Keywords: L-PBF, Alliage 625, Traitement thermique, NiCrAlY, Oxydation, Propriétés mécaniques

I. Résumé long en français

Le superalliage 625 à base de Ni a été largement utilisé pour les applications sousmarines, aérospatiales et turbines à gaz en raison de sa bonne résistance à haute température, de sa résistance au fluage, de sa résistance à la corrosion et de son excellente soudabilité. Cependant, ses bonnes propriétés mécaniques contribuent à des difficultés d'usinage, qui ont conduit à l'intérêt pour l'application de cet alliage dans le domaine de la fabrication additive (FA). La technologie AM présente de nombreux avantages. Le plus grand avantage est qu'il peut être utilisé pour fabriquer des formes complexes sans outillage; conception pour la fabrication additive (DfAM) ce qui signifie qu'il n'y a aucune contrainte dans la conception des pièces Cette caractéristique de la FA est intéressante principalement dans le domaine aérospatial car elle peut réduire le poids des pièces du moteur en éliminant les matériaux inutiles à certains endroits, ce qui serait nécessaire pour une fabrication facile par un outillage conventionnel et réduit également les dommages d'usure dus au contact des pièces. En plus du domaine aérospatial, l'application de la technologie AM s'est étendue à d'autres industries et l'intérêt pour cette technologie augmente. L'étude de base des compétences de fabrication s'est poursuivie, mais essayer d'appliquer la technologie AM dans de nouveaux contextes devient de plus en plus courant, comme le traitement de multi-matériaux, comme les alliages acier / cuivre. La présente étude a évalué la faisabilité de l'application de l'alliage 625 à base de nickel fabriqué par AM, en particulier par fusion sur lit de poudre laser (L-PBF), dans des conditions industrielles réelles en appliquant un traitement post-thermique et en développant de nouveaux multi-matériaux à base d'alliage 625. Premièrement, en ce qui concerne les propriétés mécaniques, les propriétés de traction à température ambiante et à haute température ont été évaluées, et divers traitements thermiques ont Sur la base de cette étude, l'introduction du GBS dans l'alliage corroyé 625 a été réalisée. L'alliage forgé dentelé GB 625 a montré des motifs en zigzag de GB avec des carbures M23C6 et NbC au niveau des GB. Dans une étude plus approfondie, ce traitement thermique a été appliqué à l'alliage 625 produit au L-PBF. Le traitement thermique a été adapté pour être appliqué à l'alliage L-PBF 625 compte tenu de sa microstructure particulière. En tenant compte du rôle prédominant de la microstructure, 3 les propriétés de traction à température ambiante et à haute température des échantillons produits par chaque traitement thermique ont été comparées.

Une autre méthode a été suggérée pour améliorer les propriétés des superalliages et augmenter leurs températures de fonctionnement. Les revêtements de surface ont été développés pour protéger les substrats de l'oxydation, de la corrosion et des dommages thermiques à haute température afin que leurs propriétés mécaniques puissent être maintenues pendant un fonctionnement plus long. En règle générale, les revêtements de barrière thermique (TBC) ont un effet isolant, qui diminue la température à laquelle un superalliage est réellement soumis, ce qui permet au superalliage de résister à un fonctionnement à des températures d'environ 150 ° C au-dessus de leur limite supérieure habituelle. Ces TBC consistent généralement en une couche de finition externe en céramique (TC) pour l'isolation thermique et une couche intermédiaire de liaison métallique MCrAlY (BC, où généralement M = Ni, Co, Fe ou une combinaison de ceux-ci) pour tenir compte de la différence entre les coefficients de expansion entre le TC et le substrat en superalliage, et pour faire croître une couche d'oxyde protecteur mince, compacte et adhérente favorisant une résistance optimisée à la corrosion à chaud. L'alliage MCrAlY peut être utilisé comme matériau de revêtement de liaison et comme matériau de revêtement de recouvrement lui-même. L'épaisseur du revêtement de recouvrement MCrAlY est illimitée, sa résistance à l'oxydation / corrosion est proportionnelle à son épaisseur. Ainsi, un matériau MCrAlY est généralement requis sur un substrat en superalliage exposé à des environnements difficiles pour les deux raisons.

Le processus de revêtement est totalement différent de celui de la production de substrat en superalliage; les couches de revêtement sont généralement appliquées sur le substrat par projection d'air ou de plasma sous vide (APS / VPS), dépôt physique en phase vapeur par faisceau d'électrons (EB-PVD) et placage aux ions d'arc (AIP). Les processus séparés de fabrication d'alliages revêtus nécessitent une main-d'oeuvre, un coût et du temps considérables. Pour ces considérations, le L-PBF d'une couche de

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liaison NiCrAlY sur un substrat en alliage 625, lui-même produit par le L-PBF, est également considéré. Bien qu'il y ait eu quelques études sur la fabrication multimatériaux par la technologie AM, c'est la première fois, à notre connaissance, que le L-PBF est utilisé pour produire à la fois un substrat en superalliage et une couche de liaison MCrAlY. Nous pensons que cette méthode peut améliorer la liaison entre le substrat et le revêtement en raison de la nature du processus L-PBF, qui implique la refusion typique des couches sous-jacentes. La production finale de géométries enrobées complexes, telles que des lames, est un défi car les technologies PBF nécessitent le dépôt d'une couche entière de poudre, limitant ainsi le dépôt sur des surfaces planes. Cependant, on pense que ce problème pourrait être résolu prochainement en améliorant la partie de la machine comme la liberté de mouvement de balayage laser, en utilisant un réservoir à poudre multiple, etc. Dans ce travail, l'optimisation des paramètres de processus L-PBF pour MCrAlY est décrite sur la base de plusieurs critères, et les matériaux tels que produits sont minutieusement analysés.

De plus, les propriétés d'oxydation de l'alliage 625 revêtu de MCrAlY de manière optimale ont été évaluées par analyse gravimétrique thermique (TGA). En outre, les propriétés de résistance thermique, en particulier la résistance à la dilatation et à la contraction thermiques rapides, ont été évaluées par des essais de choc thermique. Les résultats ont démontré la faisabilité du revêtement MCrAlY sur le substrat de superalliage au moyen de L-PBF, et sa fonction s'est avérée raisonnable.

Par conséquent, nous concluons que le traitement post-thermique et le revêtement MCrAlY fabriqué au L-PBF améliorent les propriétés mécaniques de l'alliage 625 fabriqué par L-PBF à des températures de fonctionnement réelles (~ 982 ° C) au lieu de l'alliage 625 conventionnel. les effets de ces procédures ont été démontrés sur la base de la microstructure et des aspects fonctionnels des échantillons fabriqués.

Publications et contribution scientifique

Objet de la recherche  Fabrication d'un système bi-matière en alliage 625 revêtu de NiCrAlY par L-PBF.

 Chaque propriété du substrat et de l'alliage revêtu uniquement doit être évaluée et améliorée pour être certifiée pour la commercialisation / l'application.

 La microstructure de divers alliages L-PBF traités thermiquement 625 a été étudiée et leurs propriétés de traction ont été comparées à température ambiante.

 Pour considérer les conditions de fonctionnement réelles de l'alliage, une étude des propriétés à haute température a été envisagée. En outre, le GBS a été choisi comme un moyen d'améliorer les propriétés mécaniques à haute température.

Par conséquent, le mécanisme de GBS doit être étudié pour la première étape à appliquer pour la première fois au superalliage à base de Nb ajouté (alliage 625).

 Le mécanisme GBS a été compris pour la première fois avec un modèle en alliage.

 L'étape suivante consiste à appliquer le GBS à l'alliage corroyé 625, l'alliage standard.

 L'application de GBS à l'alliage corroyé 625 a été réussie.

