatings performed in synthetic air have shown that the duration of the transient oxidation stage was different for the two coatings (cf. Table VI-4). After 100 hours exposure, transition aluminas were still identified for the simple aluminide coating whereas α-Al2O3 was already formed for the Crmodified coating after 3 hours exposure. These results are in good agreement with literature where Cr has long been acknowledged to limit the growth of metastable aluminas and to foster the nucleation of α-Al2O3 [48,52,53]. The presence of Cr was therefore expected to limit the inward diffusion of oxygen and internal oxidation as shown by Brumm et al. [52]. It has to be noted that Cr2O3 was also detected by XRD and Raman micro-spectrometry on the Cr-modified coating since Cr-rich phases were present on the surface (Fig. VI-3c). The formation of Cr2O3 may also have fostered the nucleation of α-Al2O3 since these two oxides have the same crystallographic structure [53,54,59]. -233-

III. Type-II hot corrosion resistance of uncoated and slurry coated nickel at 700°C

In paragraph II, the influence of 3 on the oxidation resistance of uncoated and slurry coated nickel was investigated at 700°C. The application of both the simple aluminide and Cr-modified coatings increased the resistance of nickel to oxidation-sulfidation. In this paragraph, the influence of Na2SO4 addition in the system was studied by half-embedding of the samples into Na2SO4 salt (cf. Chapter II).

As a result, type-II hot corrosion conditions were fulfilled for these experiments.

III.1. Hot corrosion of pure nickel at 700°C

 Half-embedding of pure Ni in Na2SO4 and exposed in dry air:

As presented in Fig. VI-1, a solid salt and a sufficient 3 are required to meet the type-II hot corrosion conditions. The half-embedding method was first carried out in dry air to observe the behaviour of pure nickel in these conditions and to substantiate later the influence of 3 on the corrosion attack.

According to thermodynamic considerations, the type-II hot corrosion can only occur in the SO2containing atmosphere in this work since Na2SO4-NiSO4 eutectic is not stable in dry air (Fig. VI-2).

This should prevent the formation of NiSO4-Na2SO4 eutectic phases responsible for the hot corrosion phenomenon.

The cross-section micrographs of high purity nickel samples half-embedded in Na2SO4 are given in Fig. nodules however suggested that the 3 was high enough to stabilize sulfides. In the present system, additional S can only be obtained from dissociation of the sulfate, following equation (Eq. VI-7) [8]:

Na2SO4 s → Na2O(s) + SO3(g) (Eq. VI-7) SO3 g ⇆ S2(g) + O2(g) (Eq. VI-8)
The absence of a significant 3 in the test conditions prevented the stabilization of Na2SO4-NiSO4 solution and required reaction (Eq. VI-7) to proceed from the left to the right (i.e. the sulfate releases some SO3) [8]. The characterization of the salt attested that it had remained solid after exposure and that the formation of Na2SO4-NiSO4 eutectic phases was not observed. Therefore, the existence of The extreme surface of the NiO + Ni3S2 mixed scale was covered with a NiO layer infiltrated with Na2SO4-NiSO4 as observed by other authors [6,8]. This confirmed the formation of liquid phases at the test temperature (i.e. 700°C). This NiO infiltrated layer was observed all over the sample on the atmosphere side. Extensive sulfidation of the substrate was also observed in cross-section. These observations confirmed that the Type-II hot corrosion conditions were fulfilled with the present experimental set up [8].

The SEM cross-sections of pure nickel samples subjected to Type-II hot corrosion conditions (in the -238-By comparison with the microstructures observed in paragraph II.2, more nickel sulfides were detected in the corrosion scale after exposure to Type-II hot corrosion conditions. It can therefore be concluded that in the test atmosphere, Na2SO4 accelerated the corrosion process probably by increasing the S activity (i.e. 2 ) within the scale and at the metal/scale interface.

III.2. Hot corrosion of simple and Cr-modified slurry aluminide coatings on pure nickel

The half-embedding method was also conducted on slurry coated nickel to investigate the hot corrosion resistance of the simple and Cr-modified coatings. The oxidation tests presented in paragraph II.3 confirmed that the application of the slurry coatings considerably increased the oxidation-sulfidation resistance of the nickel substrate. This was attributed to the formation of protective oxide scales on coating surfaces such as Al2O3 and Cr2O3. Whereas the Cr-modified coating was not attacked after 300 hours exposure in synthetic air + 0.5% SO2, the simple aluminide coating was prone to localized attack in the form of protrusions with internal oxidation. This was attributed to the rapid establishment of an α-Al2O3 scale on the Cr-modified coating that provided an effective barrier for SO2-SO3 molecules. These results will help to interpret the hot corrosion mechanisms with the addition of Na2SO4 in the system.

 Half-embedding of slurry coated Ni in Na2SO4 and exposed to dry air:

The half-embedding method was first performed in dry air to later investigate the influence of 3 on the corrosive attack. The top view of the samples half-embedded in Na2SO4 suggested that there was When the half-embedded samples were exposed to synthetic air + 0.5% SO2, the results were significantly different as can be seen in the crucibles (Fig. VI-32). For both coatings, there was no interaction with the salt after 20 hours suggesting that the incubation period was still ongoing. After 100 hours, the samples were highly corroded with the formation of nodules and needle-like products at the surface. The colour of the salt also changed into yellow in the corroded area. Such yellow colour was attributed to the formation of mixed Na-Ni sulfates by Lillerud and Kofstad [6]. where it sealed the gaps between the intermetallic phases. There was no formation of sulfides in such areas and that was ascribed to the presence of both β and γ' phases in the outer region of the coating [60].

Cr2O3 was also found to precipitate within the mixed sulfate layer (cf. Chromite (NaCrO2) is therefore the dominant basic solute under highly reducing conditions as expected near the oxide/salt interface [63].

Fig. VI-36. BSE cross-section images of the Cr-modified coating elaborated on pure Ni after 100 hours exposure to Type-II conditions at 700°C. Images were taken in a region considered in incubation (cf. Fig. VI-33d

).

-244-

The basic dissolution of Cr2O3 prevailing underneath the salt film at low 2 can be written following (Eq. VI-9):

Cr2O3(s) + O 2-→ 2 CrO2 -

(Eq. VI-9)

The chromite ions then oxidize in chromate ions in region with high oxygen activities following (Eq. VI-10):

2 CrO2 -+ O2 + O 2-→ 2 CrO4 2-

(Eq. VI-10)

Acting as a strong acid solute, the presence of both chromite and chromate anions serve as a buffer for the salt chemistry mainly at the scale/salt interface [63]. The chromate anions were also expected By comparison with the results obtained for the simple aluminide coating, it therefore appeared that the presence of Cr-rich phases was found to passivate the surface of the coating through the formation of Cr2O3 and to prevent the initiation of the hot corrosion process.

Fig. VI-37. BSE cross-section image and corresponding X-ray maps for nickel slurry coated with the Cr-modified coating after 100 hours Type-II hot corrosion exposure at 700°C.

After 100 hours exposure in Type-II conditions, some pits were observed on the Cr-modified coating (Fig. VI-33d). X-ray maps were done in a pitted region to interpret the corrosion mechanisms (Fig.

VI -37). The formation of Cr sulfides was highlighted at the reaction front between the pit and the metallic coating and was probably due to the extensive Al depletion forming a continuous γ' layer underneath the oxide scale [60]. As a result, depletion of the protective element (i.e. Cr) occurred in the outer region of the coating. Any oxide scale that may form in these conditions were not protective and rapidly dissolve in the sulfate phase. The dissolution of the metallic elements in the pit region was highlighted with the X-ray maps (Fig This probably illustrates the existence of a 2 .gradient across the pit region with a low 2 at the scale/coating interface and a high 2 at the salt/pit interface. Consequently, the oxide with the higher affinity for oxygen (i.e. Al2O3) precipitates in the inner region of the pit.

The existence of a 2 .gradient across the pit region probably affects the gas composition such as the decomposition of SO3 may be observed across the scale following (Eq. VI-12): SO3(g) ⇆ S2(g) + O2(g)

(Eq. VI-12)

The increased S activity (i.e. 2 ) at the scale/coating interface fosters the sulfidation of the metallic coating. Even though the sulfidation of the Cr-modified aluminide coating was observed after 100 hours exposure, the propagation was considerably lower than the corrosion attack of the simple aluminide coating (Fig. VI-33b). It is highly likely that the formation of Cr sulfides consumes the S at the reaction front thus preventing the formation of Ni3S2 at the metal/scale interface [65]. This considerably limited the formation of eutectic Ni-Ni3S2 solution and buffered the corrosion attack [65].

From the hot corrosion tests of pure Ni, it can thus be concluded that Na2SO4 increased the S activity within the growing scales and at the metal/scale interface. Increasing the S activity led to higher 2 and thus to higher sulfidation rates. This was particularly evident by comparison of the scales after exposure in the SO2-containing atmosphere only and the Type-II conditions (Fig

 Hot corrosion resistance of the slurry aluminide coatings elaborated on pure Ni:

By applying slurry coatings on the pure Ni substrate, the corrosion attack was significantly decreased.

The slurry coated samples half-embedded in Na2SO4 and exposed to dry air were not attacked and underwent "normal" oxidation. When the half-embedded samples were exposed to synthetic air + 0.5% SO2, corrosion was observed after a first incubation period which lasted at least 20 hours in our -249-It should be noted that the sulfidation of the Ni substrate was limited after 100 hours exposure. This could be explained by the different modes of transport of S during the corrosion attack [47]. For the sulfidation of pure Ni, the transport of S-bearing molecules (SO2-SO3) through the physical defects of NiO was found to control the attack [47]. In the present case, an intermetallic layer was still present on the substrate surface, which prevented the permeation of gas molecules before the attack. It could therefore be proposed that the S enrichment at the coating-substrate interface occurred by dissolution mechanisms [47]. The presence of liquid phases at temperature (i.e. 700°C) can explain the lateral decohesion observed in the microstructure after cooling since the phases must present different coefficients of thermal expansion (CTE).

For Ni samples coated with the Cr-modified coating, the corrosion attack was much slower than with the simple aluminide coating. After 100 hours exposure, some pits were randomly distributed on the sample surface (Fig. VI-33d) whereas some areas were still unaffected by the presence of Na2SO4 (Fig.

VI -36). The superior corrosion resistance of the Cr-modified coating was ascribed to the rapid formation of Cr2O3 within the Cr-rich layer. It appeared also that Cr2O3 was precipitating in the layer of mixed sulfates (Fig. VI-36b). This may indicate the presence of chromate anions in the sulfate layer as proposed in reaction (Eq. VI-9) [61]. The precipitation of Cr2O3 was found to seal the reactive sites and thereby to isolate the salt mixture from the coating layer as proposed by Otsuka and Rapp for Type-I hot corrosion tests [61]. Cr2O3 also acted as nucleation sites for α-Al2O3 as suggested by the X-ray maps on The oxidation resistance of Cr-Al double-layer coatings elaborated on CM-247 LC superalloy has been investigated under cyclic and isothermal conditions. The mechanisms of formation of the coatings were discussed in Chapter IV. The initial Cr/Al ratio of the slurry deposit was found to exert a critical influence on the resulting microstructure of the coatings. By adjusting the Cr/Al ratio and optimizing the thermal treatment, it was thus possible to obtain outward grown coatings with significant enrichment in Cr and inward grown coatings with a complete thermal barrier system. The oxidation resistance of these two systems was investigated in this work. However, only the results obtained for the low activity high temperature (LAHT) coating will be discussed here.

IV.1. Characterization of the low-activity high-temperature (LAHT) coating

The mechanisms of formation of the LAHT coating from the Cr-Al double-layer system have been presented in paragraph II.2 of Chapter IV. After deposition of the slurries by air brush (close to the 50Cr-50Al ratio in wt.%), the samples were heat treated in Ar following the conventional repair heat treatment (400°C-3h, 1080°C-6h). A secondary electrons cross-section image of the resulting coating The slurry coated samples were exposed to cyclic oxidation with 1h-cycles (45 minutes at 1000°C and 15 minutes under rapid air cooling). This oxidation test was expected to simulate industrial service conditions with the combination of both oxidation attack and thermally induced stresses [12]. Up to six slurry coated CM-247 LC samples (cf. Fig. VI-40a) were introduced in the Delta Thermique rig (cf.

Chapter II) for the cyclic oxidation tests. Among these six samples, five were gently grit blasted after heat treatment to remove residual aluminizing materials whereas an additional sample was ground down to SiC #4000 paper grade. Tolpygo et al. have shown that grit blasting can adversely affect the oxidation behaviour of a Pt-modified aluminide coating during cyclic oxidation at 1150°C [13]. Grit blasting was attributed to contaminate the coating surface by various impurities (mainly alkali elements and Ti) that become incorporated within the α-Al2O3 scale upon oxidation. This accelerated the growth rate of α-Al2O3 and decreased the spallation resistance during cyclic oxidation [13]. Therefore, The mean values for the grit blasted and the polished samples calculated after the transition period are given in Table VI-5 and compared with data from literature. The values of the Cr/Al slurry coating were in good agreement with the oxidation of Al2O3-former alloys at 1000°C (Table VI -5). Grit blasting of the coating surface was not found to significantly change the cyclic oxidation behaviour of the coating at 1000°C. Grit blasting was found to promote the formation of α-Al2O3 rather than metastable alumina polymorphs at the intermediate temperature range [13]. The rapid formation of α-Al2O3 on grit blasted samples could possibly decrease the initial oxidation rate of the coating by comparison with a polished sample. In this paragraph, only the grit blasted samples were characterized after oxidation tests since the surface preparation was not found to exert a significant influence on the oxidation kinetics.

a

This work Literature

Conditions (g².cm -4 .s -1 ) System (g².cm -4 .s -1 ) Ref.

Grit blasting 1.9x10 -13 K417G + Y/Cr-modified aluminide coating 4.7x10 -13 [19] SiC #4000 3.0x10 -13 DD32M + simple aluminide coating 5.5x10 -13 [20] NiAl-34Cr (at.%) 2.0x10 -13 [48] The SEM was carried out to investigate the morphology of the oxide scales developed on the Cr/Al

 Oxidation resistance of low-activity Cr-Al coatings elaborated on CM-247 LC:

The oxidation of Cr-modified aluminide coatings elaborated on CM-247 LC was also investigated upon cyclic and isothermal conditions at 1000°C. The coating presented high surface stability forming a thin and adherent α-Al2O3 scale on the surface. The mass gain curves confirmed the parabolic growth of the oxide scale. The thermal stresses induced by cyclic conditions were not found to significantly influence the oxidation behaviour of the slurry coating. The α-Al2O3 scales were generally well adherent and little spallation was observed. The Cr was found to segregate in the outer region of the coating upon oxidation. This might have a beneficial effect against hot corrosion considering the conditions studied here and the literature. The main degradation mechanism observed for the Crmodified coating was the formation of a secondary reaction zone (SRZ) in the sub-region of the diffusion coating. This zone was formed because of Al inward diffusion from the coating layer toward the superalloy substrate. The Al enrichment decreased the solubility of refractory elements forming then needle-shaped precipitates. This also destroyed the cube-like γ-γ' microstructure of CM-247 LC in this region. The SRZ thickened with increasing oxidation time. However, this did not impair the surface stability of the coating and little to γ' transformation was noticed after long term exposure (1000 cycles and 1000 hours at 1000°C).

To conclude, the application of slurry coatings considerably increased the high-temperature corrosion and oxidation resistance of the underlying substrates. The modification of the slurry aluminide coating with Cr was highly beneficial for hot corrosion resistance in SO2-containing atmosphere.

Conclusions and perspectives

The main objectives of the present study were first to control the aluminium activity during slurry aluminizing of nickel-based alloys by adding Cr microparticles in the slurry compositions, to get lowactivity aluminide coatings, and then to study the effect of introduced Cr on the high-temperature oxidation and the hot corrosion resistance of the coating systems.

In a first step, it was necessary to control the exothermic reactions occurring between Al microparticles and nickel-based materials upon slurry aluminizing and to investigate the influence of the coating process parameters on the formation of metallic bond coats and of complete thermal barrier systems. In a second step, Cr was added in the slurry compositions and was found to be highly beneficial on several different aspects summarized below after. From these results and based on literature, three compositions were then selected to investigate the role of Cr upon aluminizing of pure nickel substrate (67Al33Cr, 44Al56Cr and 20Al80Cr mixed slurries in wt.%). Upon aluminizing, Al preferentially reacted with Cr in the slurry rather than with the nickel substrate. This considerably limited the exothermic reactions between Al microparticles and the nickel substrate at low temperature (i.e. 650°C) by comparison with a Cr-free slurry composition.

Consequently, the addition of Cr hampered the formation of Al-rich nickel aluminides at low

Conclusions and perspectives

-270-temperature. With further annealing at high temperature (i.e. 1000°C or 1080°C), interdiffusion between the nickel substrate and the AlxCry phases synthesized in the slurry deposit was promoted.

The predominant outward diffusion of nickel at high temperatures allowed to progressively dissolve the AlxCry phases leading to the direct formation of the -NiAl phase with undissolved Cr.

Thanks to the versatility of the slurry coating method, different coating architectures were then tested to optimize the coating properties. Two systems were thus investigated on pure nickel, the Al-Cr and the Cr-Al double-layer systems. The mechanisms of formation of the coatings were described based on the experimental observations and the DSC measurements. The Cr-Al double-layer system appeared as the most adequate functional design for the aluminizing of nickel-based materials.

The results obtained on pure nickel were then extrapolated to a Ni-Cr alloy (Ni-20Cr-1.5Si in wt.%)

to CM-247 LC superalloy. The addition of Cr was found to allow the control of the exothermic reactions with the substrates upon aluminizing. As observed with pure nickel, the Cr-Al double-layer system appeared as the best coating design. It appeared that the initial Al-Cr ratio of the slurry deposit had a tremendous influence on the final coating microstructure. The threshold to achieve "low-activity" coating microstructures on CM-247 LC was between 56Cr-44Al and 67Al-33Cr compositions (in wt.%). Based on literature and on the experimental observations, the mechanisms of formation of the diffusion coatings could be explained. The results generally agree with the standard CVD aluminizing processes with two major coating microstructures: "low-activity" and "high-activity" coatings. The inert carbides present in the superalloy microstructure were used as natural markers to identify the initial surface of the substrate before aluminizing.

For the low-activity coatings (56Cr-44Al composition), the growth of the coating occurred by predominant outward diffusion of nickel at high temperature (i.e. 1080°C). The direct formation of the -NiAl phase occurred and the coating presented a typical duplex microstructure with an outer additive layer formed above the initial surface of the substrate and an inner interdiffusion resulting from the nickel withdrawal. After annealing, undissolved Cr remained in the additive layer of the coating because of its limited solubility in the -NiAl phase.

For the high-activity coatings (33Cr-67Al composition), the growth of the coating occurred by Finally, the environmental resistance of the slurry coatings developed in this work was tested in hot corrosion and oxidation conditions. The use of pure nickel in the SO2/SO3-containing atmosphere confirmed that the 3 of the reaction chamber was sufficient to sustain Type-II hot corrosion conditions at 700°C. The application of slurry coatings (simple aluminide and Cr-modified coatings)

was found to considerably increase the oxidation-sulfidation behaviour of pure nickel. After a total exposure of 300 hours in synthetic air + 0.5% SO2, the Cr-modified coating was not attacked whereas the simple aluminide coating was prone to localized attacks in the form of protrusions. The growth of these protrusions was found to occur by inward diffusion of O and outward diffusion of Al forming porous Al2O3 precipitates within the nickel aluminide phases. In Type-II hot corrosion conditions (in presence of Na2SO4), the slurry coatings presented an incubation period greater than 20 hours. After initiation of the corrosion attack, the propagation rate was significantly lower for the Cr-modified coating than that of the simple aluminide coating. The beneficial effect of Cr on the Type-II hot corrosion resistance was attributed to the saturation of the salt mixture with chromate anions and the subsequent precipitation of Cr2O3 from the melt. The formation of Cr2O3 was found to seal the reactive sites of the coatings and isolating them from the salt mixture. Such observations had not been reported yet in Type-II hot corrosion conditions. This was also ascribed to foster the establishment of a dense and continuous α-Al2O3 scale on the surface.

The cyclic and isothermal oxidation behaviour of the low-activity coatings was tested on CM-247 LC at 1000°C. The coatings presented high surface stability forming a dense and continuous α-Al2O3 scale on the surface. The thermal stresses induced under cyclic conditions had little influence on the spallation resistance of the α-Al2O3 scale. For long term exposure, outward diffusion of Ti, Ta and Hf from the CM-247 LC substrate occurred and these elements became incorporated within the α-Al2O3 scale. However, these enrichments had little influence on the performances of the coating after exposure for 1000 hours and 1000 cycles at 1000°C. The Cr was found to segregate in the outer region of the coating and may have a beneficial effect under hot corrosion conditions. The main degradation mechanism observed for the low-activity coating was the formation and the growth of a secondary reaction zone (SRZ) in the sub-region of the diffusion coating. This SRZ resulting from the inward diffusion of Al decreased the solubility of refractory elements, which led to the formation of needle-shaped precipitates.

Conclusions and perspectives

-273-However, after this work, some questions have not yet been answered that can be considered as interesting perspectives of investigation in future works. A non-exhaustive list of these perspectives is given below after.

Concerning the reactivity of the metallic powders and the synthesis of the Al-Cr donors:

-Thermodynamic calculations to determine the Al activity in the different slurry compositions and at different temperatures.

-Correlation of the thermodynamic calculations with experimental observations of slurry aluminizing.

For the reactivity between Al-Cr donors and nickel-based materials:

-Synthesis of bulk Al-Cr intermetallic compounds of different stoichiometry (e.g. Al7Cr, Al4Cr, Al8Cr5, AlCr2) to investigate their reactivity with nickel-based materials.

-Investigate the reactivity of the Al-Cr mixed slurries with a nickel-based superalloy.

For the aluminizing of nickel-based superalloys:

-Extrapolate the results obtained on CM-247 LC superalloy to other nickel-based superalloys.

-Better control of the aluminizing process based on thermodynamic calculations.

-Aluminization tests on real turbine components (e.g. turbine blades).

For the influence of the atmosphere on the formation of complete thermal barrier systems:

-The determination of the partial pressure of oxygen in the reaction chamber upon coating formation should give valuable information on the mechanisms of oxidation and of sintering of the individual aluminium microspheres.

-Extrapolate the results obtained on nickel-based superalloys.

-To form intermetallic-oxides composite materials, it would be interesting to propose a coating design based on aluminothermic reactions (with NiO and Cr2O3 especially).

-Further determine the thermal insulation performances of the ceramic top coat structures using the laser flash method (LFA).

For the environmental resistance of the slurry coatings:

-Investigate the early-stage exposure of the slurry coatings under Type-II hot corrosion conditions for a precise identification of the initiation stage.

-Use XPS (X-ray Photoelectron Spectrometry) or other extreme surface techniques to precisely identify the oxide scales formed at low temperature (i.e. 700°C) and confirm Raman analyses.

-The influence of the Cr-rich phases (i.e. residues of the aluminizing process) on the hot corrosion resistance should be evaluated.

Conclusions and perspectives

-274--The different coating microstructures obtained on CM-247 LC (Cr-modified aluminide coatings) give interesting perspectives for hot corrosion resistance. The distribution of the Cr-rich phases in the -NiAl coating matrix is presumed to influence the hot corrosion resistance of the coating.

-Realize Type-II and Type-I hot corrosion tests of the Cr-modified aluminide coatings developed during this work on nickel-based superalloys.

-Increase the oxidation time upon cyclic conditions to reach the limit of the slurry coatings developed.

For the erosion resistance and the repairability of the coatings:

-The thermal barrier coatings elaborated in Ar presented low erosion resistance. Scratch tests and tests with erosive particles (e.g. grit blasting) could therefore be conducted to evaluate the erosion resistance of the complete thermal barrier systems elaborated in air. The increased sintering of the alumina hollow microspheres should be beneficial on this aspect.

-In the industrial repair processes, grit blasting is commonly employed to remove degraded thermal barrier coatings. It would therefore be valuable to investigate the potential repairability of the thermal barrier systems following industrial standard processes.

-The chemical stripping of the nickel aluminide coatings is also feasible for repair issues.

Annexes

Annex 2

The laser flash analysis (LFA) was carried out to measure the thermal diffusivity of

Fonctionnalisation des barrières thermiques aéronautiques élaborées par barbotines

La sélection des matériaux utilisés dans les moteurs aéronautiques est un enjeu majeur pour assurer la sécurité des passagers, optimiser les performances de l'avion et maîtriser les coûts. Dans les parties les plus chaudes des moteurs (i.e. chambre de combustion et turbine), les pièces sont généralement constituées de superalliages à base nickel en raison de leurs excellentes propriétés mécaniques à haute température. Vulnérables aux phénomènes de corrosion et d'oxydation à haute température, les superalliages doivent la plupart du temps être revêtus afin de prolonger leur durée de vie (ingénierie de surface). La composition chimique et l'architecture des revêtements sont alors adaptées en fonction du régime de température et des phénomènes de dégradation rencontrés (i.e. corrosion à chaud, oxydation et/ou érosion). En vue de répondre aux nouvelles réglementations environnementales, de nouvelles voies de synthèse et de fonctionnalisation sont à l'étude comme alternatives aux procédés industriels actuels. Mots-clés: Barbotines, aluminisation, protection haute température, corrosion à chaud, oxydation.

Functionalization of aeronautical thermal barrier systems elaborated by slurry

The selection of materials is of utmost importance in gas turbine engines to ensure the security of the passengers, optimize the performances of the aircraft and be cost efficient. In the hottest region of the engines (i.e. combustion chamber and turbine), the components are usually made of nickel-based superalloys. These materials can resist to high mechanical loads at high temperature but are vulnerable to aggressive environments. Therefore, nickel-based superalloys are usually coated to increase their durability in the engine (surface engineering). The chemical composition and the coating architecture are carefully adjusted depending on the temperature regime and the mechanisms of degradation encountered (hot corrosion, oxidation and/or erosion). New synthesis routes and functionalization are currently developed as alternative solutions to industrial processes.

As a promising alternative approach, different studies were carried out in the LaSIE laboratory under the European project "PARTICOAT" and confirmed the possibility to elaborate complete thermal barrier systems (diffusion coating + thermal barrier coating) from Al-containing water-based slurries. In this work, the role of Cr as a doping agent was investigated. The addition of Cr decreased the thermodynamic activity of Al upon aluminizing and limited the exothermic reactions usually reported between Al and nickel-based materials. Different architectures of coatings were obtained thanks to the flexibility and the adaptability of the slurry coating process. The gas composition (Ar, air) and the heat treatment conditions were also investigated. Finally, the high temperature resistance of the slurry coatings developed during this work was evaluated under hot corrosion and oxidation conditions.

Keywords: Slurry coatings, aluminizing, high-temperature protection, hot corrosion, oxidation.

  electrochemical corrosion (Fig. I-1a) and oxidation (Fig. I-1b), which bring metals back to their thermodynamically stable oxidised state. Electrochemical corrosion and oxidation processes are governed by ionic and electronic exchanges at the surface of the metal.

Fig. I- 5 .

 5 Fig. I-5. Duplex γ-γ' microstructure of a nickel-based superalloy [3] and corresponding crystal structures.

  . Turbine components manufactured by directional solidification can present a columnar crystal structure (Fig. I-6b) with oriented grain boundaries or a single crystal (Fig. I-6c).

Fig. I- 6 .

 6 Fig. I-6. Cast turbine blades with (a) a polycrystalline structure (EQ), (b) a directionally solidified structure (DS) and (c) a single crystal (SX) [1].

Fig. I- 8 .

 8 Fig. I-8. (a) Solubilities of most common oxides in fused Na2SO4 at 927°C [35] and (b)representation of the negative solubility gradient[35].

Chapter 1 State of the art: context and motivations of the study - 25 -V. 2 .Fig

 1252 Fig. I-10. Binary Al-Ni phase diagram[43].

  . A schematic representation of the different steps for the mechanisms of formation is given in Fig. I-14. The diffusion layer grows very quickly due to the high exothermic reactions occurring between the nickel substrate and the Al microparticles. The reaction synthesis is mainly controlled by molten Al progressively dissolving Ni and forming Al-rich NixAly intermetallic phases at the surface. When the whole Al from the slurry deposit microparticles has reacted, solid state diffusion occurs between the Al-rich coating and the Ni substrate forming a continuous Ni2Al3 layer. As observed for other high-activity diffusion processes (cf. paragraph V.2), an annealing treatment at high temperature (i.e. 1000-1100°C) is necessary to obtain the β-NiAl phase. Simultaneously, the shells of the initial Al microparticles oxidize on the surface of the diffusion layer and sinter together. The resulting ceramic top foam

Fig. I- 14 .Fig. I- 13 .

 1413 Fig. I-14. Mechanisms of formation of the thermal barrier system on a metal substrate proposed by Bouchaud et al. [60].

  Fig. I-15. Graphical abstract of the study.

IV.

  Characterization methods .............................................................................................. IV.1. Calorimetry and thermal analyses .......................................................................................................... a) Differential Scanning Calorimetry (DSC) ....................................................................................... b) Laser Flash Analysis (LFA) .................................................................................................................. IV.2. Crystallographic and structural analyses ............................................................................................... a) X-Ray Diffraction (XRD) .................................................................................................................... b) Raman micro-spectrometry and Cr 3+ fluorescence ..................................................................... IV.3. Microscopic and elemental analyses ...................................................................................................... a) Optical microscopy .............................................................................................................................. b) Scanning Electron Microscopy (SEM) and Energy Dispersive X-ray Spectroscopy (EDS) .. c) Electron Probe Micro-Analysis (EPMA) .......................................................................................... References ............................................................................................................................... Several nickel-based materials were used in the study for the formation of diffusion and thermal barrier coatings by slurry process (Table II-1). Pure Ni and Ni20Cr were used as model materials to study the mechanisms of formation of the coatings before extrapolation to Ni-based superalloys presenting more complex chemical composition. The influence of the microstructure of the alloys was investigated after optimization of the slurry compositions. Directionally solidified superalloys such as CM-247 LC (low carbon) and DS200 are used to produce turbine components of the first LPT stages [1] whereas equiaxed superalloys, such as IN100, are employed for further LPT stages [1].

Fig. II- 1 .

 1 Fig. II-1. XRD patterns of the nickel-based materials investigated in the study.

- 45 -

 45 The metallic powders were analysed by Energy-Dispersive X-ray Spectroscopy (EDS) to detect the potential presence of impurities and determine their oxidation level. The results are summarized in Table II-4. The presence of impurities was not reported with the resolution of the EDS measurements.However, the O content was significantly higher for the Cr powders than for the Al ones. This probably resulted from the peripheral passivation/oxidation of the Cr microparticles (observation of a cellular microstructure on the coarse Cr microparticles in Fig.II-3). The three metallic powders were characterized by X-ray diffraction (XRD) to identify their crystal structures. The XRD patterns given in Fig.II-4were used as references to study the reactivity between Al and Cr microparticles (Chapter III).

Fig. II- 3 .

 3 Fig. II-3. SEM images and size dispersion of the microparticles for the three powdersinvestigated in the study.

Fig. II- 4 .

 4 Fig. II-4. XRD patterns of (a) the Al powder and (b) the two Cr powders investigated in thestudy.

Fig. II- 5 .

 5 Fig. II-5. (a) Picture of the air brush and (b) view of the samples after deposition of the slurry by air brush.

Fig. II- 6 .

 6 Fig. II-6. Architectures of the three different Al-Cr systems investigated in the study.

Fig. II- 7 .

 7 Fig. II-7. Macrographs and SEM surface images (BSE mode) of pure Ni after deposition of the different slurry compositions.

( 52 -

 52 temperature, dwell time and frequency of the cycles). The cycles are performed automatically thanks to pneumatic cylinders moving up and down the ceramic block holding the samples (Fig. II-8b). A cycle is decomposed as follows: -introduction of the samples in the furnace at the oxidation temperature, -hot dwell time at the oxidation temperature, -The cooling of the samples is ensured with four cooling nozzles flushing compressed air during the cold dwell time (Fig. II-8b). With a diameter of 200 mm, the ceramic block allows to place several alumina boats at the same time (Fig. II-8c). The right furnace (Fig. II-8a) working up to 1300°C was used in this study for the cyclic oxidation tests at 1000°C. Standard cycles of 1 hour were performed (45 minutes at 1000°C and 15 minutes of cooling). After 15 minutes of cooling, the temperature of the samples decreases below 80°C. The temperature inside the furnace was continuously measured with a type-S thermocouple (Pt-Rh10/Pt) protected in an alumina sheath (Fig. II-8c). A computer was connected to the furnace to follow the thermal cycles. The mass variation of the samples was regularly measured with a Sartorius Quintix Secura 10 -5 g precision balance to follow the oxidation kinetics.

Fig. II- 8 .III. 2 .

 82 Fig. II-8. Images of the Delta Thermique rig with (a) view of the two furnaces with the panel program on the right side, (b) view of the cooling nozzles and the ceramic block and (c) alumina boats containing the samples and placed around the thermocouple.

Fig. II- 9 .

 9 Fig. II-9. (a) Experimental set up and (b) picture of the horizontal furnace used for the hot corrosion tests in DECHEMA-Forschungsinstitut.

Fig. II- 10 .

 10 Fig. II-10. (a) Schematic of the laser flash technique from Ref. [55] and (b) image of the LFAused in the study.

  Fig. II-11. (a) Fluorescence signals of Cr 3+ in θ-Al2O3 and α-Al2O3 and (b) analysis of an α-Al2O3 powder showing the α-Al2O3 doublet and the corresponding satellite vibrations.

  Reactivity of Al and Cr powders for aluminizing purposes and mechanisms of formation of the slurry coatings on pure nickel Slurry aluminide coatings have been widely used for several years to protect metallic substrates from oxidation and corrosion. Their microstructure is comparable to high aluminium activity coatings because of the high Al content in the donor [1-3]. Such microstructure results from the high exothermal reactions occurring by self-propagating high-temperature synthesis (SHS) [1,4,5]. Montero et al. added pre-alloyed Al-Cr microparticles to the slurry composition and achieved coatings with a

First, the reactivity

  between Al and Cr microparticles was investigated by DSC with no substrate and with pure nickel as a model material for nickel-based superalloys. Then the mechanisms of formation of the coatings from mixtures of Al and Cr powders was studied on pure nickel. Finally, double-layer architectures (Al-Cr and Cr-Al systems) were elaborated to control the mechanisms of formation of the coatings. The main objective of Chapter III was to elucidate the effect of Cr addition in the slurry deposit on the mechanisms of formation of the diffusion coatings. This Chapter focused on pure nickel as a model material to highlight the reactions involved and simplify the interpretation of the results before extrapolation to a Ni-Cr alloy and an industrial nickel-based superalloy in Chapter IV.

1 .

 1 Preparation of Al-Cr mixtures Different Al-Cr compositions were selected from the Al-Cr phase diagram (Fig. III-1) and literature (Table III-1) to investigate the potential exothermic and endothermic reactions. The role of Cr as a moderator of the aluminizing reactions will be thus derived on model materials (Ni, Ni20Cr) and nickelbased superalloys from selected Al-Cr mixtures. Part of this work has been published in Ref. [8]. The Al-Cr phase diagram (Fig. III-1) is relatively complex, with six intermetallic compounds (Al7Cr, Al11Cr2, Al4Cr, Al11Cr4, Al8Cr5 and AlCr2) [9-14]. In the range 30 to 42 at.% Cr, Grushko et al. demonstrated the existence of only one high-temperature and one low-temperature phase [15]. The intermetallic compounds Al9Cr4 and Al8Cr5 therefore present a compositional continuity in this range with a high-temperature and a low-temperature sub-region [16]. Grushko et al. also reported a new

1 )

 1 in an alumina crucible. The heating and cooling rates were fixed at 10°C/min with a maximum temperature of 1300°C. Upon annealing, the reaction chamber was ventilated with a flow of Ar (20 mL/min) to limit oxidation phenomena and prevent secondary reactions in the system. For each composition, energy changes and temperatures of transformation were measured with the Calisto software coupled to the DSC technique.The DSC thermograms for the pure Al and Cr powders are given as references in Fig.III-2. For the pure Al powder (Fig.III-2a), an endothermic reaction was observed at 654 ± 1°C and can be associated with the melting point of Al. This temperature was lower than the theoretical melting point of Al (660.5°C[9]), probably due to impurities in the powder (not detected by EDS)[1,22]. The experimental enthalpy of fusion for the Al powder was 10.5 ± 0.2 kJ.mol -1 , which was close to the value given in the literature (10.7 kJ.mol -1[23]). Upon cooling, the recrystallization of Al was observed in the range 630-650°C. For the two Cr powders (Fig.III-2b), no energy change was observed up to 1300°C, which is far below the melting point of Cr (1863°C[9]).