A new observation of strain-induced grain boundary serration and its underlying mechanism in a Ni-20Cr binary model alloy

Dans cette étude, dans le cadre du renforcement GB à haute température, le 9 mécanisme du GBS a été étudié à travers un alliage modèle binaire Ni-20Cr pour préparer l'application à l'alliage 625 fabriqué par additif dans l'étude ultérieure. Un alliage entièrement traité en solution solide a été traité thermiquement sur la base de l'étude précédente adaptée pour un alliage à matrice Ni -Cr. Selon l'hypothèse du mécanisme de GBS, il n'y a eu aucun cas de GBS comme prévu. Cependant, le signe de GBS s'est produit en donnant une déformation de 5% pendant le processus de vieillissement direct / refroidissement lent dans lequel le GBS a eu lieu, qui a été conduit après le processus de solution solide. La microstructure de l'échantillon dans lequel le SGB s'est produit a été analysée par TEM et la méthode 3D-APT. Il y avait une densité élevée de dislocations autour des GB en raison d'une contrainte de 5% pendant le processus de vieillissement / refroidissement lent, ce qui a favorisé l'enrichissement en Cr près et au niveau des GB. Pour que l'échantillon présente le même comportement juste par traitement thermique, la température de dissolution a été augmentée pour induire un enrichissement en Cr près et au niveau des GB avec le concept de diffusion assistée par vide, et cela a bien fonctionné. L'analyse 3D-APT approfondie a montré que les échantillons dentelés présentaient un large enrichissement en Cr au niveau des GB dentelés. La caractéristique commune des deux échantillons était l'enrichissement en Cr, et la tension sur les GB était considérée comme le déclencheur du SGB. La déformation par calcul du premier principe était d'environ 0,97 kJ / mol. Ainsi, il a été démontré que le minimum de 0,97 kJ / mol de déformation peut conduire au SGB par les atomes interstitiels au niveau des GB.

First evidence of grain boundary serration in a specifically

Heat treatment design for superior high-temperature tensile properties of Alloy

produced by SLM

L'alliage 625 fabriqué au L-PBF a été soumis à des essais de traction à haute température. Compte tenu de l'étude récente qui décrit un essai de traction réalisé à température ambiante avec du L-PBF tel que construit et divers échantillons traités thermiquement, un traitement thermique GBS modifié (augmentation de la température de la solution solide de 1250 → 1300 ° C) a été appliqué à l'alliage L-PBF. 625 (échantillon 'GBS HT'). Ce traitement thermique a été développé pour améliorer le renforcement GB qui se produit généralement à haute température avec des conditions de faible contrainte. Grâce à ce traitement thermique, la structure colonnaire de l'échantillon construit en As a été détruite et tous les grains ont été équiaxés. De plus, des GB dentelés ont été produits avec succès et la plupart des GB ont été décorés avec des carbures NbC. En outre, un traitement thermique de recristallisation (RX HT) a également été effectué pour induire la recristallisation et un état de solution solide approprié ayant une valeur de dureté élevée. Ces deux échantillons traités thermiquement ont été soumis à des essais de traction à haute température avec des échantillons de L-PBF tel que construit (AB) et d'alliage corroyé 625. Les essais de traction ont été réalisés à température ambiante, 500 ° C, 600 ° C et 700 ° C dans les mêmes conditions dans l'air. Un essai de traction à température ambiante a été effectué pour vérifier la fiabilité de chaque échantillon et ses résultats peuvent être comparés à l'étude précédente de en ce qui concerne les conditions de fabrication de l'échantillon. Les graphiques de traction des deux échantillons traités thermiquement étaient situés entre ceux des échantillons AB et ouvrés. Ainsi, leurs forces étaient plus élevées que celles de l'échantillon forgé, et ils avaient des valeurs de ductilité plus élevées que l'échantillon AB. Ces résultats peuvent être attribués aux différences de savoir si la microstructure initiale est colonnaire ou équiaxe, s'il y a des GB droits ou des GB dentelés, et s'il y a ou non des précipitations. Comme dans une étude précédente dans laquelle des essais de traction ont été effectués à température ambiante, les trois échantillons (forgé, RX HT, GBS HT) ayant une structure de grains équiaxes ont montré une ductilité améliorée. D'autre part, les trois autres échantillons (AB, RX HT, GBS HT) ayant des précipités ont montré des valeurs de résistance améliorées. Bien que les traitements thermiques n'aient pas pu améliorer la ductilité autant que l'échantillon forgé, le phénomène DSA, qui montre une propriété de traction inattendue à une vitesse de déformation et une température typiques, a rarement été observé dans l'échantillon de GBS HT. Par conséquent, l'échantillon de GBS HT semble avoir un bon effet sur une partie de l'alliage 625 fabriqué au L-PBF.

Oxidation of NiCrAlY coating deposited by selective laser melting for thermal barrier coating application

La fabrication du L-PBF est considérée comme particulièrement appropriée pour les pièces à géométrie complexe avec une microstructure multicouche, y compris des systèmes de barrière thermique et / ou de revêtement environnemental, pour satisfaire

II. Introduction

Ni-based superalloy 625 has been widely used for submarine, aerospace, and gas turbine applications because of its good high-temperature strength, creep resistance, corrosion resistance, and excellent weldability. However, its good mechanical properties contribute to difficulties in machining, which have led to the interest in application of this alloy in the additive manufacturing (AM) field. AM technology has many advantages. The greatest advantage is that it can be used to manufacture complex shapes without tooling; design for additive manufacturing (DfAM) which means there are no constraints in designing of parts. This characteristic of AM is attractive mostly in the aerospace field because it can reduce the weight of engine parts by eliminating unnecessary materials in some locations, that would be required for easy fabrication by conventional tooling and also reduce wear damage due to part contact. In addition to the aerospace field, the application of AM technology has spread to other industries, and interest in this technology is increasing. The basic study of fabrication skills has continued, but trying to applying AM technology in new contexts is becoming more common now, such as the processing of multi-materials, like steel/copper alloys [START_REF] Liu | Interfacial characterization of SLM parts in multi-material processing: metallurgical diffusion between 316L stainless steel and C18400 copper alloy[END_REF][START_REF] Chen | Interfacial microstructure and mechanical properties of 316L/CuSn10 multi-material bimetallic structure fabricated by selective laser melting[END_REF][START_REF] Sing | Interfacial characterization of SLM parts in multi-material processing: intermetallic phase formation between AlSi10Mg and C18400 copper alloy[END_REF].

The present study assessed the feasibility of applying AM-fabricated, especially laser powder bed fusion (L-PBF), Ni-based Alloy 625 under real industrial conditions by applying post-heat treatment and to develop Alloy 625-based novel multi-materials.

First, with respect to the mechanical properties, the ambient and high-temperature tensile properties were evaluated, and various heat treatments were suggested to improve the alloy's properties. The L-PBF fabricated as-built sample was verified to have sufficiently high strength; however, it had the potential to be fractured more quickly than being desired due to the low toughness and defects due to nonhomogeneous cooling. Therefore, various post-heat treatments were applied to improve its room-temperature and high-temperature mechanical properties. With all of these heat treatments, it was possible to produce samples fabricated by L-PBF that had

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microstructures similar to that of a wrought one, which is the preferred microstructure, composed of polycrystalline grains and precipitates. They showed improved ductility and maintained reasonable strength. Although the heat-treated microstructure was much like the wrought microstructure, its grain boundaries (GBs) were typically weaker at higher temperatures. Therefore grain boundary serration (GBS) heat treatment was suggested.

GBS is a spontaneous reaction in which wavy GBs are formed; the reaction can be triggered by specifically designed heat treatments. It has been demonstrated that it is effective in inhibiting crack propagation, especially during creep of primarily solid solution-strengthened Ni-based superalloys. A specific heat treatment, including slow cooling between solution and aging temperatures successfully led to GBS in previous studies [START_REF] Yoon | Influence of initial microstructure on creep deformation behaviors and facture characteristics of Haynes 230 superalloy at 900 °C[END_REF][START_REF] Lee | A new approach to strengthen grain boundaries for creep improvement of a Ni-Cr-Co-Mo superalloy at 950 °C[END_REF]. In addition, the mechanism of GBS was studied to understand it in depth so that it could be applied to other alloy systems. Based on this study, the introduction of GBS in wrought Alloy 625 was achieved. GB serrated wrought Alloy 625 showed zigzag patterns of GBs with M 23 C 6 and NbC carbides at the GBs. In further study, this heat treatment was applied to L-PBF produced Alloy 625. The heat treatment was adapted for application to the L-PBF Alloy 625 considering its peculiar microstructure.