For

  the 82Al18Cr composition, the endothermic signal of Al melting overlaps the exothermal peak for the two systems (Fig.III-3). This means that Cr was the limiting reagent for this composition[13]. The endothermic signal of Al melting was not observed for the other Al-Cr compositions and this indicates that Al was the limiting reagent this time. After the exothermic reaction, the whole Al reacted with Cr for compositions with 33 wt.% Cr and more. Depending on the initial Al-Cr composition, different intermetallic compounds likely formed. The experimental data collected from the DSC experiments upon heating are given in Table III-2. For each transformation, three temperatures were determined experimentally (Tonset, Tpeak, Toffset).

  -3a and Fig. III-5a). This endothermic transformation was also identified for the 67Al33Cr and 55Al45Cr compositions by mixing Al with the coarse Cr powder (Fig. III-3b and Fig. III-5a). This probably refers to the decomposition of Al11Cr4 phase which was recently found to occur at 829°C [17]. The second (endothermic) transformation in the range 908-929°C should correspond to the peritectic transformation L + Al4Cr ↔ Al11Cr2 [25]. For intermediate Al-Cr compositions (67Al33Cr, 55Al45Cr, 44Al56Cr and also 30Al70Cr by mixing Al with the coarse Cr powder), an endothermic reaction was identified at a temperature of 1035-1037°C. Considering the Al-Cr phase diagram in Fig. III-1 and literature[12,17,25], this endothermic reaction probably refers to the Al4Cr melting (peritectic transformation at 33 wt.% Cr). In agreement with the Al-Cr phase diagram, the highest amount of energy released for this reaction was for the 67Al33Cr composition (Fig.III-5c). For the Al-Cr compositions with high Cr contents (30Al70Cr, 20Al80Cr and 12Al88Cr), the convolution of two endothermic reaction was observed in the temperature range 882-917°C (Fig.III-5b). This could be related to the peritectoid reaction (Cr) ↔ Al8Cr5 + AlCr2 and the congruent reaction (Cr) ↔ AlCr2

III- 6 .

 6 Fig. III-6. SEM images (BSE mode) of the three different Al-Cr mixtures before and after annealing at 650°C for 10 minutes (20 mL/min Ar flow).

  Fig.III-9). For the 82Al18Cr composition, the EDS spot measurements of the synthesized intermetallic phases were in the composition range of the Al7Cr phase (86-87.5 at.% Al[9,16]).

76 -

 76 Naplocha et al. demonstrated that the synthesis of mixed Al and Cr powders proceeds through the peritectic transformations L + Al7Cr → L + Al11Cr2 → L + Al4Cr and that the synthesis began with the formation of Al7Cr[13]. They performed DTA measurements (heating branches) for three different Al-Cr ratios (78Al22Cr, 67Al33Cr and 54Al46Cr in wt.%) and observed that the initial ratio changed the maximum temperature reached during the exothermic reaction. The maximum temperature recorded for each composition was about 777, 877 and 1000°C for respectively, 78Al22Cr, 67Al33Cr and 54Al46Cr. Therefore, the greater the Cr content in the mixture, the higher the local increase in temperature. This suggests that the Cr-rich compositions of the present study were more susceptible to form Cr-rich intermetallic compounds with a high melting point, for instance the Al8Cr5 compound.Considering these results and the Al-Cr binary phase diagram (Fig.III-1), it was then possible to predict the synthesized compounds during the exothermic reaction. For the 67Al33Cr composition, with a maximum temperature estimated at 877°C above the Al7Cr peritectic transformation, the Al11Cr2 phase should be formed. Because of the narrow concentration range between Al11Cr2 and Al4Cr

  Different compositions were tested to separate out the contribution of each reactive. For the tests, planar nickel samples (approximately 1 mm thick) were prepared from a 4.0 mm diameter nickel rod (99.9% purity). For consistency with TGA experiments (cf. Chapter II), the nickel samples were ground with #180 SiC paper then rinsed in distilled water and cleaned with ethanol under ultrasonic agitation before any further step. Three different systems were then considered for the DSC tests (Fig. III-10). The pure nickel substrate was first measured as a reference to identify the intrinsic transformations (Fig. III-11). One can see the reversible Curie point of nickel at 357 ± 2°C [28] corresponding to a change of the magnetic properties of nickel. No other transformation was identified for the nickel sample.

  Fig. III-12. Considering the test with Al powder (without binder, Fig. III-12a), a strong endothermic peak referring to the Al melting was observed at Tonset = 655°C. The exchanged heat of the reaction(5.7 kJ.mol -1 calculated from the initial amount of Al) was lower than the one for the pure Al powder (cf. I.1). This difference suggests that a fraction of the Al powder reacted with the nickel substrate.

  . III-10c), an exothermic reaction was identified at 643 ± 2°C convoluted with the endothermic signal of Al melting (Fig.III-12b).This exothermic reaction probably refers to the formation of Al-rich NixAly intermetallic phases preceded by the melting of NiAl3-Ni(Al) eutectic at 640°C[19,20]. This reaction is characteristic of the high-activity slurry aluminizing of nickel and has been demonstrated to proceed by self-propagation [1]. By comparison with the pure nickel substrate (Fig. III-11), a variation of the heat flow was observed in the range 750-900°C. Unlike for the nickel sample in contact with the Al powder (Fig. III-12a), no melting of NiAl3 nor Ni2Al3 was observed on the DSC curve of Fig. III-12b. This suggested therefore that solid-state diffusion between Ni and Al occurred after the initial exothermic reaction [1]. The samples were analysed in cross-section to observe the diffusion layers after the annealing treatment in the DSC (Fig. III-13). EDS spot measurements were performed to identify the phases formed after the heat treatment [29]. With the Al powder (Fig. III-10b), a thick diffusion zone of 210 ± 10µm thickness was formed on the nickel substrate (Fig. III-13a). This diffusion zone was composed of three different layers as depicted by the contrast difference in the BSE mode of the SEM (Al-rich β-NiAl, Ni-rich β-NiAl and γ'-Ni3Al). With the Al slurry deposit (Fig. III-13b), the diffusion zone was thinner and presented two distinct layers. A mixed outer layer of 60 ± 5µm thickness composed of both Ni-rich β and γ' phases and a continuous inner layer with the γ'-Ni3Al phase (16 ± 3µm). Such differences in the coating microstructures were attributed to different Al reservoirs initially in contact with the nickel substrate.

Fig. III- 12 .

 12 Fig. III-12. DSC thermograms for the pure nickel substrate in contact with (a) the Al metallic powder (without organic binder) and (b) the Al slurry deposit.

4 )

 4 Fig. III-10c. The deposited mass was greater for the Al-Cr slurries than for the pure Al in order to have a sufficient Al content in the donor for diffusion. The DSC thermograms are given in Fig. III-14.

Fig. III- 13 . 81 -

 1381 Fig. III-13. SEM cross-section images (BSE mode) of the pure nickel samples after DSC experiment in contact with (a) the Al metallic powder and (b) the Al slurry deposit.

  -15) of the samples for the three Al-Cr compositions and corresponding surface XRD patterns (Fig. III-16a) demonstrate that the AlxCry intermetallic compounds already formed just after 10 minutes of annealing at 650°C, as also observed by DSC (Fig.

III- 15 84 -

 1584 Fig. III-16. XRD patterns of the surfaces for the 67Al33Cr, 44Al56Cr and 20Al80Cr slurry compositions deposited on pure Ni after annealing at 650°C for (a) 10 minutes and (b) 5 hours (20 mL/min Ar flow).

"IV. 1 . 87 -

 187 low-activity" microstructure. The next part presents the mechanisms of formation of Al-Cr slurry coatings on pure nickel (as a model material) before extrapolation to a Ni-Cr alloy and a nickelbased superalloy (Chapter IV). The role of Cr addition in slurry compositions upon aluminizing was first investigated using Al-Cr mixed slurries. Then, different architectures of coatings were tested by adjusting the composition and the periodicity of the layers to control the mechanisms of formation of the coatings. IV. Mechanisms of formation of the diffusion layers on pure nickel Aluminizing of pure nickel with different Al-Cr mixed slurries: influence of Cr addition on the Al activity In order to study the role of Cr addition upon aluminizing, three Al-Cr slurry coatings (67Al33Cr, 44Al56Cr and 20Al80Cr) were deposited on pure Ni and then submitted to a diffusion heat treatment (650°C-5h). The simple Al slurry coating (with an Al activity equal to 1) is given as a reference. The SEM cross-section images and the concentration profiles of Al and Cr for the four coating systems after the diffusion heat treatment are given in Fig. III-18. For the simple Al coating (Fig. III-18a), three different layers were observed. The outer phase, with a thickness of approximately 35µm, consisted in the Ni2Al3 phase with a high Al content (approximately 60 at.%). The two other layers were thinner and contained less Al with respectively -NiAl (intermediate contrast) and γ'-Ni3Al phases (bright layer). Both -NiAl and γ'-Ni3Al layers had a thickness of approximately 3-5µm. This microstructure was typical of a high activity slurry coating [4,22]. According to the EDS analyses, Al diffused in the Ni substrate with up to 6 at.% at 10µm below the γ'-Ni3Al-substrate interface. For the three Al-Cr compositions, the diffusion of Al into the Ni substrate was considerably hampered by the addition of Cr. None of the diffusion layers that formed on the surface exceeded 10 microns in thickness (Fig. III-18). According to the EDS analyses, -NiAl and γ'-Ni3Al phases were identified. However, the diffusion layers were not continuous for the three compositions and this might be explained because either Al or Cr were in direct contact with the Ni substrate after deposition of the slurry. Kirkendall porosities were observed between the Ni substrate and the γ'-Ni3Al phase suggesting that the outward diffusion of Ni from the substrate to react with Al and Al-rich intermetallics was not compensated by Al inward diffusion. As such, approximately 3 at.% Ni incorporated to the Al-Cr intermetallic compounds above the substrate. Conversely, an increasing amount of Cr was dissolved in the coating and diffused into the Ni substrate with the greater Cr content in the mixtures (Fig. III-18f). -Upon aluminizing at 650°C, concurrent reactions occurred at the surface of the nickel substrate between Al, Cr and Ni. Al preferentially reacted with Cr within the slurry deposit to result in the formation of AlxCry intermetallics above the Ni substrate. These intermetallic compounds hampered the inward diffusion of Al into the Ni substrate. Therefore, the SHS reactions between Ni and Al usually reported for high-activity slurry coatings seemed to be limited [1,4,5]. This would explain why the diffusion layers were thin and discontinuous for all three Al-Cr compositions after 5 hours at 650°C (Fig. III-18).

Fig. III- 18 . 88 -IV. 2 .

 18882 Fig. III-18. BSE cross-sections images for (a) simple Al, (b) 67Al33Cr, (c) 44Al56Cr and (d) 20Al80Cr slurry coatings on pure nickel after the diffusion heat treatment under Ar atmosphere and corresponding (e) Al and (f) Cr EDS concentration profiles.

90 -

 90 20. A 50µm thick coating formed from the 67Al33Cr composition. The outer layer was composed of the -NiAl phase and the inner layer presented a composition close to the γ'-Ni3Al phase. The Cr content of the diffusion zone remained low (less than 2 at.%) because of the limited solubility of Cr in the -NiAl phase[36][37][38]. Some Cr-rich precipitates were indeed observed in the external region of the coating (Fig.III-20a). For the 44Al56Cr composition, the outward diffusion of nickel was clearly observed with the coalescence of Kirkendall porosities at the initial surface of the substrate (Fig.III-20b). Above the nickel substrate, a heterogeneous diffusion zone was formed with a composition close to the -NiAl phase. The precipitation of α-Cr occurred within the -NiAl phase and suggested that this diffusion zone was formed by the dissolution of AlxCry intermetallic phases due to the outward diffusion of nickel. The Al reservoir of the slurry deposit was too low for this composition to grow a continuous -NiAl layer.Chapter III -For the 20Al80Cr composition (Fig.III-20c), some islands with the β-NiAl composition were observed above the nickel substrate. The Al content of the slurry deposit was too low in this composition to grow a continuous -NiAl layer. In the slurry deposit, the initial Al-Cr phases were enriched in nickel.

91 -V.

 91 Mechanisms of formation of the diffusion layers from Al-Cr and Cr-Al double-layer systems V.1. Elaboration of Al-Cr and Cr-Al double-layer systems Double-layer systems have been studied as an alternative to Al-Cr mixed slurries developed in paragraph IV. Thanks to the flexibility of the slurry coatings, it is possible to adjust the architecture of the coatings to optimize their properties ("coatings by design"). Two double-layer systems have been elaborated (Fig. III-21) to investigate the mechanisms of formation of the coatings on pure nickel substrate. After the application of the first layer, the samples were let to dry in a fume cupboard for one hour before the application of the second layer [22]. The deposited mass of Cr was fixed at 10 mg.cm -2 for the two systems to study the influence of the Al/Cr ratio by varying the Al content (5 or 10 mg.cm -2 ), see Fig. III-21. For the sake of clarity, the Al/Cr ratio of the slurry deposits (in wt.%) was considered to differentiate the different architectures tested in this work. For the Al-Cr double-layer

  Fig. III-21. Architecture of the Al-Cr and Cr-Al double-layer systems elaborated on pure nickel.

  650°C was also prolonged up to 20 hours to emphasize the influence of the dwell time on the composition and the thickness of the diffusion layers (Fig.. With the low Al reservoir (33Al-67Cr system), only two distinct layers were observed with a composition close to β-NiAl and γ'-Ni3Al according to the EDS spot measurements. The overall thickness of the diffusion zone was not substantially changed after 20 hours. For the high Al reservoir (50Al-50Cr system), the diffusion zone presented an outer heterogeneous layer after 5 hours with γ-(Ni), β-NiAl and Ni2Al3 phases identified.The presence of γ-(Ni) in the outer region of the diffusion zone might come from the dissolution of nickel by molten Al during the SHS reactions[39]. After 20 hours of annealing, the diffusion zone was more homogeneous in composition probably due to solid state diffusion between Al and Ni. Some remnants of the Ni2Al3 phase were observed at the extreme surface of the diffusion zone whereas some grains with the γ'-Ni3Al composition were identified through the β-NiAl layer. The overall thickness of the diffusion zone was close for both Al reservoirs.

  . III-22), the unreacted Cr microparticles were oxidized after annealing. The peripheral oxidation of the Cr microparticles (Fig. III-25) probably resulted from the low partial pressure of O2 and H2O in the reaction chamber (2 vpm O2 and 5 vpm H2O). The formation of chromia was confirmed with the Raman surface analyses of the top coats showing the five characteristic peaks of chromia [40-42]. Note the progressive oxidation of the Cr microparticles on the SEM cross-section images (Fig. III-25a and Fig. III-25b).

Fig. III- 25 .

 25 Fig. III-25. BSE cross-section images showing the peripheral oxidation of Cr microparticles after (a) 650°C-5h and (b) 650°C-20h in Ar and (c), (d) corresponding surface Raman analyses.

Fig. III- 27 .V. 3 .

 273 Fig. III-27. Sequence proposed for the mechanisms of formation of the Al-Cr double-layersystem elaborated on pure nickel.

  maps of the sample are given on Fig. III-29. All the Al microparticles reacted with the Cr ones after the diffusion heat treatment resulting in a top layer with hollow alumina spheres [22] and an intermediate layer with Al-rich AlxCry intermetallic phases. A layer of Cr microparticles was still identified above the nickel substrate suggesting that Al was the limiting reagent for this composition (50Cr-50Al system). This result was relevant with the conclusions drawn from Fig. III-6. An oxygen enrichment was observed in the Cr-rich layer and probably refers to the peripheral oxidation of the Cr microparticles (Fig. III-25).

Fig. III- 28 .- 98 -

 2898 Fig. III-28. DSC thermogram for the Cr-Al double-layer system elaborated on pure nickelsubstrate.

  This indicates the low solubility of Cr in the phases formed in the diffusion zone. According to the EDS spot measurements (Fig. III-33) and the XRD patterns (Fig. III-31), the Al-rich AlxCry intermetallics formed at 650°C transformed into the stable Al8Cr5 phase with the additional annealing period at 1000°C for 3 hours. It thus seems that Al8Cr5 released Al to grow a diffusion zone by interdiffusion with nickel. The alumina hollow spheres were still observed on the surface and converted into the stable α-Al2O3 after 3 hours of annealing at 1000°C [22] (Fig. III-31). At higher magnification (Fig. III-33), three distinct layers were observed in the diffusion zone for the 50Cr-50Al composition. The outer dark layer in BSE mode was identified as the Al-rich -NiAl phase. The intermediate layer was composed of the Ni-rich β-NiAl phase with dissolved Cr (approximately 5 at.%). The third layer was very thin and located just above the nickel substrate with a composition close to the γ'-Ni3Al phase. The transformation of AlxCry into NixAly phases by reaction with Ni results

  33Al system was too low to grow a diffusion zone (Fig. III-33). In other words, the Al activity of the slurry deposit was too low to grow a continuous -NiAl layer. The phases observed at the interface between the nickel substrate and the deposit were indeed rich in Cr (Fig. III-33) suggesting the α-(Cr) solid solution.

103 -

 103 Fig. III-34. When the temperature reached approximately 640°C (Fig. III-28), the exothermic reaction between Al and Cr microparticles started in the slurry deposit. During this exothermic reaction, different Al-rich Al-Cr intermetallic compounds were formed at the interface between Cr and Al layers. After annealing for 5 hours at 650°C (Fig. III-29), the whole Al had reacted with the Cr microparticles to form these intermetallic phases (Fig. III-31). However, a layer of unreacted Cr microparticles was still observed above the nickel substrate. The 2 of the reaction chamber was sufficient to oxidize the surface of the Cr microparticles and the hollow alumina spheres.In the presence of a Cr interlayer, the SHS reactions usually observed between the substrate and Al microparticles were thus completely supressed with this architecture. Increasing the temperature to 1000°C resulted in the formation of a diffusion zone of approximately 25µm of thickness (Fig. with the 50Cr-50Al system. The presence of undissolved Cr particles within the growing diffusion zone confirmed that the outward diffusion of nickel was governing the formation of the diffusion zone. The Al-Cr intermetallic phases were progressively dissolved and became richer in Cr with time (Fig.III-31and Fig.III-32). In the ceramic top coat, the oxidation of the alumina shells was fostered at 1000°C and converted the transition aluminas into the stable α-Al2O3 phase[22,45].-The Cr-Al double-layer system was found to successfully decrease the activity of Al on the surface of the substrate upon aluminizing. The diffusion of Al at low temperature (650°C) into the nickel substrate was hampered by the layer of Cr microparticles through the formation of Al-Cr intermetallic compounds in the slurry deposit. This architecture avoided the SHS reactions usually observed for high-activity slurry aluminizing of nickel substrates[1,4,5]. With the additional annealing at 1000°C for 3 hours, interdiffusion between the nickel substrate and the AlxCry intermetallic phases synthesized in the slurry deposit was promoted. Therefore, a diffusion zone composed of the -NiAl phase grew by outward diffusion of nickel.

Fig. III- 34 .

 34 Fig. III-34. Sequence proposed for the mechanisms of formation of the Cr-Al double-layersystem elaborated on pure nickel.

  45-50µm (Fig. III-35b). The precipitation of α-Cr occurred in the outer region of the coating whereas Cr-rich phases remained above the diffusion zone (Fig. III-35d). In both microstructures, a γ'-Ni3Al network was observed within the -NiAl phase (Fig. III-35a and Fig. III-35b). This probably referred to the grain boundaries of the coatings. Formation of γ'-Ni3Al within the grain boundaries would also confirm the hypothesis of predominant outward diffusion of nickel upon aluminizing.



  Fig. III-35. (a), (b) BSE cross-section images and X-ray maps for the Cr-Al double-layer systems elaborated on pure nickel and (c), (d) corresponding EDS concentration profiles after the conventional low-activity heat treatment in Ar.

4 )-

 4 through the formation of AlxCry phases above the substrate. Thanks to the flexibility of the slurry technique, different architectures of coatings were elaborated to elucidate the mechanisms of formation of the coatings on pure nickel. Two double-layer systems (Al-Cr and Cr-Al double-layer systems) were thus investigated by alternative deposition of Al and Cr microparticles. For the Al-Cr double-layer system, the Cr top layer was found to deplete the Al reservoir by forming AlxCry intermetallics. As a result, the diffusion of Al towards the substrate was limited at 650°C and the formation of Al-rich NixAly intermetallic phases (namely NiAl3 and Ni2Al3) was considerably hampered on the substrate surface. The additional annealing at high temperature (i.e. 1000°C-3h) did not lead to further Al ingress in the diffusion zone since the layer of hollow alumina spheres acted as a barrier to diffusion. This layer, composed of thin alumina shells, was found to be brittle and not adherent to the underlying substrate.For the Cr-Al double-layer system, the Cr interlayer acted as a diffusion barrier for Al. For the two Cr-Al ratios investigated (approximately 50Cr-50Al and 67Cr-33Al in wt.%), unreacted Cr microparticles were observed above the nickel substrate after 5 hours of annealing at 650°C. With this architecture, the SHS reactions between Al microparticles and the nickel substrate were thus completely suppressed. Heat treatment at high temperature (i.e. 1000°C or 1080°C) resulted in the formation of a continuous diffusion zone with mainly the -NiAl composition. These results suggested that the coatings were formed by major outward diffusion of nickel from the substrate. The outward motion of nickel progressively dissolved the AlxCry phases synthesized in the slurry deposit and foster the direct formation of Ni-rich -NiAl. Since Cr has a low solubility in the -NiAl phase, undissolved Cr was observed in the form of Cr-rich phases in the diffusion zone. The sequence of reaction can be described following equations (Eq. III-3) and (Eq. III-4):xAl + yCr → AlxCry (Eq. III-3) Ni + AlxCry → -NiAl + α-(Cr) (Eq. III-The initial Cr-Al ratio of the slurry deposit therefore played an important role on the composition and the homogeneity of the diffusion layers formed on nickel substrate. The conventional "lowactivity" heat treatment (400°C-3h + 1080°C-6h) was particularly effective to form homogeneous diffusion layers with the desired microstructure. Therefore, this heat treatment was selected for the elaboration of Cr-Al slurry coatings on a Ni-Cr alloy and CM-247 LC superalloy (Chapter IV). Slurry aluminizing of nickel-based alloys from Al and Cr microparticles I. Slurry aluminizing of polycrystalline Ni20Cr .............................................................. I.1. From Al microparticles ...........................................................................................................................

  . IV-1. In Fig.IV-1a, the raw substrate was clearly identified (bright areas) together with the random distribution of diffusion spots (dark areas). In the cross-section (Fig.IV-1c), the island-like diffusion morphology reported by B. Rannou[12] was observed. After the complete heat treatment (Fig.IV-1d), the aluminized layer appeared heterogeneous and exhibited diffusion islands. In addition, the top coat was not adherent regardless of the annealing treatment.At higher magnification, the elongated morphology of the diffusion islands (Fig.IV-2a) suggests that Cr acted as a barrier of diffusion preventing further Al ingress into the substrate. However, after the complete heat treatment, the aluminized layer was discontinuous and Cr-rich phases precipitated at the interface of the coating with the substrate (Fig.IV-2b). The composition of these phases was in the α-(Cr) domain[19] according to EDS spot measurements (approximately 3 at.% Al, 76 at.% Cr and 21 at.% Ni). The Cr concentration of β-NiAl and γ'-Ni3Al phases of the aluminized layer did not exceed 5 at.%.

  performed on representative areas to get the elemental distribution after annealing (Fig. IV-5). For both 44Al56Cr (Fig. IV-5a) and 20Al80Cr (Fig. IV-5b) slurry mixtures, a nickel-depleted area was clearly identified at the interface between the aluminized layer and the substrate, where the voids are observed. A Cr enrichment was also observed in the nickel-depleted area and suggested that the outward diffusion of Ni was faster than that of Cr at 1080°C [28]. As a result, the outward diffusion of nickel was clearly observed above the continuous aluminized layer. The EDS concentration profiles of the cross-section of Ni20Cr substrate slurry coated with the 44Al56Cr and the 20Al80Cr compositions are given in Fig. IV-6a and Fig. IV-6b. Note that the interface between the substrate and the coating was taken as the origin to perform the EDS spot analyses. As observed with the pure nickel substrate, nickel dissolved the Al-Cr intermetallic phases resulting in the direct formation of the β-NiAl phase (Fig. IV-6). Nickel indeed presents a low solubility in Al4Cr and

Fig. IV- 6 .I. 3 .

 63 Fig. IV-6. EDS concentration profiles for Ni20Cr substrate slurry coated with (a) 44Al56Cr and (b) 20Al80Cr mixtures after annealing (400°C-3h + 1080°C-6h).

  Alternatively, they can correspond to oxidized Al microparticles but the resolution of the X-ray map of O does not allow to depict the periphery. Unlike the X-ray maps obtained for Al-Cr mixed slurries in Fig. IV-5, the ones of Ni and Al in Fig. IV-8 with the Cr-Al doublelayer attested the formation of an even β-NiAl layer of about 45µm thickness.

  Fig. IV-8. BSE cross-section image and corresponding EDS maps of (a) Ni, (b) Al, (c) Cr and (d)O. The map of Si is not presented since this element did not segregate in the coating.

Fig. IV- 9 .

 9 Fig. IV-9. (a) Low magnification BSE image of a corner and (b) details of the interdiffusion zone (electrolytically etched with 10 vol.% H2C2O4) after elaboration of the 45Cr-55Al double-layer system on Ni20Cr substrate.

II. 1 .

 1 Fig. IV-10. Micrographs of the Cr-Al double-layer system elaborated on Ni20Cr substrate after complete heat treatment (400°C-1h + 700°C-2h + 1100°C-2h) with (a), (c) 50Cr-50Al (wt.%) and (b), (d) 40Cr-60Al (wt.%) initial compositions.

For

  Fig. IV-12. As observed with the Ni20Cr substrate (Fig. IV-3), the thickness of the diffusion zone after annealing increased with increasing the Al content of the slurry mixture.

Fig. IV- 12 .

 12 Fig. IV-12. Coating thickness measurements (diffusion zone) after annealing of the three Al-Cr mixed slurries on CM-247 LC superalloy.

  Because of the random distribution of the Al microparticles on the substrate surface after slurry deposition (Fig. II.7. of Chapter II), this reaction was only possible on local spots where Al was in direct contact with the substrate. The initial phases formed upon aluminizing where thus different from one area to another, which modified the motion of Al and Ni. The initial formation of Ni2Al3 would facilitate Al motion whereas the direct formation of Ni-rich β-NiAl phase would promote the nickel outward diffusion upon aluminizing[45,46]. This would explain the rough surface of the coating after annealing.For the 20Al80Cr composition, the overall thickness of the aluminized layer did not exceed 20µm and was mainly composed by the interdiffusion zone (Fig.IV-14e). As depicted with the EDS concentration profile on Fig.IV-14f, coarse precipitates with different compositions were formed in a β-NiAl matrix.The nickel enrichment of the phases formed above the aluminized layer suggested that nickel outward diffusion occurred upon aluminizing. However, the Al content of the 20Al80Cr slurry was too low to grow a thicker diffusion zone. These results are in good agreement with the observations done on pure nickel substrate (paragraph IV.2. of Chapter II) and on Ni20Cr substrate (paragraph I.2. of Chapter III).

  For the three compositions, Crrich phases present in the top coat were easily identified and contained great amounts of Al. For the 20Al80Cr slurry mixture (Fig. IV-15c), Cr and Ta (W) coarse precipitates were identified through the β-NiAl coating matrix. Above the aluminized layer, both β-NiAl and α-(Cr) phases (Fig. IV-14f) were distinctively observed. This confirmed the preferential formation of NixAly phases over NixCry phases upon nickel outward diffusion from the substrate into the AlxCry phases present above the diffusion zone (cf. Fig. III-20 in Chapter III).

131 -II. 2 .

 1312 Low-activity high-temperature (LAHT) coatings from the Cr-Al double-layer system a) Low-activity conventional heat treatment (400°C-3h + 1080°C-6h) The aluminizing tests on CM-247 LC from Al-Cr mixed slurries demonstrated that the 44Al56Cr composition was suitable to limit the inward diffusion of Al at low temperature and to foster the outward diffusion of nickel at high temperature. The Al activity of this composition seemed to be the threshold to obtain low Al activity coatings from elemental powders at 1080°C. Therefore, the 44Al56Cr composition was selected to elaborate a low Al activity coating on CM-247 LC superalloy based on the Cr-Al double-layer system (cf. Fig. II-6. In Chapter II). For this, a first layer of approximately 10 mg.cm -2 of Cr and a second layer of approximately 8 mg.cm -2 of Al were deposited on the superalloy by air brush (close to the 56Cr-44Al ratio in wt.%). The sample was subsequently heat treated in flowing Ar at 400°C for 3 hours and 1080°C for 6 hours. A macrograph of the sample after annealing is given on Fig. IV-17a. The alumina top coat structure was not adherent and the surface micrograph suggested the presence of intermetallic phases formed in the slurry deposit upon annealing (Fig. IV-17b). A cross-section micrograph of the coating after annealing is given on Fig. IV-18a.

  zone (IDZ), an intermediate layer with fine dispersion of Cr and an outer layer with Cr-rich phases in the coating matrix (Fig.IV-18a). The thicknesses of both additive layer (A.L.) and IDZ were very close over the whole sample with an IDZ/A.L. ratio close to one[37]. The final microstructure of the coating was very close to the Cr-modified aluminide coating obtained by Gleeson et al. from pack co-deposition process (Fig.IV-18b)[32]. The Cr enrichment of the additive layer was confined to approximately 20% of the total coating thickness, which is also in good agreement with the results obtained by Gleeson et al.[32]. The substrate carbides highlighted on Fig.IV-18awere entrapped in the interdiffusion zone after annealing. This suggested the predominant outward diffusion of nickel upon aluminizing[40,41].

Fig. IV- 18 .

 18 Fig. IV-18. Cross-section optical micrograph of the 56Cr-44Al double-layer system elaborated on CM-247 LC superalloy after annealing in flowing Ar (400°C-3h + 1080°C-6h) and (b) microstructure of the Cr-modified aluminide coating produced by Gleeson et al. [32].

  . IV-20. Note that the EDS spot measurements in the outer region of the coating were realized in the -NiAl coating matrix between the α-Cr precipitates. The determination of the exact composition of the α-Cr precipitates was not possible by EDS because of the relatively large spot resolution. Ta and W enrichments were confined in the IDZ (Fig.. The larger atomic sizes of Ta and W and their very low solubility limits into -NiAl can probably explain their low concentrations in the additive layer[49,50,51]. This observation is in favour of outward grown coating where Cr and Co are the faster diffusing elements with Ni and Al[40,41,49].

  Fig. IV-18a, the substrate carbides were confined in the IDZ after annealing suggesting that the growth of the coating predominantly occurred by outward diffusion of nickel at high temperature. By comparison with the sample heat treated with the conventional low-activity heat treatment (Fig. IV-18a), the additive layer on Fig. IV-22 presented a darker contrast under light optical microscopy. This darker contrast suggested the formation of the Al-rich β-NiAl phase as observed on the Ni20Cr substrate (Fig. IV-10c and d). The distribution of bright phases under light microscopy within the additive layer (Fig. IV-22) suggested that the Cr enrichment was greater than the one observed after the conventional low-activity heat treatment (Fig. IV-18a). These bright phases were observed over the whole thickness of the additive layer, which means up to approximately 50% of the total coating

  Fig. IV-18a and Fig. IV-22 suggested that the growth predominantly occurred by outward diffusion of nickel at high temperature [40]. As previously described by Goward et al., outwardly grown coatings foster the direct formation of the Ni-rich β-NiAl compound on the substrate surface[1,40]. Since Ni diffusivity is more significant than that of Al in Ni-rich β-NiAl phase at high temperature[45,46], Ni outward diffusion from the substrate is promoted and an additive layer based on the Ni-rich β-NiAl phase is formed above the initial surface of the substrate by reaction with Al[1,40]. Because of nickel withdrawal from the underlying substrate, an interdiffusion zone is formed where substrate elements having low solubility in β-NiAl precipitate (Ta and W typically). The inert carbides originally present in the substrate could also be observed in the coating microstructure but remained in the interdiffusion zone[40,41]. The β-NiAl additive layer was therefore "cleaner" than inward-grown coatings. This therefore suggests that the composition of the additive layer should be in the Ni-rich β-NiAl domain after annealing as observed on the concentration profile of Fig.

1 ) - 139 -

 1139 Fig. III-34 in Chapter III). Since nickel has low solubility in these Al-Cr phases (about 1 at.% in Al4Cr and 3 at.% in Al8Cr5)[16], and because the enthalpy of formation of NixAly phases is more exothermic than the one of NixCry phases, nickel dissolved the Al-Cr phases and preferentially combined with Al. Consequently, the outward diffusion of nickel fostered the direct formation of Nirich β-NiAl with a small fraction of dissolved Cr. With a low solubility in the β-NiAl phase (less than 8 at.%[2,13,55]), a large fraction of Cr remained undissolved in the additive layer. Therefore, the growth of the coating with the 56Cr-44Al double-layer system seemed to be governed by the simplified reaction (Eq. IV-1):Ni + AlxCry → -NiAl + α-(Cr) (Eq. IV-The continuous outward diffusion of nickel upon aluminizing progressively consumed the AlxCry phases (i.e. by combination with Al). Therefore, reaction (Eq. IV-1) occurred until the AlxCry donors were completely dissolved. The SEM characterization of the samples after annealing also showed that a fraction of Al was still present in the slurry deposit and did not incorporate in the diffusion coating (Fig.. This external Al content was mainly coming from AlCr2 and α-(Cr) phases after the aluminizing process (Fig.IV-13). Since the AlCr2 phase is not stable at temperatures higher than 910°C (Fig. III-5b in Chapter III), this phase was not present in the microstructure at the aluminizing temperature (1080°C or 1100°C) and was formed upon cooling[56][57][58]. Considering the Ni-Al-Cr ternary equilibrium[19] and the observations of Huang et al.[10], only β-NiAl and α-(Cr) phases were stable at the aluminizing temperature. Therefore, the growth of the coating should occur until the AlxCry donor of reaction (Eq. IV-1) is converted into α-(Cr). As soon as α-(Cr) constitutes the main fraction of the slurry deposit (i.e. the Al donor), the kinetics of aluminizing should decrease because of the very low driving force of the new simplified reaction (Eq. IV-2): Ni + α-Cr(Al → -NiAl + α-(Cr) (Eq. IV-2)

  Fig. III-15) and Cr-Al double-layer systems (cf. Fig. III-31). It can therefore be estimated that the 2 hours annealing at 700°C

  . III-15), both slurry mixtures were found to mainly form the Al8Cr5 phase with an excess of Cr (most probably α-(Cr) solid solution). After further annealing at 1000°C for 3 hours (Fig., Al8Cr5 was still identified for the 44Al56Cr mixture whereas the 20Al80Cr one was mainly composed of α-(Cr) at temperature. This therefore implies that the driving forces of reactions (Eq. IV-1) and (Eq. IV-2) were significantly different between the two compositions. For the 20Al80Cr mixed slurry, the major presence of α-(Cr) phases as Al donors at the initial stage of aluminizing considerably limited the growth of the coating. This was well illustrated in Fig. IV-12 and Fig. IV-15c with the very thin aluminized layers formed on the CM-247-LC superalloy. For the 44Al56Cr mixed slurry, the initial Al content was low enough to ensure the major outward diffusion of nickel and high enough to grow an aluminized layer of significant thickness (Fig. IV-12 and Fig. IV-15c).

  IV-25a. Some cracks highlighted by white circles, were observed in the top foam structure composed of hollow alumina microspheres (Fig.IV-25b)[12,18,27,44,47,54] while the diffusion area looked quite homogeneous in the optical cross-section (Fig.IV-26a).

  Fig. IV-26b [41]. The different layers identified by Das et al. after such coating formation are reported on Fig. IV-26a. The coating can be divided in three distinct layers: an outer layer rich in precipitates predominantly formed by inward diffusion of Al (INL + ONLS), an intermediate layer free of precipitates formed by outward diffusion of nickel (ONLO) and an interdiffusion layer (IDZ).

  . IV-27). The bright phases observed in the outer region of the diffusion layer (D.L.) were enriched in Ta (Fig. IV-27d) and Hf (cf. Annex 1) and confirmed the presence of the substrate carbides (SC). This attested that the initial surface of the substrate remained at the extreme surface of the coating upon aluminizing (cf. Annex 1 and Ref. [41]). The X-ray map of Cr emphasized the fine precipitation of α-(Cr) in the outer region of the diffusion layer (Fig. IV-27c). EDS spot measurements were also performed in cross-section to get the concentration profile of the coating layer (Fig. IV-28a). The outer region of the diffusion layer was mainly composed of Al-rich -NiAl phase with local Cr enrichments, as suggested from the optical micrograph. Note that the concentrations of W and Ta in the diffusion layer were not negligible here in comparison with the outward grown coatings of paragraph II.2. The detailed view of the outer region of the coating in BSE mode (Fig. IV-28b) emphasized the fine distribution of Cr-rich phases with an intermediate contrast.

  treatment (400°C-1h + 700°C-2h + 1100°C-2h) tailored for the formation of the full thermal barrier system on nickel-based superalloys[12,27,44,47]. The surface macrograph attested the great homogeneity of the sample and the adherence of the top coat structure (Fig.IV-29a). The characteristic top coat structure composed of hollow alumina microspheres was observed on the sample surface by optical microscopy (Fig.IV-29b)[12,27,44,47].