Taking the predominant role of microstructure into account, the room-temperature as well as high-temperature tensile properties of samples produced by each heat treatment were compared.

Another method has been suggested to enhance the properties of superalloys and increase their operating temperatures. Surface coatings have been developed to protect substrates from oxidation, corrosion, and thermal damages at high temperature so that their mechanical properties can be maintained for longer operation [START_REF] Daroonparvar | The role of formation of continues thermally grown oxide layer on the nanostructured NiCrAlY bond coat during thermal exposure in air[END_REF][START_REF] Leng | Microstructural evolution of an aluminide coating on alloy 625 during wet air exposure at 900 °C and 1000 °C[END_REF][START_REF] Khalid | Microstructural study on oxidation of aluminized coating on Inconel 625[END_REF][START_REF] Schike | Advanced gas turbine materials and coatings[END_REF]. Typically, thermal barrier coatings (TBC) have an insulation effect, which decreases the temperature that a superalloy is actually subjected to, which permits the superalloy to withstand operation at temperatures about 150 °C above their usual upper limit [START_REF] Clarke | Thermal-barrier coatings for more efficient gas-turbine engines[END_REF][START_REF] Donachie | Superalloys: A Technical Guide 2nd[END_REF][START_REF] Reed | The Superalloys: Fundamentals and Applications[END_REF][START_REF] Padture | Thermal barrier coatings for gas-turbine engine applications[END_REF][START_REF] Osorio | Thermal barrier coatings for gas turbine applications: failure mechanisms and key microstructural features[END_REF]. These TBCs commonly consist of an external ceramic top coat (TC) for thermal insulation and an intermediate MCrAlY metallic bond coat (BC, where typically M = Ni, Co, Fe or a combination of those) to accommodate the difference between the coefficients of thermal expansion between TC and superalloy substrate, and to grow a thin, compact, and adherent protective oxide scale promoting optimized hot corrosion resistance [START_REF] Clarke | Thermal-barrier coatings for more efficient gas-turbine engines[END_REF][START_REF] Donachie | Superalloys: A Technical Guide 2nd[END_REF][START_REF] Reed | The Superalloys: Fundamentals and Applications[END_REF][START_REF] Padture | Thermal barrier coatings for gas-turbine engine applications[END_REF][START_REF] Osorio | Thermal barrier coatings for gas turbine applications: failure mechanisms and key microstructural features[END_REF][START_REF] Zotov | Effects of annealing on the microstructure and the mechanical properties of EB-PVD thermal barrier coatings[END_REF]. The MCrAlY alloy can be used as a bond coat material and as an overlay coating material itself. The thickness of MCrAlY overlay coating is unlimited, its resistance to oxidation/corrosion is proportional to its thickness. Thus, MCrAlY material is generally required on a superalloy substrate being exposed to harsh environments for both reasons.

The coating process is totally different from that of superalloy substrate production; the coating layers are generally coated on the substrate by air or vacuum plasma spray (APS/VPS), electron beam-physical vapor deposition (EB-PVD), and arc ion plating (AIP) [START_REF] Donachie | Superalloys: A Technical Guide 2nd[END_REF][START_REF] Reed | The Superalloys: Fundamentals and Applications[END_REF]. The separate processes of coated alloy manufacturing require considerable manpower, cost, and time. For these considerations, L-PBF of an NiCrAlY bond coat on an Alloy 625 substrate, itself produced by L-PBF, is considered as well. Although there have been some studies on multi-material manufacturing by AM technology, this is the first time, to the best of our knowledge, that L-PBF has been used to produce both a superalloy substrate and an MCrAlY bond coat. We believe that this method can improve the bonding between the substrate and coating due to the nature of the L-PBF process, which involves the typical remelting of underlying layers. The ultimate production of complex coated geometries, such as blades, is challenging because PBF technologies require the deposition of an entire layer of powder, thus limiting deposition on flat surfaces. However, it is believed that this issue could be addressed soon by improving the machine part like the freedom of laser scanning motion, using multiple-powder tank, etc. In this work, the optimization of the L-PBF process parameters for MCrAlY is described on the basis of several criteria, and the as-produced materials are thoroughly analyzed. Additionally, the oxidation properties of the optimally MCrAlY-coated Alloy 625 were assessed by thermal gravimetric analysis (TGA). Furthermore thermal resistance properties, especially resistance to rapid thermal expansion and contraction, were assessed by thermal shock testing. The results demonstrated the feasibility of MCrAlY coating on the superalloy substrate by means of L-PBF, and its function was found to be reasonable.

Therefore, we conclude that post-heat treatment and L-PBF manufactured MCrAlY coating improve the mechanical properties of Alloy 625 manufactured by L-PBF at real operating temperatures (~982 °C) as an alternative to conventional Alloy 625.

Furthermore, the positive effects of these procedures were demonstrated based on the microstructure and functional aspects of the fabricated samples.

19

III. Literature review

Additive manufacturing

AM is "a process of joining materials to make objects from 3-dimensional (3D) model data, usually layer upon layer, as opposed to subtractive manufacturing methodologies" [START_REF] Alcisto | Tensile Properties and Microstructures of Laser-Formed Ti-6Al-4V[END_REF], which means it can be used for all kinds of production, maintaining and conserving its required properties. Moreover, AM can achieve the manufacture of complex and near net shape without tooling. This advantage would lead to saving of material and reduction of the product weight. Hence, this technology is identified as a key factor leading the 4th industrial revolution. On the other hand, it also has some disadvantages, such as high investment costs (for metals), low volume production rate, expensive raw materials, and post-processing. Despite these problems, the revenues generated by the AM technique are growing every year. According to one report [START_REF]Increasing adoption of Additive Manufacturing in automotive and aerospace industry is[END_REF], the global AM market is expected to reach 41.59 billion dollars or more by 2027. As seen in Figure 1, the AM market will continue to grow and, accordingly, prospects for its growth are evident in the material sector as well as in services related to AM. Ti-alloys have been the most studied materials used in AM, as seen in Figure 2, along with Ni-based superalloys especially Alloy 718 and 625 due to their excellent weldability. Although nickel-based superalloys are known to be widely used for hightemperature applications due to their excellent high-temperature mechanical properties, literature reviews have shown that AM-fabricated Ni-based superalloys are mainly based on microstructure analysis or research on room-temperature mechanical properties according to parameter condition differences (Fig. 3). Therefore, it is believed that more AM research cases are needed for the real application of these alloys.

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Families of additive manufacturing technologies

AM technologies are broadly divided into seven families (Fig. 4). These families use different manufacturing methods or heat sources. In the binder jetting method, an inkjet printer head is used to spread selectively a liquid binder onto a powder bed to glue particles together and create the desired geometry layer by layer; subsequently, heat treatment is applied to remove binder and to sinter the particles together the particles together [START_REF] Zenou | 3-Additive manufacturing of metallic materials[END_REF]. Direct energy deposition (DED) is a method in which melt metal feedstock (either powder or a wire) is projected with directed continuous coaxial nozzles to the melt pool created on the part by the beam of a high energy source (laser or electron beam) [START_REF] Carroll | Anisotropic tensile behavior of Ti-6Al-4V components fabricated with directed energy deposition additive manufacturing[END_REF]. Material extrusion additive manufacturing, whose representative method is fused deposition modeling (FDM), uses a continuous filament of a thermoplastic material as a base material. A coil feeds the filament via a heated moving extruder head, and the molten material is deposited onto a platform [START_REF]3D experience, 3D printing-additive[END_REF]. Material jetting is a process in which layers of a photopolymer or a wax are selectively deposited using either a continuous or drop on demand (DOD) approach onto a build platform via inkjet printer heads. It allows forms to be manufactured with

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different materials within the same part [START_REF]3D experience, 3D printing-additive[END_REF][START_REF] Moore | Fatigue properties of parts printed by PolyJet material jetting[END_REF]. The PBF method employs highpower laser/electron beams focused on a metal powder bed [START_REF]3D experience, 3D printing-additive[END_REF][START_REF] Hou | Overview of additive manufacturing process[END_REF]. Sheet lamination uses welded, glued, or assembled sheets of a material that are cut away to form the desired object using a digitally guided laser [START_REF] Hou | Overview of additive manufacturing process[END_REF]. In vat polymerization, a vat of liquid photopolymer is used to constructs a model layer by layer through a curing process by ultraviolet (UV) light [START_REF] Hou | Overview of additive manufacturing process[END_REF]. Among various AM methods, PBF and DED are usually employed in industry where deal with metallic materials because of their suitability for processing metallic materials, in the form of powder [START_REF]Additive manufacturing (A game changer for the manufacturing industry[END_REF].