  146-A low magnification micrograph of the cross-section is given on Fig. IV-30. The overall thickness of the aluminized layer was of approximately 80µm, which is twice thicker than the outward grown coatings of Fig. IV-22. The diffusion layer, highlighted here with a dark contrast, mostly composed the coating whereas the IDZ at the interface additive layer-substrate was only 18±2µm thick. These observations were in good agreement with the microstructures obtained by Das et al. from single-step high Al activity packs after 2 hours of aluminizing at 1034°C (cf. Fig. IV-26b). Unlike the 56Cr-44Al doublelayer system (Fig. IV-22a), substrate carbides were observed in the outer region of the coating. This confirmed the predominant inward diffusion of Al at the early stages of the aluminizing process [41].The reaction front between the IDZ and the underlying substrate was not even for the 33Cr-67Al double-layer system compared to the 44Cr-56Al one. This could also be related to the initial inward growth of the coating and the delayed outward diffusion of nickel. The schematic representation of the coating formation proposed by Das et al. seems to very well fit with the present study (Fig. IV-26b).

  microstructure of the coatings (Fig. IV-32) as observed in Chapter III for pure nickel substrate. Such differences resulted from the different mechanisms of formation of the coatings and especially from the predominant outward diffusion of Ni or inward diffusion of Al upon aluminizing [1,32,40,41,55].

  Fig. IV-32a and Fig. IV-32c), the coating was formed by predominant outward diffusion of nickel at high temperature. The Ni-rich β-NiAl phase was readily formed on the substrate surface with an outer additive layer free of precipitates above the initial substrate surface (ISS), cf. Annex 1. Because of the design of the present system, undissolved α-(Cr) phases were entrapped in the additive layer. The growth of the coating with the 56Cr-44Al double-layer system can therefore be simplified by the reaction (Eq. IV-3):

Fig

  Fig. IV-32c is most likely to form upon cooling[41,45,46]. This layer was not observed for the sample heat treated with the conventional heat treatment. After 6 hours of aluminizing at 1080°C, the slurry deposit should be composed of α-(Cr) phases thus considerably limiting the Al activity. Upon cooling, the Al activity was then not high enough to foster the formation of an Al-rich β-NiAl layer by inward diffusion of Al. Even though the aluminizing temperature was higher for the complete heat treatment (1100°C), the shorter time of diffusion (2 hours) did not allow the full decomposition of the AlxCry phases (confirmed with EDS measurements). The Al activity of the slurry deposit was therefore high enough upon cooling to cause inward diffusion of Al into the β-NiAl coating matrix. This Al enrichment fostered the formation of Al-rich β-NiAl phase, composing nearly the whole additive layer after annealing (Fig.IV-32c) because of the greater diffusivity of Al than of Ni in such Al-rich β-NiAl phase[45,46].

  Fig. IV-34. (a) BSE cross-section image and (b) EDS concentration profile of CM-247 LC slurry coated with the 33Cr-67Al double-layer system after annealing in flowing Ar (400°C-3h + 1000°C-6h).

  Fig. III-15 in Chapter III). The high Al activity of this phase seemed to be high enough to foster the inward diffusion of Al into the substrate. The evidence of inward diffusion of Al after annealing at 1000°C confirmed that the reaction Al4Cr → L + Al8Cr5 was not responsible of the coating formation. The Al-Cr phases synthesized in the slurry deposit therefore released the Al diffusing inwardly into the underlying substrate. Consequently, the composition of the intermetallic phases changed from Al4Cr to Al8Cr5 and AlCr2 after 6 hours annealing at 1000°C. This change in composition attested the progressive Al depletion of the Al-Cr donors to diffuse in the substrate. Since the Ni2Al3 phase has a very low domain of stability at 1000°C [16], Al-rich β-NiAl formed that expelled Cr out. Further Al ingress combined with the segregated Cr and led to a fine precipitation of Cr-rich phases (presumably Al8Cr5 [10]) in the coating matrix (Fig. IV-35a).

  Fig. IV-36a and Fig. IV-36b suggested that the aluminizing temperature had a significant influence on the resulting microstructure of the coating. The interdiffusion zone (IDZ) was much thicker for the sample annealed at 1080°C (Fig.

  Influence of the atmosphere (Ar, air) on the formation of thermal barrier coatings (TBC) from simple Al and Cr-Al double-layer slurry systems I. General background on the oxidation of Al particles and the formation of hollow alumina spheres .................................................................................................................... II. Influence of the atmosphere (Ar, air) on the aluminizing of pure nickel from Al microparticles ....................................................................................................................... II.1. Diffusion heat treatment (400°C-1h + 700°C-2h) ........................................................................... II.2. Complete heat treatment (400°C-1h + 700°C-2h + 1100°C-2h) ................................................ II.3. Complete heat treatment with the introduction of synthetic air upon annealing .................... II.4. Discussion .................................................................................................................................................. III. Influence of the atmosphere (Ar, air) for the Cr-Al double-layer system ............... III.1. Characterization of the top coat structure after heat treatment in Ar ...................................... III.2. Influence of the atmosphere (Ar, air) on the formation of the Cr-Al double-layer system on pure nickel ............................................................................................................................................................. III.3. Reactivity between Al microparticles and oxide powders (NiO, Cr2O3 and Al2O3) .............. III.4. Discussion .................................................................................................................................................. IV. Summary and outlook .................................................................................................. References ............................................................................................................................. background on the oxidation of Al particles and the formation of hollow alumina spheres Recently, Pedraza et al. investigated the influence of the annealing conditions on the formation of hollow alumina microspheres from Al microparticles [1]. The composition of the atmosphere was found to play a significant role on the release of Al from the microparticles and their simultaneous oxidation as reported by other authors [2-5]. The release of Al from the particles was associated with concomitant mechanisms including the peripheral oxidation of the spheres and a volume variation due to the expansion of the Al core upon heating [1,6]. Depending on the atmosphere (inert or oxidizing), the thickness of the oxide shells and the amount of Al flowing out from the microspheres were different. These concomitant mechanisms were first proposed by Levitas et al. for nano-sized and micro-sized Al particles [7-10] and are presented on Fig. V-1. For micro-sized particles, the initial amorphous alumina shell grows until it reaches a critical thickness of approximately 4 to 5 nm [3,11,12].

  Fig. V-2 summarizes the different stages of oxidation of Al microparticles

  The surface morphology of the top coat structure after the diffusion heat treatment in Ar and in synthetic air is given in Fig.V-4. With Ar atmosphere (Fig.V-4a), hollow spheres can be observed on the surface with a morphology comparable to the one presented by Pedraza et al. in 120 Pa He+4%H2[1]. The Al and O contents of the surface were very close according to the EDS area analyses of the surface (Fig.V-4c). A small fraction of nickel was also detected. For the sample heat treated in synthetic air (Fig.V-4b), the microparticles did not fully release the metallic Al as depicted with the EDS surface analyses (Fig.V-

by

  Trunov et al. who estimated the formation of thicker oxide shells for smaller microspheres in dynamic dry air atmosphere [2]. The bright spots in the Al-containing microspheres were attributed to Ni enrichment (Fig. V-4b). The cross-section micrographs and the corresponding XRD patterns of the surface for both systems are given in Fig. V-5. For the sample annealed in Ar, a continuous diffusion layer with a thickness of approximately 50-60µm was formed on the nickel substrate (Fig. V-5a). The top coat was composed of hollow spheres with thin oxide shells (barely visible with optical microscopy). The crystallographic structure of the Ni2Al3 phase was identified by XRD (Fig. V-5c). This phase resulted from the selfpropagating high-temperature synthesis (SHS) occurring between Ni and Al [21,28]. Additional NixAly phases were identified at the reaction front between the nickel substrate and the Ni2Al3 layer as shown by Pedraza et al. [21]. According to literature, the crystallization of the natural amorphous alumina layer into γ-Al2O3 takes place at around 550°C [2,3,13,14]. A special attention was therefore given on the XRD pattern of Fig. V-5c to identify the characteristic patterns of γ-Al2O3, by comparison with literature [2,20,29-31]. However, γ-Al2O3 and the other polymorphs of alumina were not detected by XRD and the top coat structure observed in Fig. V-4a seemed to be transparent to the X-rays. This might come from the limited thickness of the oxide shells developed on the microspheres. As a result,

  pattern in addition to Ni2Al3 (Fig. V-5c). As suggested by the SEM surface characterization (Fig. V-4b) and the cross-section micrograph (Fig. V-5b), metallic Al was still present in the top coat after the diffusion heat treatment (Fig. V-5c). This suggests that the oxide grown at the surface of the Al microparticles blocked the full release and/or reaction of Al with the gas [7,10]. As a result, metallic Al was still contained in the smaller microparticles in the top coat structure of Fig. V-5b. The NiAl3

  surface images of the top coat structure of the coatings after the complete heat treatment are given in Fig.V-7. For the sample heat treated in Ar (Fig.V-7a), the top coat structure was very close to the one obtained by Pedraza et al. after exposure to 1150°C in He-4%H2 atmosphere[1]. Even though the microspheres were partially sintered after heat treatment in the inert gas, they were also deformed and broken. This could be related to the volume shrinkage occurring during the γ-Al2O3 → α-Al2O3 transition[1,6,20,27]. A Ni enrichment of the surface was observed by EDS after the final step at 1100°C (Fig.V-7c) in comparison with the results obtained after the diffusion heat treatment (Fig.V-4c). Since the whole metallic Al reacted with the nickel substrate after the two hours dwell at 700°C in Ar (Fig.V-5c), this Ni enrichment probably resulted from the outward diffusion of nickel during the two hours dwell at 1100°C. For the sample heat treated in synthetic air, the microspheres were less deformed and more sintered (Fig.V-7b) than the ones observed after heat treatment in Ar. Some microspheres were burst after annealing and resulted in oxide crusts similar to the ones observed by other authors[6,19]. The formation of these oxide crusts confirmed the important thickening of the alumina shells in synthetic air. Although the bright contrasted areas were found to be rich in Ni (about 10 at.% Ni), the overall Ni content of the surface was lower for the sample heat treated in synthetic air (Fig.V-7c). This difference might come from the thickness of the oxide shells formed after heat treatment. Indeed, the oxide shells formed in synthetic air were thicker than the ones obtained in Ar (Fig. V-8a and b) resulting in higher quantifications of Al and of O with respect Ni at the surface of the top coat in contact with the gas. This was also confirmed in the XRD patterns of Fig. V-8c, where the characteristic peaks of α-Al2O3 were better resolved in synthetic air than in Ar (Fig. V-8c). The microstructures of the diffusion layers formed after the complete heat treatment were similar in both atmospheres (Fig. V-8a and Fig. V-8b). The coatings were composed of two different layers: an external layer with the β-NiAl main composition and the γ'-Ni3Al one at the grain boundaries and an inner layer with the γ'-Ni3Al composition at the interface with the nickel substrate. Some β to γ' transformation was observed at the extreme surface of the coatings. This transformation was probably a consequence of Al outward diffusion to form the α-Al2O3 TGO scale at the extreme surface of the coating (confirmed by Raman-microspectrometry). For the same amount of slurry deposited on the surface, the overall thickness of the coating was larger after heat treatment in Ar (Fig. V-8a) than in synthetic air (Fig. V-8b). This suggests that the unreacted fraction of Al with the nickel substrate was greater in synthetic air than in Ar and therefore remained in the top coat structure to oxidize into alumina (Fig. V-8c).

Fig. V- 7 .

 7 Fig. V-7. Surface morphology of the top coat structure after the complete heat treatment in (a) Ar and (b) synthetic air and (c) corresponding EDS surface analyses.

  performed along the gas flow direction of the thermobalance (Fig. V-10). For the sample heat treated in synthetic air (Fig. V-10b), only α-Al2O3 was detected whereas both γ-Al2O3 and α-Al2O3 were detected for the sample heat treated in Ar (Fig. V-10a). The gas flow was found to exert an influence on the oxidation rate of the microspheres in the inert atmosphere. The presence of γ-Al2O3 after exposure to 1100°C for two hours was not relevant with the time-temperature diagram of alumina

Fig. V- 9 .

 9 Fig. V-9. Thermogravimetric curves obtained for the complete heat treatment of pure nickel slurry coated with Al in Ar and in synthetic air.

  Fig. V-10. Raman surface analyses of pure nickel slurry coated with Al after complete heat treatment in (a) Ar and (b) synthetic air.

V - 173 -II. 3 .

 1733 Complete heat treatment with the introduction of synthetic air upon annealingIn paragraphs II.1 and II.2, it was shown that the composition of the atmosphere (Ar or synthetic air) had a critical effect on both the oxidation of the Al microparticles and their sintering in the top coat structure. The objective was here to propose a single heat treatment to elaborate a functional thermal barrier system (diffusion zone + top coat structure) by switching the composition of the atmosphere during the heat treatment. Two different heat treatments were performed (Fig. V-13) as follows:-heating from ambient temperature to 400°C, one hour dwell at 400°C and then heating from 400°C to 700°C in flowing Ar (20 mL/min), -stop of the flow of Ar and purge of the reaction chamber with a vacuum pump for 8 seconds after reaching 700°C ("Switch 1") or at the end of the two hours dwell at 700°C ("Switch 2"), -introduction of synthetic air in the reaction chamber with a flow of 16 mL/min, -second part of the heat treatment (1100°C) in synthetic air.The complete rejuvenation of the atmosphere requires approximately 15 minutes and can be identified on both thermogravimetric curves by a step in Fig. V-13 (synthetic air is denser than Ar). As observed in Fig. V-6 and Fig. V-9, the binder was fully decomposed after one hour at 400°C. The interpretations of the thermal analyses were based on the work of Trunov et al. (Fig. V-2). For the "Switch 1" sample, a stepwise oxidation was observed as soon as synthetic air was introduced in the reaction chamber at the beginning of the plateau at 700°C. In contrast, the "Switch 2" sample presented a slow mass increase during the two hours of annealing at 700°C similar to the samples heat treated in Ar in paragraphs II.1 and II.2. These observations confirmed that the oxidation rate attributed to the crystallization of the amorphous alumina into γ-Al2O3 and its subsequent healing (stage II in Fig. V-2)

  The surface morphology of the top coat structures obtained after annealing are given on Fig. V-14a and Fig. V-14b. The hollow alumina microspheres were easily identified on both samples [1,21]. The top coat structure of "Switch 2" sample (Fig. V-14b) was very close to the one obtained for the sample heat treated in Ar (Fig. V-7a). Shrinkage of the oxide shells occurred and suggested the occurrence of the γ-Al2O3 → α-Al2O3 transition[1,6,20,27]. Shrinkage of the oxide shells was also observed for the "Switch 1" sample (Fig.V-14a). However, the introduction of synthetic air at the beginning of the 2 hours dwell at 700°C seemed to result in the thickening of the oxide shells. These observations are in good agreement with data from Ref.[1] after exposure of the Al microspheres at 1150°C in inert or oxidizing atmospheres. This would explain the significantly higher mass gain observed for the "Switch 1" sample in comparison with the "Switch 2" one on the thermogravimetric curves (Fig.V-13). EDS window analyses were performed on several areas of the samples (results not shown). No significant difference was revealed by the EDS measurements except that some Ni-rich spots were observed for the "Switch 1" sample. This could indicate the outward diffusion of nickel upon annealing at 1100°C. The cross-section micrographs of the full thermal barrier coatings elaborated on pure nickel after annealing are given in Fig. V-15. The microstructure of the diffusion layer was similar in both conditions. The light optical microscopy allowed to distinguish the different NixAly phases formed on the surface through differences of contrast (Fig. V-15a and Fig. V-15b). The diffusion layer was composed of a thin intermediate γ'-Ni3Al layer at the interface with the Ni substrate and the β-NiAl coating matrix. Some β to γ' transformation was observed at the grain boundaries and at the extreme surface of the coating.These observations were confirmed by EDS spot measurements and by XRD analyses (Fig. V-15c).

Fig. V- 13 .

 13 Fig. V-13. Thermogravimetric curves obtained for the complete heat treatment of pure nickel slurry coated with Al in Ar and in synthetic air. The sharp mass variation corresponds to the purge of the reaction chamber before introduction of synthetic air.

Fig. V- 14 .Fig. V- 16 .

 1416 Fig. V-14. Surface morphology of the top coat structure after the complete heat treatment in Ar and synthetic air with (a) "Switch 1" and (b) "Switch 2" and (c), (d) corresponding Ramansurface analyses.

  . V-17b). Since the introduction of synthetic air was done at the end of the two hours dwell time for this sample (Fig. V-13), the amount of Al flowing out from the spheres must have been the same than for the sample presented in Fig. V-5a. However, the overall mass gain was close to the "Switch 1" sample suggesting the growth of a γ-Al2O3 layer. It is important to mention here that for these three samples the conversion of the alumina shells into α-Al2O3 was complete. This was confirmed with both XRD and Raman analyses. The TGO scales developed in synthetic air (Fig. V-16) were expected to have little influence on the overall mass variation since the specific surface area of the top coat structure was much greater than the surface of the sample. Fig. V-18 allows to appreciate the difference in thickness of the oxide shells developed after the different annealing conditions.

  the particles and the release of Al from the metallic core. Based on the observations of the present work and from literature, the mechanisms of formation of the individual hollow alumina microspheres in Ar and in synthetic air are proposed in Fig. V-19. At room temperature, a native amorphous oxide layer covers the Al microparticle. The amorphous alumina layer grows until reaching a critical thickness of 4 to 5 nm to crystallize into γ-Al2O3 at about 550°C[3,11,12,16]. This transformation induces nanocracks (fractures) in the alumina shell since γ-Al2O3 has a greater density than the initial amorphous layer[6,15]. The composition of the atmosphere (i.e. 2 ) was then critical on the mechanisms of formation of hollow alumina microspheres. The mechanisms presented on Fig. V-19 are described separately for the sake of clarity.

  forming a continuous Ni2Al3 layer with hollow alumina microspheres on top (Fig. V-11). At this stage, the alumina shells are very thin (Fig. V-11) with smooth surfaces (Fig. V-4a). Increasing the temperature up to 1100°C led to the progressive transformation of the γ-Al2O3 into α-Al2O3. Note that this

III. 1 .

 1 results in the thickening of the alumina shell and in an enhanced sintering of the individual alumina microspheres through the formation of a α-Al2O3 connecting network (Fig. V-7b and Fig. V-18b). The second acceleration of the oxidation rate upon heating between 700°C and 1100°C (stage III on Fig. V-17b) could therefore be attributed to the fast oxidation of molten Al in the top coat structure. of the atmosphere (Ar, air) for the Cr-Al double-layer system Characterization of the top coat structure after heat treatment in Ar For this study, the Cr-Al double-layer system has been developed on pure nickel after heat treatment in Ar with a ratio close to 50Cr-50Al in wt.% (cf. paragraph V.3. in Chapter V). Fig. V-20 presents the cross-section images of the top coat structure obtained after different annealing treatments in flowing Ar. As observed for the simple Al slurry coating elaborated in Ar (Fig. V-18a), the alumina oxide shells remained very thin with a relatively weak sintering. These results are in good agreement with data from Ref. [1] after annealing in He-4%H2 inert atmosphere. Even though the Cr-Al double-layer system was highly beneficial to limit the SHS reactions between Ni and Al upon aluminizing, the Cr interlayer did not significantly change the top coat structure in flowing Ar. A BSE surface image of the top coat structure after the "conventional low-activity" heat treatment is given on Fig. V-21a. The differences in contrast allowed to distinguish the hollow alumina microspheres (dark areas) from the intermetallic phases (bright areas). This suggested that a fraction of the top coat structure spalled off upon annealing. Fig. V-21b shows an isolated alumina hollow sphere at higher magnification.

Ar. III. 2 .

 2 Fig. V-20. BSE cross-section images of the top coat structure after annealing for (a) 400°C-3h + 650°C-5h, (b) 400°C-3h + 650°C-5h + 1000°C-3h and (c) 400°C-3h + 1080°C-6h in flowing Ar.

  The top coat structure was close to the one obtained by Montero et al. on IN738 Ni-based superalloy after exposure at 800°C for 100 hours [23]. The oxide shells were characteristic of the microstructure of α-Al2O3. The formation of α-Al2O3 was confirmed by XRD (Fig. V-25) and Raman surface analyses.

  Fig. V-19 in synthetic air.

Fig. V- 23 .

 23 Fig. V-23. Thermogravimetric curves for the Cr-Al double-layer system elaborated on pure nickel in the different conditions (cf. Fig. V-22) for the (a) complete heat treatment, (b) isothermal section at 400°C, (c) heating ramp at 5°C/min between 400°C and 1080°C and (d)isothermal section at 1080°C.

Fig. V- 26

 26 Fig. V-26 presents the cross-section micrographs of the coating systems obtained after the different annealing conditions on pure nickel substrate. As suggested by the thermogravimetric curves and the surface characterization of the samples, synthetic air had a critical effect on the coating microstructure.

  Fig. V-26. Cross-section micrographs of the 50Cr-50Al double-layer system elaborated on pure nickel after the different annealing conditions (TC: Top coat, OL: Oxide layer, DL: Diffusion layer and IO: internal oxidation).

Fig. V- 30 .Fig

 30 Fig. V-30. DSC curves (heating branches at 5°C/min) obtained for the mixtures of Almicroparticles and oxide powders.



  . V-26). Note that the microstructure of the coatings was completely different between sample S2 and sample S4 (Fig. V-26). This suggested that the nature of the atmosphere (inert or oxidizing) had a tremendous influence upon heating between 400°C and 1080°C (Fig. V-22). Since the heating ramp was performed in Ar for sample S4 (Fig. V-22), molten Al could be released from the Al microparticles (Fig. V-19). Simultaneously, the peripheral oxidation of Cr microparticles was limited, which fostered the synthesis of AlxCry phases in the slurry deposit (cf. paragraph V.3. in Chapter III). When air was introduced into the reaction chamber (at about 1080°C), interdiffusion between nickel from the substrate and the Al-Cr phases just initiated resulting in the formation of isolated diffusion zones (Fig. V-26). An excess of Cr microparticles did not react with Al and oxidized on the surface between the AlxCry phases and the nickel substrate (cf. Fig. in Chapter III-34). The nickel substrate also oxidized into NiO on the uncoated areas. Then Cr2O3 and NiO reacted together to result in NiCr2O4 spinels in the top coat [68]. This would explain the fast mass gain of sample S4 after introduction of synthetic air (Fig. V-23c).  Influence of the Cr interlayer on the mechanisms of formation: Upon heating in synthetic air (sample S2), the peripheral oxidation of Al and Cr microparticles was considerably promoted as depicted by the sharp mass gain on thermogravimetric curves of Fig. V-23c. The peripheral oxide developed on Cr microparticles acted as a diffusion barrier for Al since the formation of Cr2O3 delayed the initiation of their exothermic reaction (Fig. V-30). Consequently, metallic Al remained in the top coat and did not react with the underlying nickel substrate. The presence of Al-Cr intermetallic phases surrounded by α-Al2O3 in the top coat structure of sample S2 (Fig. V-27) suggested that the reduction of Cr2O3 by molten Al occurred in the top coat. Molten Al was probably coming from the volume shrinkage of the alumina shells upon heating releasing liquid Al. This hypothesis is in good agreement with the XRD surface analyses of Fig. V-25. Therefore, the Cr interlayer promoted the oxidation of Al into α-Al2O3 in the top coat and the sintering of individual hollow alumina microspheres (Fig. V-24). Influence of the Cr-Al ratio on the mechanisms of formation of complete thermal barrier systems: In order to investigate the influence of the initial Cr-Al ratio on the microstructure of the complete thermal barrier coatings, pure nickel samples were slurry coated with a 25Cr-75Al double-layer system and annealed (Fig. V-22). Except for the sample heat treated in flowing Ar (S6 on Fig. V-31), the microstructure of the coatings was completely different than that of the 50Cr-50Al system (Fig. V-26).

  . The SEM characterization of sample S1 showed the random distribution of diffusion layers and oxide layers on the substrate surface. A BSE cross-section image and the corresponding X-ray maps of Ni, Al, O and Cr are given on Fig. V-32. The evidence of alternating β-NiAl diffusion layers and oxide layers can be observed along the substrate surface on Fig. V-32. The β-NiAl diffusion layers were surrounded with α-Al2O3 whereas the oxide layers were composed of a Ni-rich oxide (NiO) and a Cr-rich one (Cr2O3), according to local Raman analyses. The top coat structure was composed of hollow alumina spheres with thick oxide shells and presenting a good sintering at the naked eye (Fig. V-32b and Fig. V-32c). Note that AlxCry and NixAly intermetallic phases were also observed in the top coat structure. Their random distribution suggested that highly exothermic reactions occurred upon annealing (i.e. aluminothermic reactions). The oxidation of Cr microparticles into Cr2O3 was considered to cause the heterogeneous growth of the diffusion layers on the surface. Thick Cr2O3 layers were thus formed between the β-NiAl diffusion layers (Fig. V-32d). Cr2O3 acted as diffusion barrier for Al which preferentially reacted with NiO. Since the reaction between Al and NiO and eventually Ni is highly exothermic, the heat released could subsequently melt the Al around the ignition zone. This resulted in the formation of random diffusion islands surrounded by α-Al2O3. When the whole metallic Al is consumed, oxidation of the nickel substrate into NiO occurs in the "uncoated" areas.



  Fig. V-32.BSE cross-section image of the sample S1 (25Cr-75Al double-layer system) after annealing in synthetic air and corresponding X-ray maps of (a) Ni, (b), Al, (c), O and (d) Cr.

Fig. V- 33 .

 33 Fig. V-33. (a) BSE cross-section image of sample S4 (25Cr-75Al double-layer system) after annealing and (b) details of the TGO scale and the top coat structure formed.

-IV. 1 .

 1 investigated for the formation of a complete thermal barrier coating. Upon annealing in Ar, the Cr interlayer did not affect the microstructure of the top coat layer since fast release of molten Al from the microspheres occurred upon heating. This resulted in the formation of AlxCry phases in the slurry deposit and subsequent interdiffusion with nickel from the substrate to form continuous NixAly diffusion layers. When annealing was performed in synthetic air, the mechanisms of formation of the coatings were completely different. The high 2 in synthetic air promoted the oxidation of the nickel substrate and of the Al and Cr microparticles thus modifying their reactivity. Peripheral oxidation of Cr microparticles appeared to suppress the exothermic reaction associated with the formation of AlxCry phases at about 640°C. Consequently, the formation of Cr2O3 acted as a diffusion barrier and the Al was contained in the top coat structure. This resulted in the thickening of the hollow alumina spheres and their extensive sintering. After annealing, a thick oxide layer grew between the nickel substrate and the top coat structure with Cr2O3, NiCr2O4 and NiO oxides. This layer acted as a ceramic bond coat and ensured the adherence of the alumina top coat structure. Outlook:The present study opened a new way for the synthesis of intermetallic-ceramic composites based on aluminothermic reductions of NiO and of Cr2O3. The appearance of such reactions resulted in increased sintering of alumina top coat structures for thermal barrier coatings applications. The influence of the annealing conditions and of the architecture of the coatings was investigated for a better comprehension of the driving forces for the functionalization of thermal barrier systems. Since the optimization of the complete thermal barrier systems was not completed for functional applications, Chapter VI will be focused on the environmental resistance of the diffusion coatings (i.e. aluminide bond coats). The main objective of Chapter VI was to emphasize the functional applications of the diffusion coatings elaborated during this work by investigating their environmental resistance in different conditions. Environmental resistance of diffusion coatings elaborated by slurry I. State of the art............................................................................................................... I.1. Simulation of Type-II hot corrosion conditions at laboratory scale ............................................ I.2. Oxidation resistance of simple and modified aluminide coatings ................................................. II. Oxidation resistance of uncoated and slurry coated nickel at 700°C ....................... II.1. Characterization of the diffusion coatings elaborated on pure nickel ......................................... II.2. Influence of pSO3 on the oxidation-sulfidation of pure nickel ....................................................... a) Oxidation of pure nickel in synthetic air ................................................................................... b) Oxidation-sulfidation of pure nickel in synthetic air + 0.5% SO2......................................... c) Discussion.......................................................................................................................................... II.3. Influence of pSO3 on the oxidation resistance of simple and Cr-modified slurry coatings .... a) Oxidation in synthetic air .............................................................................................................. b) Oxidation in synthetic air + 0.5% SO2 ........................................................................................ c) Discussion.......................................................................................................................................... III. Type-II hot corrosion resistance of uncoated and slurry coated nickel at 700°C ... III.1. Hot corrosion of pure nickel at 700°C............................................................................................... III.2. Hot corrosion of simple and Cr-modified slurry aluminide coatings on pure nickel .............. III.3. Discussion .................................................................................................................................................. IV. Oxidation resistance of Cr-Al double-layer system developed on CM-247 LC ....... Characterization of the low-activity high-temperature (LAHT) coating .................................... IV.2. Cylic oxidation at 1000°C ...................................................................................................................... IV.3. Isothermal oxidation at 1000°C ........................................................................................................... IV.4. Discussion .................................................................................................................................................. V. Summary and outlook .................................................................................................. References ............................................................................................................................. As presented in Chapter I, different modes of degradation can be observed on materials exposed to high-temperature environments. Depending on the temperature regime and the activity of the aggressive substances, the corrosion rate can be significantly modified (cf. Fig. I-7 in Chapter I). In this

II- 9

 9 of Chapter II (with the collaboration of Dr. Ing. Mathias C. Galetz and Dr. Xabier Montero from DECHEMA-Forschungsinstitut). The oxidation of SO2 into SO3 following reaction (Eq. VI-2) was fostered using iron zeolite catalysts placed upstream of the samples in the hot zone of the horizontal furnace. The calculated equilibrium 3 from data in Ref. [4] was 7.5x10 -3 atm at 700°C. According to the Na-Ni-O-S stability diagram (Fig. VI-2a), the available 3 was therefore high enough to provoke the acidic reaction of NiO following reaction (Eq. VI-1). Considering now the Na-Al-O-S (Fig. VI-2b) and the Na-Cr-O-S (Fig. VI-2c) stability diagrams, both Al2O3 and Cr2O3 are stable in the SO2-containing atmosphere. Therefore, the acidic reaction of Al2O3

Fig. VI- 2 .

 2 Fig. VI-2. Superimposed stability diagrams of (a) Na-Ni-O-S, (b) Na-Al-O-S and (c) Na-Cr-O-S systems at 700°C from Ref. [7]. The blue cross refers to synthetic air equilibrium and the red one to synthetic air + 0.5% SO2 according to this work. The black one refers to the O2 + 0.1%SO2 equilibrium from Ref.[7].

II. 1 .

 1 resistance of uncoated and slurry coated nickel at 700°C Characterization of the diffusion coatings elaborated on pure nickel The SEM cross-section images of the simple aluminide and the Cr-modified aluminide coatings elaborated on pure nickel are respectively given in Fig. VI-3a and Fig. VI-3b. EPMA line scans were also performed across the coating thickness to identify the different phases formed on the nickel substrate after heat treatment in Ar. The mechanisms of formation for the high-Al activity slurry aluminizing in

(

  Fig. VI-3. BSE cross-section images and corresponding EPMA line scans of (a), (b) simple and (c), (d) Cr-modified slurry aluminide coatings elaborated on pure nickel.

Fig. VI- 4 .II. 2 .

 42 Fig. VI-4. SEM surface images of (a), (b) the simple and (c), (d) the Cr-modified aluminidecoating elaborated on pure nickel after light grit blasting.

  99%) was submitted to oxidation at 700°C in synthetic air under atmospheric pressure. The weight gain curves, recorded by TGA, are presented in Fig. VI-5. The parabolic rate constant , calculated from the transformed curves, was in the same range of order that data from literature (Table VI-1).

Fig. VI- 5 .

 5 Fig. VI-5. Oxidation curves of pure nickel obtained by TGA in synthetic air at 700°C.

Fig. VI- 7 .

 7 Fig. VI-7. (a) XRD patterns and (b) Raman surface analyses of pure nickel oxidized in synthetic air at 700°C.

Fig. VI- 11 .

 11 Fig. VI-11. (a), (b) BSE cross-section images of the corrosion scales formed on pure nickel after exposure to synthetic air + 0.5% SO2 at 700°C for 300 hours and (c) corresponding EPMA mappings of Ni, S and O.

  2). The Raman local analyses performed in the cross-sections (Fig. VI-12b) confirmed the formation of NiO in the dark regions of Fig. VI-12a. The reference signal of Ni3S2 (highlighted in Fig. VI-12b) was obtained from these analyses with four peaks at approximately 190, 225, 306 and 352 cm -1 .

Fig. VI- 12 .

 12 Fig. VI-12. (a) BSE cross-section image of the nickel sulfide layer formed on the substrate surface (cf. Fig. VI-11b) and (b) Raman spot analyses in the different areas.

  observations are in good agreement with the thermodynamic data from Gheno et al. (Fig. VI-2a) considering the stability domains of NiO and NiSO4 with respect of the 3 . The sulfation of NiO fostered the formation of porous and unprotective oxide scales as observed on the surface macrographs (Fig. VI-9) and in cross-section (Fig. VI-10). The formation of NiSO4 was also attributed to increase the S activity on the corrosion scale [46]. The presence of Ni3S2 at the scale/substrate interface (Fig. VI-11c and Fig. VI-12a) and within the NiO scale (Fig. VI-11c) confirmed

4 6 )II. 3 .

 63 Ni(s) + NiSO4 s → 4 NiO s + [Ni-S](l) (Eq. VI-5) Ni(s) + [Ni-S] l → Ni3S2(s) (Eq. VI-The Ni-S solution was not detected in the corrosion scales but was included in the reaction sequence since it may form an intermediate reaction as proposed by Kofstad and co-workers [43-45]. The microstructures observed in Fig. VI-11a and Fig. VI-11b strongly supported the formation of a liquid solution at temperature (i.e. 700°C). According to the Ni-S phase diagram (Fig. VI-13), an eutectic is formed at 635°C and for a composition between 20 and 23 wt.% S [43,44]. The continuous supply of S-bearing molecules from the scale/gas interface towards the metal/scale interface probably sustained the formation of a Ni-S liquid solution at the reaction front. When the 2 was high enough (i.e. above the nickel sulfide layer), Ni-S protrusions oxidized into NiO (cf. Fig. VI-11b). Influence of pSO3 on the oxidation resistance of simple and Cr-modified slurry coatings a) Oxidation in synthetic air The slurry coatings presented in paragraph II.1 were oxidized in synthetic air to follow the mass gain curves and compare with the results obtained with pure Ni substrate. The oxidation curves of the different systems are given in Fig. VI-14. The oxidation rate of pure Ni decreased with the application of slurry coatings and suggested a parabolic growth rate with both simple and Cr-modified aluminide coatings. The values of the coatings were calculated from the transformed curves and are summarized in Table VI-3. By comparison with data from literature, the value of the simple

Fig. VI- 14 .

 14 Fig. VI-14. Oxidation curves of uncoated and slurry coated nickel with a simple and a Crmodified aluminide coating at 700°C in synthetic air.

4 ,

 4 Fig. VI-16 and compared with the values of pure Ni and the oxidation curves in synthetic air. Both coatings decreased the oxidation rate of the nickel substrate. The addition of 0.5% SO2 in the gas composition seemed to have no influence on the oxidation behaviour of Ni + Cr/Al (Fig. VI-16). On the opposite, the mass gain of the simple aluminide coating was significantly higher with 0.5% SO2 and reached approximately 10 mg.cm -2 after 300 hours exposure.

Fig. VI- 16 .

 16 Fig. VI-16. Mass gain as a function of square root of time for the different systems in synthetic air + 0.5% SO2 at 700°C. The oxidation curves in synthetic air are also given for comparison.

Fig. VI- 18 .

 18 Fig. VI-18. Macrographs and BSE surface images of (a), (b) the simple aluminide and (c), (d)the Cr-modified coating after 100 hours exposure in synthetic air + 0.5% SO2 at 700°C.

Fig. VI- 21 .

 21 Fig. VI-21. BSE cross-section image and corresponding EPMA mapping of the Cr-modifiedcoating after 300 hours exposure at 700°C in synthetic air + 0.5% SO2.

Fig. VI- 22 . 1 . 2 . 3 . 4 . 5 .

 2212345 Fig. VI-22. Secondary electrons surface image of a protrusion formed on the simple aluminide coating after 100 hours exposure at 700°C in synthetic air + 0.5% SO2 and corresponding Xray maps.

VI- 24 .

 24 Fig. VI-24. Light optical cross-section micrographs of high purity nickel half-embedded inNa2SO4 and exposed in dry air at 700°C.

  Fig. VI-25. BSE cross-section image and corresponding EPMA mapping of pure nickelembedded in Na2SO4 after 100 hours exposure at 700°C in dry air.

Fig. VI- 26 .