Additive manufacturing for metallic components

The DED technique, also referred to as laser-engineered net shaping (LENS), direct laser metal deposition (DLMD), and laser cladding is carried out by feeding the metal powder into coaxially operating nozzles with a laser beam directed onto the part surface [START_REF] Liu | Fabrication of functionally gaded TiC/Ti composites by Laser Engineered Net Shaping[END_REF] as shown in Fig. 5. It has the advantage that it can be employed for diverse material processing activities, such as metallic coating, high-value component repair, prototyping, and even low-volume manufacturing with a closed-loop control 23 system to maintain dimensional accuracy and material integrity. Moreover, it can be used to quickly create large parts [START_REF] Shim | Effect of layer thickness setting on deposition characteristics in direct energy deposition (DED) process[END_REF][START_REF] Javidani | Additive Manufacturing of AlSi10Mg Alloy Using Direct Energy Deposition: Microstructure and Hardness Characterization[END_REF]. This method can be used for repairing and coating due to the nature of the technique; a flat starting base plate is not necessary [START_REF] Zenou | 3-Additive manufacturing of metallic materials[END_REF]. Bi-material manufacturing by DED has been already studied in relation to two different kinds of Ti alloys (TC11/Ti2AlNb) [START_REF] Zhang | The interface microstructure and tensile properties of direct energy deposited TC11/Ti2AlNb dual alloy[END_REF], pure copper (Cu), stainless steel 304L [START_REF] Ahang | Joining of copper and stainless steel 304L using direct metal deposition[END_REF], molybdenum alloy [START_REF] Bhardwaj | Direct Energy Deposition-Laser Additive Manufacturing of Titanim-Molybdenum alloy Parametric studies, Thesis for the Degree of Doctor Philosophy in CWNT & IMT Mines-Albi 94 microstructure and mechanical properties[END_REF], and other materials. However, it has significant disadvantage of a rough surface finish, due to the large width of the deposited bead.

Powder DED results in a better surface finish than the wire DED technique, with typical XY resolutions of 100 to 1,000 µm and 2,000 to 50,000 µm, respectively [START_REF]Digital Alloys' Guide to Metal Additive Manufacturing-Part11[END_REF]. On the other hand, the PBF process produces a much better surface finish (higher resolution) than the DED method, with a resolution of 20 to 200 µm [START_REF]Digital Alloys' Guide to Metal Additive Manufacturing-Part11[END_REF]. This technology uses a small-diameter of heat source as laser or electron beam [START_REF] Gasman | 15-Additive aerospace considered as a business[END_REF] to selectively melt thin layers of powders (L-PBF: 20 μm -100 μm, EBM: 50 μm -200 μm). It can build parts with complex geometries but not as large as those that can be built by DED. Generally, the recoating system is spreading over a thin layer of metallic powder, which is commonly made up of the same material as that of the powder bed, and selected areas are fused together layer by layer [START_REF] Zenou | 3-Additive manufacturing of metallic materials[END_REF]. There are three representative PBF-type techniques, selective laser sintering (SLS), electron beam melting (EBM), and laser powder bed fusion (L-PBF,) roughly depending on their type of beam used, the power of the beam, and the melting ratio, as shown in the Table 1. [START_REF] Serin | A review of additive manufacturing technologies[END_REF].

Table 1. Comparison of characteristics of SLS, EBM and L-PBF

SLS EBM L-PBF

Type of Beam Laser beam Electron beam Laser beam

Melting Ratio Partial Full Full

Material

All (especially for polymers)

All (especially for Ti-alloy) All

The SLS differs from EBM and L-PBF in that this method is used for rapidprototyping due to low volume density of parts produced. Moreover, metallic powder is not commonly used in SLS currently [START_REF] Hongjun | A note on rapid manufacturing process of metallic parts based on SLS plastic prototype[END_REF]. The major difference between EBM and L-PBF is the heat source; power is supplied by transferring kinetic energy (electron beam) and photon absorption (laser), respectively [START_REF] Zenou | 3-Additive manufacturing of metallic materials[END_REF]. The power of the energy source of EBM can be about 10 times higher than L-PBF [START_REF] Wysocki | Laser and Electron Beam Additive Manufacturing Methods of Fabricating Titanim Bone Implants[END_REF]. Because of the high power input to the thin metallic powder bed, the officially known cooling rate is up to 10 6 K in both cases (or higher in L-PBF since it does not have a pre-scanning system).

This rapid cooling rate can raise the residual stresses in the part interior, causing detachment from the base plate or cracking at the edge side. To reduce the cooling rate, heating of the powder bed is conducted as follows. The temperature should be kept higher than 600 °C in EBM by pre-scanning of every layer prior to melting. In L-PBF, the sole base plate is heated up to 200-500 °C [START_REF] Gokuldoss | Additive Manufacturing Processes: Selective Laser Melting, Electron Beam Melting and Binder Jetting-Selection Guidelines[END_REF] accordingly, the supporters connecting the base plate and the part surface, which is expected to be subject to high stress, are heated together. Furthermore, the beam power difference requires the appropriate adjustment of the powder size to be melted well, thus minimizing the porosity of the volume under the given power of the scan. The powder should be in the size range of 15-45 µm and 45-105 µm in L-PBF and EBM, respectively [START_REF] Zenou | 3-Additive manufacturing of metallic materials[END_REF]. In L-PBF, the powder directly absorbs the heat energy from the laser and melts, whereas in EBM, the powder is penetrated by an electron beam that produces heat energy from the kinetic energy transfer, causing the particles to melt and adhere to each other [START_REF] Fousová | Influence of Inherent Surface and Internal Defects on Mechanical Properties of Additively Manufactured Ti6Al4V Alloy: Comparison between Selective Laser Melting and Electron Beam Melting[END_REF].

Because the large beam spot size requires a larger powder size, it can produce larger layer thickness resulting in higher production rate than L-PBF while it has a rougher surface finish. The characterization results obtained for samples produced by the two methods are presented in Table 2. Schematics of both methods are shown in Fig. 6.

Materials

Introduction of Ni-based superalloys

A superalloy is an alloy that can withstand operating temperatures usually over 500 °C. Such an alloy retains good mechanical properties at a high temperature that is over 70% of their melting temperature compared to conventional structural alloys [START_REF] Yoo | Nickel Base Superalloys[END_REF][START_REF] Byun | Research Trend of Mo based Superalloys[END_REF]. These superalloys are made of group VIIIB elements of the periodic table with relatively low reaction rates for mechanical strength and oxidative improvement [START_REF] Sims | Superalloys Ⅱ[END_REF].

Therefore, most superalloys are employed for power generation, aero-engines, exhaust valves, steam turbines, and so forth.

Strengthening mechanisms of Ni-based superalloys

The 4. The precipitation strengthening mechanism is divided into two, 'Orowan bypassing'

and 'shearing' depending on interaction between precipitates and dislocations.