 26 Fig. VI-26. Tentative reaction paths for the exposure of high purity nickel in Na2SO4 and in dry air at 700°C based on Ref. [8].

  Fig. VI-30 are also in good agreement with the observations of Gheno et al. [8].

Fig. VI- 31 .

 31 Fig. VI-31. Top view of the Ni samples half-embedded in Na2SO4 and cross-section micrographs in the salt side for (a), (b) the simple aluminide and (c), (d) the Cr-modifiedcoating after 300 hours exposure at 700°C in dry air.

Fig. VI- 32 .

 32 Fig. VI-32. Top view of the Ni samples half-embedded in Na2SO4 after exposure at 700°C in synthetic air + 0.5% SO2 with (a), (b) the simple aluminide and (c), (d) the Cr-modified coating.

Fig. VI- 33 .

 33 Fig. VI-33. BSE cross-section images of (a), (b) the simple aluminide and (c), (d) the Cr-modified coatings elaborated on pure Ni after exposure to Type-II hot corrosion conditions at 700°C.

Fig. VI- 34 .

 34 Fig. VI-34. SEM cross-section image of the simple aluminide coating exposed to Type-II hot corrosion conditions for 20 hours at 700°C showing thin Al2O3 scale and Al2O3 intrusions.

Fig. VI- 35 .

 35 Fig. VI-35. BSE cross-section image and corresponding X-ray maps for nickel slurry coatedwith simple aluminide after 100 hours Type-II hot corrosion exposure at 700°C.

  Fig. VI-36b, confirmed withRaman local analyses). This indicates that Cr-containing ions were present in the salt chemistry and precipitated as Cr2O3 in a region where the solubility of the oxide is low[61,62]. According to Rapp and Otsuka[63], chromia solubilizes in fused Na2SO4 with chromate anions (CrO4 2-) as the dominant solute at high oxygen activities whereas chromite anions (CrO2 -) dominate at low oxygen activities.

  to precipitate as Cr2O3, cf. (Eq. in the reducing sites and thereby inhibit the corrosion attack by preventing the contact between the salt and the metal (i.e. prevent sulfidation)[61]. This would explain the sealing of the intermetallic coating with Cr2O3 which in turn fosters the nucleation and growth of a protective α-Al2O3 scale (Fig. VI-36b and Fig. VI-37).2 CrO4 2-→ Cr2O3(s) + 2 O 2-+ O2 (Eq. VI-11)

  . VI-37). The very low concentration of Na and the low concentration of S in the pit region suggest that it was mainly composed of oxides (i.e. dissolution/reprecipitation mechanisms). Similar microstructures were observed by Task et al. under Type-II hot corrosion conditions[64]. On the X-ray maps (Fig., it appeared that Al was mostly concentrated near the reaction front whereas Ni was observed above the initial surface of the coating.

  . VI-38). The difference of brightness of NiO and Ni3S2 with optical light microscopy highlighted the greater amount of sulfide formation after exposure to Type-II conditions (Fig. VI-38c and Fig. VI-38d) than without Na2SO4 (Fig. VI-38a and Fig. VI-38b).

  Fig. VI-38. Light optical micrographs of the corrosion scales formed on Ni substrate after exposure at 700°C to synthetic air + 0.5% SO2 (a), (b) without Na2SO4 and (c), (d) with Na2SO4.

  Fig. VI-37 and local Raman analyses. This considerably increased the hot corrosion resistance of the Cr-modified coating with a fast healing of the oxide scales. The Ni solubility in the salt was considerably hampered thus preventing the formation of Na2SO4-NiSO4 eutectic phases and Ni sulfidation (Fig. VI-37). The formation of Cr sulfides at the reaction front of the attacked regions was found to buffer the corrosion attack[65]. resistance of Cr-Al double-layer system developed on CM-247 LC

  is given on Fig. VI-40a. A light grit blasting of the surface using #220 alumina mesh was performed to remove the residues of the slurry after coating process (cf. paragraph II.2 in Chapter IV). As shown previously, the coating was composed of two main layers with an outer additive layer (A.L.) formed above the substrate surface based on β-NiAl and an inner interdiffusion zone (IDZ) enriched with precipitates containing refractory elements (Fig. VI-40b). The final OPS polishing allowed to reveal the precipitation of α-(Cr) phases within the β-NiAl additive layer.

Fig. VI- 40 .

 40 Fig. VI-40. (a) SEM cross-section image and (b) EDS concentration profile of the low-activity coating elaborated on CM-247 LC before oxidation tests.

  Fig. VI-41. (a) Weight change as a function of time and (b) instantaneous kp for Cr/Al slurry coated CM-247 LC upon cyclic oxidation at 1000°C in air.

  slurry coated CM-247 LC (Fig.. The SEM surface images suggested the great adherence of the oxide scales and very limited spallation or crack formation on the coating surface after cyclic oxidation. Lu et al. obtained similar oxide morphologies for the cyclic oxidation of Y/Cr-modified aluminide coatings [19]. At higher magnification (Fig. VI-42), a platelet-like morphology was identified and suggested the formation of α-Al2O3. Raman micro-spectrometry was carried out to characterize the oxide scales developed on the Cr/Al coating. Several measurements were performed on the sample surface to ensure reproducibility (Fig. VI-43). Raman surface analyses confirmed the formation of an exclusive α-Al2O3 scale even after 100 cycles at 1000°C (Fig. VI-43c).

Fig. VI- 42 .

 42 Fig. VI-42. SEM surface images of the oxide scales developed on the LAHT coating elaborated on CM-247 LC after cyclic oxidation at 1000°C in air.

Fig. VI- 44 .Table VI- 6 .

 446 Fig. VI-44. Summary of the EDS window analyses performed on the surface of the samples after elaboration of the coating (AS-COATED) and cyclic oxidation at 1000°C in air.

Fig. VI- 45 .

 45 Fig. VI-45. BSE cross-section images at the extreme surface of the Cr/Al coating showing the growth of the α-Al2O3 scale and the enrichment of Cr and Hf at the oxide subscale.

Fig. VI- 46 .

 46 Fig. VI-46. (a) SEM cross-section image and (b) EDS concentration profile of the low-activity coating elaborated on CM-247 LC after 100 cycles at 1000°C in air.

Fig. VI- 47 . 3 .

 473 Fig. VI-47. Microstructural evolutions of the LAHT coating elaborated on CM-247 LC after cyclic oxidation at 1000°C in air (electrolytically etched with 10 vol.% H2C2O4).

Fig. VI- 48 .Fig. VI- 49 .

 4849 Fig. VI-48. Weight gain as a function of time for Cr/Al slurry coated CM-247 LC under isothermal conditions at 1000°C in air.

  The reactivity between Al and Cr microparticles was first determined by DSC using Al-Cr mixtures of different compositions. A highly exothermic reaction occurred by reaction of Al and Cr microparticles at a temperature below the melting point of Al (i.e. 660°C). This exothermic reaction was associated with the formation of AlxCry intermetallic phases of different stoichiometry. The formation of these intermetallic compounds was found to be fast, involving liquid phases, and following self-propagating high-temperature synthesis (SHS) reactions. The following peritectic reactions sequence was proposed to account for the formation of intermetallic compounds: L + Al7Cr → L + Al11Cr2 → L + Al4Cr. The heat change of the system and the stoichiometry of the synthesized AlxCry phases was dependent on the initial Al-Cr ratio and on the size of the microparticles involved in the reaction. After the initial exothermic reaction, solid state diffusion mechanisms of Al and of Cr were mainly involved. The reactivity between Al-Cr mixtures and pure nickel, as a model material, was also studied and demonstrated that the addition of Cr limited the exothermic reaction usually observed between pure Al microparticles and nickel-based materials.

  predominant inward diffusion of Al at low temperature and outward diffusion of nickel at high temperature. The initial inward diffusion of Al resulted in the precipitation of low solubility elements (i.e. Cr, Ta and W) in the outer region of the coating. The annealing temperature (1000°C or 1080°C) had a tremendous impact on the final microstructure of the coating and was ascribed to the enhancedConclusions and perspectives-271-outward diffusion of nickel at 1080°C than the one at 1000°C. The high-activity coatings presented a fine distribution of Cr-rich phases and large substrate carbides in a -NiAl matrix.The different aluminizing treatments carried out on nickel-based materials in Chapters III and IV focused on the formation of the metallic bond coats. The different heat treatments were performed in Ar to prevent the oxidation of the microparticles. However, the ceramic top coats composed of hollow alumina spheres were generally very brittle after heat treatment in Ar. Therefore, the influence of the atmosphere (Ar, air) on the formation of complete thermal barrier systems was investigated (Chapter V).The nature of the atmosphere (Ar, air) had a significant influence on the oxidation of Al microparticles and on the release of Al upon aluminizing. The thermogravimetric measurements were used to interpret the experimental observations. The introduction of air in the reaction chamber was found to significantly increase the growth of the oxide shells surrounding the Al microparticles. The thickening of the alumina shells in air considerably increased the sintering of the microspheres and their adherence to the underlying bond coat by comparison with the samples annealed in Ar. Consequently, the mechanical resistance of the ceramic top coat structure was greatly improved by performing the annealing treatment in air. On the opposite, the atmosphere had little influence on the aluminizing of the nickel substrate when only Al microparticles were used as donors. The metallic bond coats were then based on the -NiAl phase after annealing at high temperature (i.e. 1100°C).When a Cr interlayer was placed between the nickel substrate and the Al microparticles, the mechanisms of formation were significantly different in Ar and in air. The oxidation of the Cr microparticles into Cr2O3 was found to considerably decrease their reactivity with Al at temperatures lower than 900°C. The peripheral oxidation of Cr microparticles suppressed the exothermic reaction identified at about 640°C between Al and Cr microparticles. Consequently, the Cr interlayer acted as a barrier for diffusion for the Al, which was oxidized in the top coat structure. This resulted in the formation of a highly sintered ceramic top coat structure based on hollow alumina spheres. The introduction of air in the reaction chamber was also susceptible to initiate aluminothermic reactions.

Fig. 1 .

 1 Fig. 1. BSE cross-section images and corresponding X-ray maps of Hf for (a), (b) an outwardly and (c), (d) an inwardly grown coating on CM-247 LC superalloy. The initial surface of the substrate (ISS) is highlighted with the position of the substrate carbides after the aluminizing process.

  uncoated and coated samples to evaluate the thermal insulation behaviour of the slurry coatings (cf. paragraph IV.I. in Chapter II). The thermal diffusivity measurements for uncoated pure Ni and Ni slurry coated with "PARTICOAT" are given as examples in Fig. 2a (2 samples for each system). The thickness of the alumina foam (hollow alumina spheres, cf. Fig. I-13 in Chapter I) was approximately 45µm for the two slurry coated samples.The thermal conductivity (Fig.2b) was calculated following (Eq. II-2) (cf. Chapter II) from the specific heat capacity and the density of pure nickel (values taken from the literature). The thermal conductivity of the slurry coated samples was lower than pure Ni samples, which highlighted the beneficial role of the alumina foam on the thermal insulation. The data given in Fig.2aand Fig.2bwas recalculated to estimate the thermal diffusivity (Fig.2c) hence the thermal conductivity (Fig.2d) of the alumina foam only. The values were in the same range of order than conventional YSZ (Yttria-Stabilized Zirconia) Air plasma spray coatings.

Fig. 2 .Fig. 3 .

 23 Fig. 2. (a) Thermal diffusivity measurements and (b) calculated thermal conductivity of pure Ni samples with no coating and with "PARTICOAT" and (c), (d) calculated values for the alumina foam (comparison with self-standing YSZ APS coatings).

  II.2. Composition and preparation of the slurries ...................................................................................... II.3. Deposition of the slurries by air brush on metallic substrates ....................................................... II.4. Thermal treatments ...................................................................................................................................

III. Environmental degradation of the slurry coatings ....................................................... III.1. Isothermal and cyclic oxidation............................................................................................................... III.2. Hot corrosion tests ...................................................................................................................................

Table II -1. Chemical composition (wt.%) of the investigated nickel-based materials.

 II

	Chapter II								Experimental methods
	Material	Ni	Cr	Si	Al	Co	W	Mo	Ti	Ta	Hf
	Pure Ni (EQ)	Bal.
	Ni20Cr (EQ)	Bal.	20.0	1.5
	CM-247 LC (DS)	Bal.	8.0	-	5.6	9.3	9.5	0.5	0.7	3.2	1.4
	DS200 (DS)	Bal.	8.4	-	5.0	10.0	12.5	-	1.8	-	2.2
	IN100 (EQ)	Bal.	10.0	-	5.5	15.0	-	3.0	4.7	-	-
						-42-

For pure Ni, Ni20Cr and CM-247 LC materials, the samples were prepared from cylindrical rods with a Struers Secotom-50 cutting blade in the form of coupons of approximately 2.0±0.2 mm thickness. Different Ni substrates were used depending on the application. The details are given in Table

II

-2. To investigate the reactivity between pure Ni and the metallic powders by DSC, a 4.0±0.1 mm diameter substrate was used allowing the introduction of the samples in the alumina crucibles. Due to the availability of the material, the development of the coatings and the hot corrosion tests were done using the Ni substrate with a diameter of 12.7±0.1 mm. The DS200 and IN100 materials were used to extrapolate the slurry coatings developed during the study on aircraft nickel-based superalloys (results not presented in the manuscript). For these two materials, the samples were prepared from nonrepairable gas turbine blades kindly provided by SR Technics Airfoils Services (Cork, Ireland). The use of engine run components as parent materials was considered as a first step to transfer the research of the present study to an industrial scale.

Table II -2. Summary of the substrates used for the different applications in the study.

 II

	Application	Material	Diameter (mm)	Supplier
	Reactivity between Ni and powders (Chapter III)	Pure Ni	4.0±0.1	Goodfellow
		Pure Ni	12.7±0.1	Goodfellow
	Development of the coatings (Chapters III, IV and V)	Ni20Cr	12.0±0.1	Goodfellow
		CM-247 LC	20.0±0.1	Siemens
	Hot corrosion tests (Chapter VI)	Pure Ni	12.7±0.1	Goodfellow
	Oxidation tests (Chapter VI)	Pure Ni CM-247 LC	12.7±0.1 20.0±0.1	Goodfellow Siemens
	Extrapolation to gas turbine blades	DS200 IN100	Airfoil Airfoil	SR Technics SR Technics

Pure Ni and Ni20Cr materials present a γ-(Ni) matrix whereas the CM-247 LC superalloy display a γ-

Table II -3. Characteristics of the metallic powders investigated in the study.

 II

	Powders	Supplier	Theoretical size (µm)	Experimental size (µm)	Geometry
	Al	Sibthermochim	<20	2-17	Spherical
	Cr (fine)	Goodfellow	<10	2-10	Complex
	Cr (coarse)	Goodfellow	38-45	20-70	Complex

Table II -4. EDS window analyses of the different metallic powders investigated in the study.

 II

		Al		Cr		O
	Powders	wt.%	at.%	wt.%	at.%	wt.%	at.%
	Al	96 ± 1	94 ± 2	-	-	4 ± 1	6 ± 2
	Cr (fine)	-	-	91 ± 2	75 ± 2	9 ± 2	25 ± 2
	Cr (coarse)	-	-	92 ± 2	77 ± 2	8 ± 2	23 ± 2

Table II -5. Compositions of the different slurries investigated in the study using the Al powder delivered by Sibthermochim and the fine Cr powder delivered by Goodfellow.

 II

	Slurry	Composition of the metallic part	Amounts in the pill box	Metallic part in the slurry composition (wt.%)
	Al	100 wt.% Al	2g Al	45
	Cr	100 wt.% Cr	4g Cr	61
	67Al33Cr	67 wt.% Al + 33 wt.% Cr	1.6g Al + 0.8g Cr	48
	44Al56Cr	44 wt.% Al + 56 wt.% Cr	1.5g Al + 1.9g Cr	57
	20Al80Cr	20 wt.% Al + 80 wt.% Cr	0.9g Al + 3.6g Cr	63

-6. Table II-6. Details of the heat treatments performed in the TGA to investigate the mechanisms of formation on pure Ni. Slurry systems Decomposition of the organic binder Diffusion step Homogenization step Pure Al Al-Cr mixtures Al-Cr double-layer Cr-Al double-layer

	400°C-3h	650°C-5h	-

Pure Al Al-Cr mixtures Al-Cr double-layer Cr-Al double-layer

		400°C-3h	650°C-5h	1000°C-3h
	Al-Cr mixtures Cr-Al double-layer	400°C-3h	1080°C-6h	-
	Pure Al [20,25,26]	400°C-1h	700°C-2h	1100°C-2h

Table II

 II

	-7. Therefore, by comparison of experiments 1a and 2a, the influence of Na2SO4 in a SO3-
	containing atmosphere was investigated whereas the comparison of experiments 1a and 1b was useful
	to study the influence of 3 on the oxidation behaviour of the materials and the coatings. The
	experiment 1b was done in a box furnace in DECHEMA-Forschungsinstitut calibrated with a type-K
	thermocouple (Ni-Cr/Ni-Alumel).
	For the hot corrosion tests, the slurry diffusion coatings were elaborated on pure Ni to clearly
	investigate the role of each alloying element on the hot corrosion resistance. Indeed, the complex
	composition of nickel-based superalloys (Table II-1) may disturb the corrosion mechanisms with the
	inclusions of substrate elements in the oxide scales developed on the surface. Since the experimental
	set up was build up for this study, pure nickel samples were submitted to each batch as reference
	samples. By comparison with literature [29,47,48], it was then possible to evaluate the corrosiveness
	of the atmosphere and validate the experimental set up.

Chapter II Experimental methods -55- Table II-7. Synthetic table for the different corrosion tests done in DECHEMA- Forschungsinstitut. The thermogravimetric measurements were done in La Rochelle.

	Test	Principle	Atmosphere	Temperature (°C)	Furnace
	1a	Half-embedding	Synth. air + SO2/SO3		Horizontal furnace (Fig. II-9b)
	1b	Half-embedding	Lab air		Box furnace
	2a	Without salt	Synth. air + SO2/SO3	700	Horizontal furnace (Fig. II-9b)
			Synth. air		TGA
	2b	Without salt
			Lab air		Box furnace
	IV

. Characterization methods IV.1. Calorimetry and thermal analyses a) Differential Scanning Calorimetry (DSC)

Table II -8. Parameters for the different heat treatments performed with the DSC.