Shearing mechanism usually occurs when size of the coherent precipitate with matrix is small. The APB and SF take place during dislocation shearing. However incoherent parts are rather subject to bow-out by the Orowan mechanism as size of precipitates increase, dislocations left rings around a precipitate since it is difficult for a dislocation to overcome [START_REF] Hosford | Mechanical Behavior of Materials[END_REF][START_REF] Kelly | Strengthening methods in crystals[END_REF]. Each precipitation contributing to alloy strengthening has critical size (r 0 ) where the mechanism changes from shearing to Orowan bypassing as indicated in Figure 8. It is characterized by an elliptical shape aligning along the <110> direction of the matrix. In addition, because it is a metastable phase, if it is exposed for a long time at the temperature of 650 °C or higher, the phase transforms to an acicular δ-phase its formation rate increases when the material is stressed [START_REF] Kim | Creep oxidation behavior and creep strength prediction for Alloy 617[END_REF]. Table 5 summarizes the composition and crystal structure of the above-described phases and some other phases that may be formed in Ni-based superalloys.

Introduction of Alloy 625

Alloy 625 was invented around 1964 for application in high-pressure steam lines in power plants, and its application in industry was gradually broadened due to its robustness in harsh environments. As presented in Figure 9, this alloy was originally developed from Alloy 600 to improve the corrosion properties; it can endure highly acidic environments, marine conditions, and high temperature conditions (from cryogenic to 980 °C in theory) [53]. In addition to these properties, highly alloyed Alloy 625 can provide high strength in harsh environments. Its nominal composition is presented in Table 6. Mo has the solid-solution-strengthening effect on its matrix with Fe and Nb. In addition, it contributes to the good corrosion resistance together with high level of Cr content. Al and Ti are added in small amounts mainly for refining purposes to enhance weldability by suppression of the γ´ phase in comparison to other alloys, such as Alloy 718, while a sufficient Nb+Ti+Al content has a positive effect on precipitation hardening owing to the presence of the γ´´ phase [START_REF][END_REF].

On the other hand, large amounts of alloy elements do not always have good effects. When this alloy is solidified from the liquid state, the dominating solidification reaction is the enrichment of the remaining interdendritic liquid in Nb, and the consequent formation of Nb-rich laves phase and/or Nb carbide during the final stages of solidification. The phase formation during solidification is governed by the C/Nb ratio; the formation of γ+NbC occurs with no laves phase at a high C/Nb ratio. First γ+NbC is formed, followed by laves phase formation at the end of solidification at an intermediate C/Nb ratio; γ+laves is formed with no NbC at a low C/Nb ratio. The primary NbC carbide that appears when it is solidified is normally very large, and it can be serve as a crack initiation site in the presence of an external load. Thus, it is preferable that this compound be dissolved during the solid-solution process to the greatest extent possible. respectively [57]. γ˝ is a main strengthening precipitate of this alloy whose strengthening mechanism comes from having coherence with the γ-FCC matrix. This precipitate presents an elliptical disc shape from the tetragonal crystal structure. Its strengthening mechanism is, analogous to γ´, formation of APB or SF between two partial dislocations when a dislocation tries to pass through this particle. Formation of γ˝ requires aging treatment in a temperature range of 600 to 650 °C for over 20 hours, whereas it can be formed by short-term aging when it is aged at higher temperature about 650 to 750 °C [START_REF] Sundararamman | Precipitation of an intermetallic phase with Pt2Mo-type structure in alloy 625[END_REF]. This phase is thermally unstable; it ultimately changes to equilibrium phase δ. The δ orthorhombic incoherent phase presents an acicular shape.

It can be formed by transformation from the γ˝ phase or directly by aging treatment [START_REF] Mathew | Microstructural changes in alloy 625 during high temperature creep[END_REF].

L-PBF manufacturing of Alloy 625

Alloy 625 is one of the most widely used Ni-based superalloys along with 718 for

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AM fabrication by PBF. This alloy, in particular, is usually used for the manufacture of parts through the PBF method. On the surface, unmelted powder particles are often observed because of the low chance of re-melting which is something in common of most additive manufactured part. In addition, for samples with curved or sloped surfaces, due to the AM nature of layer-by-layer scanning, step marks appear on the surface called staircase effect requiring proper surface treatment according to the application environment of the sample. The L-PBF method results in relatively better surface properties because of the smaller melt pool size from the lower heat source compared to EBM.

Alloy 625 produced by L-PBF varies slightly depending on the processing conditions, but generally Nb and C are segregated on the liquid/solid interface during manufacture, resulting in NbC carbides. However, because this is caused by rapid heating/cooling, segregation in the form of a very thin film appears, and the resulting carbide size is as small as submicron size. This has a better effect on the mechanical properties than a large primary NbC carbides found in wrought Alloy 625, which is likely to be a crack initiation site, which appears during standard solidification. In addition, high density of dislocations generated by rapid heating/cooling contributes to high strength together with fine NbC. A high density of dislocation is observed in columnar arrangements according to the scanning direction; NbC of submicro size is mainly observed at the columnar interface (figure 11). Alloy 625 made by L-PBF is known to begin to recrystallize after heat treatment at temperatures higher than 1000 °C, and heat treatment alone can lead to recrystallization because of the high density of dislocations generated during the manufacturing process. With a high density of dislocations, the as-built sample has very high strength with low elongation capacity; thus, it requires heat treatment to recover some ductility. Accordingly, several studies have been conducted to compare the mechanical properties of Alloy 625 after heat treatment, and the findings are summarized in the table below. All of the above mentioned studies were conducted at room temperature; very few studies have been conducted on the mechanical properties of the alloy at high temperatures, which are the conditions under which this alloy is actually used. Furthermore, most of the studies were performed with L-PBF as-built without any microstructure control through post-heat treatment.

Introduction of coatings on superalloys

Superalloys are generally used in environments in which oxidation and hot corrosion can easily occur, such as high temperature and seawater applications. In addition to this, degradation caused by oxidation tends to be worse with increased temperature. However, application at very high temperatures is required with the development of technologies. In addition, it is desirable for superalloys to keep their mechanical properties as long as possible in harsh environments [START_REF] Reed | The Superalloys Fundamentals and Applications[END_REF].

For the long-term performance of the alloy, degradation due to external factors must be minimized. Coating on the alloy surface is essential to maintain its mechanical properties. Surface coating improves resistance to oxidation/corrosion, and reduction of temperature enables the alloy to operate above its incipient melting temperature. A ceramic material used as a coating is able to drop the surface temperature due to its low thermal conductivity [START_REF] Muktinutalapati | Materials for gas turbines-An overview[END_REF]. Some coating possesses an elements reservoir for the formation of a protective dense oxide layer at the interface between the coating surface and air. The elements react with detrimental factors attacking alloys to protect them [START_REF] Tawancy | Failure mechanism of a thermal barrier coating system on a nickel-base superalloy[END_REF]. There are three types of coatings generally used in industry; diffusion coatings, overlay coatings, and thermal barrier coatings (TBCs) as shown in Figure 12.

Diffusion coatings are most commonly used; they use at least one of the major constituents from the base metal by diffusion for the coating property. This type of coating is composed of an external diffusion zone and an internal diffusion zone. An external diffusion zone is developed by reaction with the elements in the alloy substrate, and the internal diffusion zone is a mixed zone of the substrate and coating elements. Diffusion of Cr, Si, and other elements can be used for the prevention of Type Ⅱ corrosion (650 to 800 °C), and Al diffusion with Pt, Pd, and/or rare earth elements can contribute powerful resistance to Type Ⅰ corrosion (800 to 950 °C).

However, it has an oxidation resistance limit, which led to greater interest in overlay

Novel Fabrication of Alloy 625 and MCrAlY Bond Coat by L-PBF and Microstructure Control

41

coatings [START_REF] An | Effect of the θ-α-Al2O3 Transformation in Scales on the Oxidation Behavior of a Nickel-Base Superalloy with an Aluminide Diffusion Coating[END_REF].

Overlay coating, also called MCrAlY coating (M stands for Ni or Co or Ni+Co), has independence from the substrate elements as opposed to diffusion coatings. It supplies all the constituents in the coating itself. This method has the advantage that it provides resistance to both Type Ⅰ and Type Ⅱ corrosion and oxidation. The coating thickness is not limited by process considerations, and the thicker the coating is, the better oxidation/corrosion properties it can have. High-temperature heat treatment is required to homogenize the coating and ensure its adherence to the substrate [START_REF] Taylor | An oxidation study of an MCrAlY overlay coating[END_REF]. The coated MCrAlY shows four types of representative phases as seen in Table 8, namely, γ, β-NiAl, γ´, and α-Cr. Based on the predicted phase map (Figure 13) of NiCrAl (typically 22 wt% Cr and 10 wt% Al, bal. Ni, the very low amount of Y was not taken into account), γ and β-NiAl phases appear at high temperature, and they would transform to γ´ and α-Cr phases around 975 °C having volumetric change. This volumetric change may lead to internal stress within the coating part; therefore, care must be taken when it is subjected to thermal treatment [START_REF] Richard | The influences of heat treatments and interdiffusion on the adhesion of plasma-sprayed NiCrAlY coatings[END_REF][START_REF] Fritscher | Investigation of an as-sprayed NiCoCrAlY overlay coating: A thermoanalytical approach[END_REF]. Meanwhile, the coating property depends on the diffusion rates of active metal elements and oxygen atoms. If the diffusion rate of active metal elements, especially Al and Cr, is faster than that of oxygen atoms attacking the alloy, external oxidation, TGO, takes place at the surface of the coating part. Otherwise, if the reverse is true, internal oxidation may take place at the interface of the substrate and coating. This internal oxidation is highly undesirable because it does not protect the alloy substrate from the external environment, and it generates volumetric change, which gives rise to tension stresses at the surface, which causes susceptibility to thermomechanical fatigue and the development of an inhomogeneous surface oxide layer [START_REF] Villafane | Oxidacion en alta temperatura[END_REF]. The present study, to improve the high-temperature properties of Alloy 625 produced by L-PBF, a multi-material was fabricated by bonding MCrAlY to the alloy surface in the same way. Recent research on the production of multi-materials by AM includes studies on stainless steel and copper alloys [START_REF] Liu | Interfacial characterization of SLM parts in multi-material processing: metallurgical diffusion between 316L stainless steel and C18400 copper alloy[END_REF][START_REF] Chen | Interfacial microstructure and mechanical properties of 316L/CuSn10 multi-material bimetallic structure fabricated by selective laser melting[END_REF][START_REF] Sing | Interfacial characterization of SLM parts in multi-material processing: intermetallic phase formation between AlSi10Mg and C18400 copper alloy[END_REF], a composite material mixing two different powders [START_REF] Demir | Multi-material selective laser melting of Fe/Al-12Si components[END_REF][START_REF] Wei | 3D printing of multiple metallic materials via modified selective laser melting[END_REF], NiCrAlY cladding [START_REF] Partes | High temperature behavior of NiCrAlY coatings made by laser cladding[END_REF][START_REF] Tobar | Laser cladding of MCrAlY coatings on stainless steel[END_REF][START_REF] Vilar | Structure of NiCrAlY coatings deposited on single-crystal alloy turbine blade material by laser cladding[END_REF], a developed 45 methodology based on the selective laser melting (SLM) concept [START_REF] Chivel | New approach to multi-material processing in selective laser melting[END_REF] innovative AM of a steel-ceramic multi-material [START_REF] Koopmann | Additive manufacturing of a steel-ceramic multi-material by selective laser melting[END_REF] among others. In this work, MCrAlY was introduced for the L-PBF production of MCrAlY coated Alloy 625 for the first time, to the best of our knowledge; by verifying its validity, it will be able to take the first step to contribute more efficiently to the production of coating alloys in the industrial field, and further lay the groundwork for future equipment improvement and research directions in the AM field.

Introduction of MCrAlY coatings

Grain boundary strengthening of γ matrix

GBs play a key role when a sample is subjected to creep condition. Generally, high-angle GBs are very weak at high temperatures due to the GB sliding effect, especially under creep condition. Therefore, single-crystal alloys which have no grain boundaries being produced by directional solidification method are preferred in very high-temperature applications, such as turbine blade. Each of their brief microstructure is indicated in Figure 15. However, the particular manufacturing process of singlecrystal alloys has led to interest in GB strengthening of polycrystalline alloys due to its difficult manufacturing process. Therefore, two representative grain-boundarystrengthening methods were described in the following section.

Thesis for the Degree of Doctor Philosophy in CWNT & IMT Mines-Albi

Improvement of GB cohesion strength in Ni-based superalloy

There are various ways to achieve GB strengthening by improving GB cohesion.

The first method is adding GB-strengthening elements, which are generally known elements, such as B, C, Zr, Hf, and Y. These elements in small quantities can improve the mechanical properties of alloys significantly due to their effects on GBs. They tend to move to the GBs on account of their different atomic sizes and electronic properties, which results in GB stabilization, providing high cohesive strength [START_REF] Sims | Superalloys Ⅱ[END_REF][START_REF] Brooks | Heat Treatment, Sturcture and Properties of Nonferrous Alloys[END_REF]. The effect of each element has been studied. B is good for ductility and having a highly strengthening influence with Hf, whereas C has little or no effect. Regarding Zr, it has been reported that it shows a slight embrittlement effect. Moreover, regarding their contribution to the high-temperature mechanical properties, they inhibit the formation of GB carbides, such as M 23 C 6 and M 6 C, which are brittle when they are coarsened and have the potential to initiate cracks easily when the Ni-based superalloy is exposed to high temperatures [START_REF] Razumovskiy | First-principles-aided design of a new Ni-base superalloy: Influence of transition metal alloying elements on grain boundary and bulk cohesion[END_REF].

In the second method, the proportion of special GBs is increased, known as GB engineering (GBE), typically for the face-centered cubic (FCC) system. A special GB is a low-index boundary plane or low-coincidence site lattice (CSL) plane like twin boundaries (Σ3) that have resistance to intergranular degradation due to their very low disorientation and stable configuration of the interface. Its mechanism is explained by the term 'Σ3 regeneration' based on Σ3 n + Σ3 n+1 →Σ3. This method requires a threshold level of Σ3s and a driving force for selective GB (mainly incoherent Σ 3s and some Σ 9s) migration. It can be achieved by generating a proper amount of annealing twins and replacing the random boundary (high CSL) with the incoherent Σ3s [START_REF] Randle | Grain boundary engineering: an overview after 25 years[END_REF].

GB serration is another method to improve GB cohesive strength. The concept of the GB serration method is similar to GBE, which produces special boundaries.

Through this method, GBs are produced with a wavy configuration via specialized heat treatment. Although the total area of the GB plane is increased, so it is not expected to have a positive effect on grain boundary sliding occurring condition, it has been proved by some studies that serrated GBs are arranged along the low index plane (close to {111}) of each neighboring grain alternately. Thus, the alloy was highly resistant to crack initiation and propagation. This method is based on the GB phasetransition theory suggested by Hsieh et al. who observed GBs with a flat and smoothly curved interface with temperature change [START_REF] Hsieh | Observations of roughening/de-faceting phase transitions in grain boundaries[END_REF].

GB serration

Grain boundary serration (GBS) refers to wavy grain boundaries (GBs) having an amplitude and a period, usually occurring during cooling after homogenization heat

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treatment [START_REF] Koul | On the mechanism of serrated grain boundary formation in Ni-based superalloys[END_REF][START_REF] Larson | Metallurgical factors affecting the crack growth resistance of a superalloy[END_REF]. When a Ni-based superalloy is under creep condition, cracks initiate and propagate along the GBs, which are vulnerable at high temperature, while the serrated GBs make the sample resistant to the crack propagation due to its specific GB configuration; therefore, creep life is extended by improvement of creep rupture strength [START_REF] Larson | Metallurgical factors affecting the crack growth resistance of a superalloy[END_REF].