 II

	System	Heating rate	Maximum temperature	Cooling rate
	Elemental powders	10°C/min	1300°C	10°C/min
	Mixture of powders	10°C/min 5°C/min	1300°C 650°C	10°C/min 50°C/min (quench)
	Substrate + powders	10°C/min	1300°C	10°C/min
	Substrate + slurry deposit	5°C/min	1200°C	5°C/min

-56- b) Laser Flash Analysis (LFA)

Chapter III -Reactivity of Al and Cr powders for aluminizing purposes and mechanisms of formation of the slurry coatings on pure nickel I. Reactivity of Al and Cr powders with no substrate ..................................................... I

  .1. Preparation of Al-Cr mixtures ................................................................................................................ I.2. Reactivity of Al and Cr powders measured by DSC .........................................................................

II. Synthesis of Al-Cr intermetallic compounds ................................................................ III. Reactivity of Al, of Cr and of Al-Cr mixtures with pure nickel substrate .................. III

  .1. Reactivity of the metallic powders with pure nickel substrate ....................................................... III.2. Reactivity of Al-Cr mixtures with pure nickel substrate .................................................................. III.3. Synthesis of Al-Cr intermetallic compounds from Al-Cr mixed slurries for aluminizing .........

IV. Mechanisms of formation of the diffusion layers on pure nickel ................................. IV

  .1. Aluminizing of pure nickel with different Al-Cr mixed slurries: influence of Cr addition on the Al activity ..................................................................................................................................................................

	IV.2. Outward diffusion of nickel at high temperature (i.e. 1000°C and 1080°C) ...............................
	V

. Mechanisms of formation of the diffusion layers from Al-Cr and Cr-Al double-layer systems ....................................................................................................................................

  Mechanisms of formation for the Cr-Al double-layer system .........................................................

	V.1. Elaboration of Al-Cr and Cr-Al double-layer systems ......................................................................
	V.2. Mechanisms of formation for the Al-Cr double-layer system .........................................................
	V.3.

VI. Summary and outlook .................................................................................................. References .............................................................................................................................

Al-Cr mixture (wt.%) EDS spot location EDS spot measurements Possible phase Crystal structures observed by XRD Phase composition of synthesized products Al

							Dominant	Secondary
							aluminide	aluminide(s)
	67Al33Cr	1 2 3	81.7 83.3 80.4	18.3 16.7 19.6	Al11Cr2 [11] Al4Cr [9]	Al11Cr2, Al4Cr [26], (Cr)	Al11Cr2	Al4Cr
		4	17.5	82.5	(Cr) [9]
	44Al56Cr	1 2 3	82.1 82.1 12.2	17.9 17.9 87.8	Al11Cr2 [11] (Cr) [9]	Al11Cr2, Al4Cr [26], Al11Cr4 [17], (Cr)	Al11Cr2	Al4Cr, Al11Cr4
	20Al80Cr	1 2 3	75.2 76.2 70.9	24.7 23.8 29.1	Al11Cr4 [11] Al8Cr5 [11]	Al11Cr4 [17], (Cr)	Al11Cr4	Al8Cr5
		4	2.5	97.5	(Cr) [9]

(at.%) Cr (at.%)

Table III -4. Composition for the different slurries prepared in this work and corresponding Al activity in the powder mixtures. Slurry Al:Cr (wt.%) Al:Cr (at.%) Al activity Deposited mass (mg.cm -2 )

 III

	Al	100:0	100:0	1	10
	67Al33Cr	67:33	80:20	0.8 (High)	10-12
	44Al56Cr	44:56	60:40	0.6 (Intermediate)	15-18
	20Al80Cr	20:80	33:67	0.33 (Low)	25-30

  I.2. From Al-Cr mixed slurries ..................................................................................................................... I.3. Functional design: Cr-Al double-layer system ...................................................................................

II. Slurry aluminizing of the CM-247 LC nickel-based superalloy ................................. II.1. Aluminizing from Al-Cr mixed slurries ............................................................................................... II.2. Low-activity high-temperature (LAHT) coatings from the Cr-Al double-layer system .......... a) Low-activity conventional heat treatment (400°C-3h + 1080°C-6h) ................................. b) Complete heat treatment (400°C-1h + 700°C-2h + 1100°C-2h) ....................................... c) Discussion.......................................................................................................................................... II.3. High-activity high-temperature (HAHT) coatings from the Cr-Al double-layer system ........ a) Low-activity conventional heat treatment (400°C-3h + 1080°C-6h) ................................. b) Complete heat treatment (400°C-1h + 700°C-2h + 1100°C-2h) ....................................... c) Discussion.......................................................................................................................................... III. Summary and outlook .................................................................................................. References .............................................................................................................................................................

Table IV -1. Summary of the EDS spot measurements (Fig. IV-9b) and suggested phases.

 IV

	EDS spot		EDS spot measurements		Possible
	location	Al (at.%)	Si (at.%)	Cr (at.%)	Ni (at.%)	phase(s)
	1 2 3 4	9.0 6.0 17.3 32.6	0.9 1.1 2.2 1.9	70.0 71.6 49.7 7.2	20.1 21.3 30.8 58.3	α-(Cr) α-(Cr) β-NiAl + α-(Cr) β-NiAl

Table VI -1. Comparison of the parabolic rate constant calculated in this work with data from literature.

 VI

		This work	[39]	Literature	[41]
	(g 2 .cm -4 .s -1 )	4.3x10 -12	4.2x10 -12		8.9x10 -12

Table VI -2. Summary of the EDS spot measurements done in cross-section (cf. Fig. VI-11c). EDS spot location O (at.%) S (at.%) Ni (at.%) Possible phase

 VI

	Dark oxide scale	48.0	0.3	51.7	NiO
	Nickel sulfide layer	5.4	39.9	54.8	Ni3S2
	Nickel substrate	5.0	0.6	94.4	γ-(Ni)

Table VI -3. Comparison of the parabolic rate constant calculated in this work with data from literature. kp values (g².cm -4 .s -1 ) This work Literature

 VI

	Pure Ni	4.3x10 -12	4.2x10 -12 8.9x10 -12	[39] [41]
	Ni + Al	6.3x10 -15	3.0x10 -15 (bulk Ni-50Al, at.%)	[48]
	Ni + Cr/Al	2.2x10 -13	6.4x10 -14 (Pure Cr)	[49]

Table VI -5. Comparison of the parabolic rate constant calculated in this work with data from literature at 1000°C in dry air.

 VI

  Dans le cadre du projet Européen « PARTICOAT », le LaSIE a démontré la faisabilité d'élaborer des systèmes barrières thermiques complets (couche de diffusion + barrière thermique) en une seule étape à partir de barbotines (« slurries ») à base aqueuse contenant des microparticules d'Al. Dans cette étude, l'ajout de Cr comme dopant a été étudié. L'addition de Cr a permis d'abaisser l'activité de l'Al lors de l'étape d'aluminisation et de limiter les réactions exothermiques entre Al et substrat à base de nickel. L'optimisation des ratios entre Al et Cr a permis d'obtenir différentes microstructures de revêtement. Diverses architectures de dépôts ont également pu être testées grâce à la souplesse d'élaboration des revêtements par barbotines. L'influence de l'atmosphère (Ar, air) et celle des conditions de traitement thermique ont également été étudiées. Enfin, la durabilité des revêtements développés au cours de la thèse a été évaluée au cours d'essais de corrosion à chaud et d'oxydation.

Remerciements

Chapter 1 State of the art: context and motivations of the study -24-V. High-temperature protection of nickel-based superalloys

Ces travaux de thèse se sont déroulés long de ces trois années de thèse (bourse n° 2014Mes sincères remerciements à Monsieur Bruno Mortaigne, membre de l'état-major des armées et directeur scientifique au sein de la DGA, pour son accompagnement Je remercie chaleureusement le Docteur Xabier Montero (Chargé de recherche HDR à John de votre bienveillance lors de mon séjour sur l'île d'Emeraude et de m'

 Formation of "high-activity" coatings from the Cr-Al double-layer system:

The mechanisms of formation of the coatings for the 33Cr-67Al double-layer system were different than the ones observed for the 56Cr-44Al explaining the different microstructures observed (Fig.

IV-32b and Fig. IV-32d). One key aspect here is that inert substrate carbides were found in the outer region of the aluminized layer, which indicated that the initial substrate surface remained at the aluminized layer-slurry deposit interface (cf. Annex 1). Based on the work of Das et al. [41], the coating formation involved inward diffusion of Al upon heating through the formation of Al-rich β-NiAl phase.

Once the high temperatures reached (1034°C in the work of Das et al.), outward diffusion of nickel from the substrate was promoted. This differs from the state-of-the art "two-step" high activity processes where Al inward diffusion occurs in a first step at low temperature (700°C-900°C) and nickel outward diffusion is promoted in a second step at high temperature (1000°C-1150°C) to stabilize the β-NiAl phase [33,40,55]. Das et al. observed that the single-step process resulted in a nonstationary reaction front moving outwardly during the aluminizing process. The mechanisms of coating formation are summarized on Fig. IV-26b.

Since the coating formation of the 33Cr-67Al double-layer system started by inward diffusion of aluminium, the reaction (Eq. IV-3) was not governing the aluminizing process. In this system, the Cr interlayer acted as a barrier for Al ingress into the substrate (cf. Chapter III). One could therefore expect that this reaction is likely to happen in the slurry deposit. However, this reaction was presumably not the only driving force to explain the coating formation through the appearance of a liquid phase. To confirm that, a CM-247 LC slurry coated sample with the 33Cr-67Al double-layer system was annealed at a temperature lower than 1040°C. A SEM crosssection image of the sample after annealing (400°C-3h + 1000°C-6h) is given on Fig. IV-33. The microstructure of the aluminized layer was very close to the one obtained by Das et al. with singlestep high activity pack in the early stage of aluminizing [41]. A thin IDZ was formed under the main coating diffusion layer. The substrate carbides were found in the whole thickness of the aluminized layer attesting that its formation predominantly involved inward diffusion of Al. The AlxCry intermetallic compounds were still present above the aluminized layer after annealing. For the three hours dwell at 400°C (Fig. V-23b), a similar mass loss was observed for all samples and was attributed to the removal of the organic binder [38]. The composition of the atmosphere seemed to exert no significant influence on the thermogravimetric curves up to 400°C. However, the heating curves presented in Fig. V-23c suggested that the samples S1 and S2 were affected by synthetic air from approximately 550°C. The mass gain observed from 550°C for these two samples was attributed to the growth of the amorphous alumina layer [2,3,16]. A stepwise oxidation was then observed and centred at approximately 615°C suggesting the crystallization of the amorphous alumina layer into γ-Al2O3 [2,3,16]. For samples S1 and S2, the oxidation rate considerably increased after reaching 750°C.

 Thermogravimetric analyses:

The overall mass variation was much higher than the one recorded in 20c). For samples S1 to S4, the oxidation rate progressively decreased suggesting a parabolic growth rate. This was also observed for sample S5 with the introduction of synthetic air one hour after reaching 1080°C. The significant differences on the overall mass variation of the samples upon annealing suggested that the resulting coatings must be completely modified by the introduction of synthetic air at different stages of the annealing treatment.

For samples S3 and S4, the top coat presented a white colour (see the macrographs) suggesting the formation of α-Al2O3. The SEM surface images attested that the thickness of the oxide shells progressively decreased from sample S1 to sample S5. The volume shrinkage induced by the γ-Al2O3 to α-Al2O3 transformation led to the deformation of the alumina shells. This volume shrinkage was more pronounced for thin alumina shells (samples S3 to S5) because of their lower mechanical strength.

For sample S5, the top coat structure was no longer adherent but powdery-like. Some intermetallic phases were observed underneath the alumina hollow spheres presenting thin oxide shells.

As suggested with the Al X-ray map (Fig. This could be attributed to the significant growth of the NiO scale that masked the γ-(Ni) structure of the nickel substrate from the X-rays. The characteristic vibration modes of NiO were observed on the Raman spectra after 20 hours and 100 hours of exposure [42].

Table V-1. Raman shifts (cm -1 ) measured for the different spot analyses (Fig. V-26b).

This work

Fig. VI-20. BSE cross-section image and corresponding EPMA mapping of the aluminide coating after 300 hours exposure at 700°C in a protrusion region.

-232-A similar degradation mode was observed by Task et al. in Ref. [51] for the exposure of Ni-36Al alloy (at.%) in Type-II hot corrosion conditions (Na2SO4 deposit in a gas containing O2 + 0.1% SO2 at 700°C).

Since Na2SO4 deposit was not used in this work, it can be concluded that the formation of such protrusion is due to the presence of SO2-SO3 rather than Type-II hot corrosion mechanisms.

Therefore, in the work of Task et al., Na2SO4 acted as a catalyst to foster local breakage of the Al2O3 scale and internal oxidation of the nickel aluminide phase by selective oxidation of Al. This can be attributed to a higher 3 and a lower 2 at the metal surface when a salt deposit is present as explained by F.S. Pettit in Ref. [4].

If the propagation stage was clearly identified in this work, the initiation stage was more difficult to understand as it was not directly observed in the microstructures. However, the S enrichment (at.%) in low 2 atmospheres [52]. Thus, the cyclic degradation process could initiate and lead to the growth of the protrusion by selective oxidation of Al.

 Beneficial effect of Cr addition on the oxidation resistance in synthetic air + 0.5% SO2:

The presence of Cr in the aluminide coating was highly beneficial for the oxidation resistance in the SO2-containing atmosphere. Unlike the simple aluminide coating, the Cr-modified one was not attacked even after 300 hours exposure. This enhanced oxidation resistance was attributed to the fast formation of the stable α-Al2O3 phase on the coating surface. Since Cr2O3 was also detected with both XRD and Raman surface analyses, it seems that the nucleation of α-Al2O3 was promoted because of the initial formation of Cr2O3 [48,53,54,59]. Therefore, even though the extreme surface of the coating was quite rough (cf. Fig. , the formation of a pure α-Al2O3 scale was ascribed to limit Al depletion in the near-surface region and to prevent internal oxidation [52]. As a result, the Al depletion in the coating mainly occurred by interdiffusion with the pure nickel substrate because of the great concentration gradient. This led to the to γ' transformation at the grain boundaries of the coating layer (Fig. VI-21).

-246-

III.3. Discussion

Using the half-embedding method, it was possible to investigate three different zones on the samples: the atmosphere side (no contact with Na2SO4), the salt side (embedded in Na2SO4) and the interface between these two areas.

 Hot corrosion of pure Ni:

The hot corrosion tests of pure Ni using this method confirmed that a sufficient 3 is necessary to reproduce the Type-II conditions. When pure Ni half-embedded in Na2SO4 was exposed to dry air at 700°C, very little sulfidation of the substrate occurred even after 300 hours of exposure. In the atmosphere side (sample without contact with Na2SO4), a NiO scale was growing with a cellular morphology as observed for the oxidation of pure Ni in synthetic air (Fig. VI-6). In the salt side (sample embedded in Na2SO4), mixed NiO + Ni3S2 nodules have grown above the substrate surface. However, the microstructures suggested that there was no formation of liquid phases at temperature. The formation of nickel sulfides within the oxide was attributed to the decomposition of SO3 at the low 2 prevailing within the thick oxide scale following (Eq. VI-12) [8]. In these conditions, the sulfides can form within the oxide scale or at the metal/scale interface where the 2 is minimum and the 2 is maximum [8,47].

When pure Ni half-embedded in Na2SO4 was exposed to synthetic air + 0.5% SO2 at 700°C, the Type- VI -30). The formation of intermixed NiO + Ni3S2 layers within the corrosion scale suggested that the sulfide layer was growing together with NiO at the sulfide/oxide interface [8]. In such microstructures, Ni3S2 was ascribed to form a continuous network [66]. Since Ni diffusivity is much higher within the sulfides than within the oxides, the reaction kinetics was governed by outward diffusion of Ni in the Ni3S2 network [8]. The reaction of pure Ni with synthetic air + 0. -260-

The oxide scales were also characterized in cross-section using SEM (Fig The Cr-modified aluminide coating elaborated on CM-247 LC superalloy presented good thermal stability during cyclic and isothermal oxidation at 1000°C. The high-temperature oxidation resistance was conferred by the formation of a pure α-Al2O3 scale on the coating surface. This scale was found to exhibit good spallation resistance and a parabolic growth rate even in cyclic conditions. The light grit blasting process performed after coating elaboration to remove intermetallic phases resulting from the aluminizing treatment, was not found to influence the oxidation behaviour of the coating. The formation of the stable α-Al2O3 polymorph rapidly occurred. For longer exposure time (more than 500 hours), Ti, Hf and Ta were found to incorporate as TiO2 and HfO2 (and presumably Ta2O5) in the α-Al2O3 scale. Ti and Ta were usually ascribed to decrease the oxidation resistance of aluminide coatings at temperatures above 1100°C [70,72] whereas Hf increased the spallation resistance of α-Al2O3 scale [70,71]. However, they did not significantly affect the adherence of the α-Al2O3 scale for the oxidation time and the temperature investigated here.

 Microstructural evolutions upon oxidation:

The (Type-II hot corrosion conditions), the corrosion rate was even higher. This was attributed to an increase in S activity (i.e. 2 ) in the scale and at the substrate/scale interface because of the formation of Na2SO4-NiSO4 eutectic phases at the scale surface.

 Hot corrosion resistance of simple aluminide and Cr-modified coatings:

The application of slurry coatings on pure Ni considerably increased its oxidation-sulfidation resistance.

The Cr-modified coating offered a better protection than the simple aluminide coating in the SO2containing atmosphere at 700°C. The latter was prone to localized attacks in the form of protrusions with extensive internal oxidation. The mechanisms of formation and growth of the protrusions were described and were found to be controlled by both inward diffusion of O and outward diffusion of Al from the coating layer. The Cr-modified coating was not attacked after 300 hours exposure in the SO2-containing atmosphere. The protection efficiency of the coatings was attributed to their ability to form and rejuvenate protective oxide scales on the surface. When exposed to Type-II hot corrosion conditions, the corrosion attack was preceded by an incubation period going from 20 hours to 100 hours for the two slurry coatings (simple aluminide and Cr-modified). After 100 hours exposure, both coatings have shown corrosion attack. However, the propagation rate of the attack was considerably lower for the Cr-modified coating than for the simple aluminide. This was attributed to the rapid oxidation of the Cr-modified coating and its ability to reform protective scales (i.e. Cr2O3 and α-Al2O3).

On the other hand, the simple aluminide underwent fast degradation in Type-II conditions. The corrosion attack was found to initiate locally and then to propagate rapidly through the formation of liquid phases. This led to an important S enrichment at the interface between the substrate and the coating and an extensive metal dissolution (i.e. Ni and Al) within the sulfate deposit.

Based on the experimental results of this work, it appeared that the faster the establishment of protective oxide scales, the better the protection against the salt mixture. This therefore indicates that the pre-oxidation of the slurry coatings should increase their performances in hot corrosion conditions.