A study to figure out the mechanism of GBS is still ongoing. The mechanism for generating GBS [START_REF] Koul | On the mechanism of serrated grain boundary formation in Ni-based superalloys[END_REF][START_REF] Larson | Metallurgical factors affecting the crack growth resistance of a superalloy[END_REF][START_REF] Henry | The dendritic growth of γ′ precipitates and grain[END_REF][START_REF] Mitchell | On the formation of serrated grain boundaries and fan type structures in an advanced polycrystalline nickel-base superalloy[END_REF] has been attributed to the interaction between GBs and precipitates, which is preferentially precipitated around GBs during slow cooling. For example, as the γ´ phase grows asymmetrically in the direction of a GB, the GB is pushed to be moved by the γ´ phase [START_REF] Henry | The dendritic growth of γ′ precipitates and grain[END_REF], the suppression of GB motion when the γ´ phase is pre-precipitated around the GB [START_REF] Danflou | Formation of Serrated Grain Boundaries and Their Effect on the Mechanical Properties in a Nickel Base Superalloy[END_REF], or the GB is led and moved as the γ´ phase moves toward the GB [START_REF] Larson | Metallurgical factors affecting the crack growth resistance of a superalloy[END_REF]. This GBS mechanism theory assumes that the γ´ phase has to be preferentially precipitated before the formation of GBS.

Hong et al. [START_REF] Yoon | Influence of initial microstructure on creep deformation behaviors and facture characteristics of Haynes 230 superalloy at 900 °C[END_REF][START_REF] Lee | A new approach to strengthen grain boundaries for creep improvement of a Ni-Cr-Co-Mo superalloy at 950 °C[END_REF] reported that high-temperature mechanical properties were improved by introducing GBS into solid-solution-strengthened alloys, namely, Alloy 230 and Alloy 617, which both have very low γ´ volume fractions. In addition, they demonstrated that GBs become serrated without any influence of inter/intra granular precipitates during slow cooling of the heat-treatment process. They also confirmed that the direction of the serrated GBs was almost {111}, which is a close-packed plane in FCC; thus, the GBs were stabilized by lowering the interface energy. According to this new theory, GBS can be introduced considering following conditions. First, the newly attempted GBS formation theory is unlike the conventional GBS formation theory [START_REF] Randle | Grain boundary engineering: an overview after 25 years[END_REF][START_REF] Koul | On the mechanism of serrated grain boundary formation in Ni-based superalloys[END_REF]; GBs are independent of precipitates, and GBS is formed preferentially by GB motion prior to GB precipitation. Therefore, for a GB moving freely, the precipitates that were inhibiting the free movement of the GB must be completely dissolved during solid-solution treatment, and precipitation should be delayed as much as possible during the process of GBS formation. Secondly, sufficient thermal activation energy should be provided to allow the GB itself to move smoothly. Finally, the GBS phenomenon is similar to phase transition, which occurs spontaneously to reduce free energy at each temperature during cooling [START_REF] Hsieh | Observations of roughening/de-faceting phase transitions in grain boundaries[END_REF][START_REF] Hong | Significant decrease in interfacial energy of grain boundary through serrated grain boundary transition[END_REF]; that is, sufficient time should be provided to maintain an equilibrium configuration of GBs at each temperature during cooling from high temperature to low temperature.

Several studies related to inducing GBS have been reported [START_REF] Yoon | Influence of initial microstructure on creep deformation behaviors and facture characteristics of Haynes 230 superalloy at 900 °C[END_REF][START_REF] Lee | A new approach to strengthen grain boundaries for creep improvement of a Ni-Cr-Co-Mo superalloy at 950 °C[END_REF][START_REF] Hong | Significant decrease in interfacial energy of grain boundary through serrated grain boundary transition[END_REF]; however, no study about the Nb-rich solid-solution strengthened Alloy 625 has been reported ever.

The target alloy has a relatively high Nb content, which easily appears in the NbCtype carbide as opposed to the previous studies [START_REF] Yoon | Influence of initial microstructure on creep deformation behaviors and facture characteristics of Haynes 230 superalloy at 900 °C[END_REF][START_REF] Lee | A new approach to strengthen grain boundaries for creep improvement of a Ni-Cr-Co-Mo superalloy at 950 °C[END_REF][START_REF] Hong | Significant decrease in interfacial energy of grain boundary through serrated grain boundary transition[END_REF]. Therefore, the above theory of GBS formation was newly introduced for the Nb-added solid-solution-strengthened Alloy 625, and in-depth analysis of the mechanism of GBS formation was conducted.

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IV. Methodology

The commonly used equipment and methods are described here, but special methods and machines are indicated in each paper (Section 4.3 Introduction of articles)

Alloy 625 and MCrAlY powders

The element composition of each powder was analyzed by Inductively coupled plasma atomic emission spectroscopy (ICP-OES), and the results are shown in Table 9.

Alloy 625 powder provided by the L-PBF machine manufacturer was used to build substrate material. Most particles were spherical with some satellite particles. The powder-size distribution was measured by using a laser particle-size analyzer. The particles had a size distribution between 20 µm (D10) and 44 µm (D90) with an average volume diameter around 29 µm. Additionally, the powder flow rate was measured by the Hall flow method (standard ASTM B 213); 50 g of powder was passed through a funnel, and the time it took was registered as the flow rate. The funnel had a 30° cone angle and a ø2.5-mm outlet hole. The results of the powder flow rate were 12.6 sec.

NiCrAlY powder provided by DUCAL Corporation composed with generally known element composition (Table 9) has same powder morphology as that of Alloy 625. The powder-size distribution resulted in between 24 µm (D10) and 51 µm (D90)

with an average volume diameter around 35 µm (Figure 16). The powder flowability, which was measured in the same manner as Alloy 625 was 17.4 sec.

Additive manufacturing

A SLM 125 HL (SLM Solutions) L-PBF machine, which has a build chamber size of 125 mm × 125 mm × 125 mm was used to manufacture all samples in this study.

The machine system is equipped with a 400-watt fiber laser as the energy source (80μm laser beam diameter). Argon was used as an inert gas in the build chamber. The stripe scanning strategy with stripe length of 10 mm and 33° orientation rotated incrementally between layers was used for Alloy 625 fabrication by adopting default parameter condition (power: 275 W, scanning speed: 760 mm/s, hatch distance: 120 µm, layer thickness: 50 µm) and non-rotation for NiCrAlY (for the oxidation test samples, rotation was exceptionally adopted with 67° to give same condition with substrate manufacturing). For the parameter condition of NiCrAlY, the optimization of record the mass evolution during the high-temperature exposure. Samples were directly hung in the furnace using a platinum suspension. For NiCrAlY-coated Alloy 625 samples, thermogravimetric analysis was performed using a NABERTHERM N 11/H furnace. The temperature was verified with a K-type thermocouple, and the mass evolution was measured using a precision balance (SARTORIUS MC5). The samples were hung in an alumina crucible with platinum wires to expose all sample faces to the environment. Interruptions at definite exposure times were carried out at 5, 10, 20, 50, 100, and 200 hours to document the evolution of the mass gain and the nature of the oxidation products.

-Thermal shock test: The machine used for the thermal shock test was a specifically designed 'Burner Rig' produced by Hysytech SRL Corporation. For each sample, 450 cycles were carried out. Considering the nature of the superalloy substrate, a maximum temperature of 950 °C (2 min plateau) and quenching to 300 °C was conducted. Heating to 950 °C took approximately 1 h, and quenching to 300 °C took under 1 min. Only methane was burnt, and the air/methane ratio was set to 1.2 to provide an oxidizing environment. The specimens were weighted with a precision balance (0.1 mg) and photographed after each 30-cycle session to study the evolution of the materials.

V. Publications and scientific contribution 5.1 Focus of research

 Manufacturing an operational NiCrAlY coated Alloy 625 bi-material system by L-PBF.

 Each property of only substrate and coated alloy has to be evaluated and improved to certify for commercialization/application.

 The microstructure of various heat treated L-PBF Alloy 625 was investigated and their tensile properties were compared at room temperature.

 For considering real operation conditions of the alloy, high temperature properties investigation was considered. Furthermore, GBS was chosen as a one way to improve the high temperature mechanical properties. Therefore, GBS mechanism is needed to be investigated for the first step to apply to Nb added Nibased superalloy (Alloy 625) for the first time.

 GBS mechanism was understood for the first time with a model alloy.

 The next step is applying GBS to wrought Alloy 625, the standard alloy.

 Application of GBS to wrought Alloy 625 was succeeded.

 Finally it was ready to apply GBS to L-PBF Alloy 625.

 GBS was successfully induced in L-PBF Alloy 625 and it's high temperature tensile properties were assessed with recrystallized L-PBF, L-PBF As-built and wrought Alloy 625. Microstructure controlled L-PBF Alloy 625 showed higher strength than wrought one and higher ductility than L-PBF As-built at high temperature. Especially GBS induced sample has resistance to DSA occurrence. determined to be at least 0.97 kJ/mol by first-principle calculation. Thus, GBS can be brought about by introducing a proper amount of atoms having a large diameter, such as Cr, Zr, and Nb, to provide enough strain energy as its trigger.

We then wondered if we could apply this GBS to L-PBF Alloy 625, thus the first trial to induce GBS of Alloy 625 was made on a wrought alloy as a standard. Based on its mechanism, firstly it was realized that precipitates have to be dissolved for employment of enough solute atoms in the Ni-Cr matrix and letting the GBs move freely, and then the solute atoms have to approach to the GBs as much as possible during slow cooling. Accordingly the starting temperature for solution heat treatment was higher than in previous study related to GBS of Alloy 617 (1200 °C→1250 °C) to dissolve as much as possible precipitations, especially primary NbC, which has a relatively high solvus temperature. The modified heat treatment introduced GBS to wrought Alloy 625 successfully. The successfully induced serrated GBs were decorated with M 23 C 6 as well as NbC. After the new attempt to promote GBS in wrought Alloy 625, it we assessed whether GBS induced in Alloy 625 would be more effective than the conventional wrought alloy for high-temperature mechanical properties and the possibility to introduce GBS to Alloy 625 fabricated by L-PBF as well.

The introduction of GBS in L-PBF manufactured Alloy 625 has been for the first time ever, the microstructure of the L-PBF Alloy 625 was firstly analyzed. First of all, it was necessary to removing precipitates as much as possible so that the GBs could move freely. L-PBF Alloy 625 is composed of a fine columnar structure with NbC at the interface of each columnar structure. This columnar structure with NbC, which has a GB pinning effect is not preferred. Thus, a higher solutionizing temperature than that of wrought Alloy 625 was chosen (1250 °C→1300 °C). Furthermore, the final aging

Introduction of articles 1. A novel approach to the production of NiCrAlY bond coat onto IN625 superalloy by selective laser melting

This study was about the feasibility of producing MCrAlY bond coat material directly on an Alloy 625 substrate produced by L-PBF. In the present study, L-PBF was used to fabricate both the Ni-based superalloy substrate and the MCrAlY bond coat, which is the first time this has been studied, to the best of our knowledge, in relation to AM fabrication. Fifteen different processing parameters were tested, and all cases showed substantial dilution between NiCrAlY and superalloy substrate due to the typical remelting characteristic of the L-PBF process, ensuring an excellent metallurgical bonding. The optimum process parameters, layer thickness t = 50 μm, hatch spacing h = 120 μm, laser power P = 250 W, and laser scanning speed v = 800 mm/s, were selected based on several criteria. An approximately 250-μm-thick NiCrAlY coating was obtained with a very low residual porosity ≤0.1%, including a ∼35% dilution zone, a smooth hardness profile between 275 and 305 Hv, low residual stress levels, and a progressive Al concentration distribution. These results confirmed the feasibility of L-PBF as an alternative method for the deposition of coatings in TBC systems or MCrAlY coating itself.

Influence of heat treatments on microstructure evolution and mechanical properties of Inconel 625 processed by laser powder bed fusion

This study was about the mechanical behavior at room temperature and microstructure of as-built and heat-treated Alloy 625 samples fabricated by L-PBF. On contrary to the rest of the work, in this case the samples were produced with different machine (EOSINT M270 Dual Mode machine, which is equipped with a 200W Yb fiber laser with a spot diameter of 100 μm) with powder (size distribution between d 30 of 16μm and d 90 of 48μm) from a different source (laser power of 195 W, scan speed of 1200 mm/s, hatching distance of 0.09mm and layer thickness of 0.02 mm) as one part of collaborated study with Politecnico di Torino in Italy. Although the L-PBF manufactured as-built sample showed good tensile properties, post-heat treatment is recommended to obtain specific mechanical properties to satisfy industrial requirements. Various types of heat treatments were performed to produce each microstructure and selected each representative heat treatment by comparison of Brinell hardness. γ˝ and M 23 C 6 carbides with recrystallized grains appeared in direct aging (DA) treated sample (700 °C/24 hrs), produced from the as-built microstructure, which showed columnar structures with high density of dislocation and NbC carbides. Solution (S) treatment (1150 °C/2 hrs) made the sample have a recrystallized equiaxed austenitic grain structure with primary and secondary MC carbides. Furthermore, the solution/aging (SA) treated sample (1150 °C/2 hrs -700 °C 24 hrs) showed the same microstructure as that of the S sample, but with γ˝ and M 23 C 6 carbides. The optimal temperature for the heat treatments was selected by a comparison of Brinell hardness values and microstructure observation. Four types of samples were subjected to tensile testing at room temperature with the as-built sample as a counterpart. The as-built sample showed yield strength (YS) and ultimate tensile strength (UTS) values of 783 MPa and 1,041 MPa respectively, with a ductility of 33%. The DA sample exhibited YS and UTS values of 1,012 MPa and 1,222 MPa, with the lowest ductility of 23%. The S sample showed the lowest YS and UTS values of 396 MPa and 883 MPa with the highest increment of ductility of 55%. The SA sample showed a significant increase in strength: YS = 722 MPa, UTS=1,116 MPa and correlate ductility of 35%. According to the tensile results, solution treatment contributes to the ductility increment, and aging treatment improves the strength. Therefore, it was demonstrated that the microstructure evolution of L-PBF Alloy 625 under various heat treatments provides a wide spectrum of possible mechanical properties.

A new observation of strain-induced grain boundary serration and its underlying mechanism in a Ni-20Cr binary model alloy

In this study, as a purpose of the GB strengthening at high temperature, the mechanism of GBS was investigated through a Ni-20Cr binary model alloy to prepare applying to additive manufactured Alloy 625 in the further study. A fully solidsolution-treated alloy was heat treated based on the previous study adapted for an Ni-Cr matrix alloy [START_REF] Yoon | Influence of initial microstructure on creep deformation behaviors and facture characteristics of Haynes 230 superalloy at 900 °C[END_REF][START_REF] Lee | A new approach to strengthen grain boundaries for creep improvement of a Ni-Cr-Co-Mo superalloy at 950 °C[END_REF][START_REF] Hong | Significant decrease in interfacial energy of grain boundary through serrated grain boundary transition[END_REF]. According to the hypothesis of the GBS mechanism, there was no occurrence of GBS anywhere as expected. However, the sign of GBS occurred by giving a 5% strain during the direct aging/slow cooling process in which GBS took place, which was conducted after the solid-solution process. The microstructure of the sample in which GBS occurred was analyzed by TEM and the 3D-APT method. There was a high density of dislocations around the GBs due to 5% straining during the aging/slow cooling process, which promoted Cr enrichment near and at the GBs. To make the sample show the same behavior just by heat treatment, the solutionizing temperature was increased to induce Cr enrichment near and at the GBs with the vacancy-assisted diffusion concept, and it worked well. The in-depth 3D-APT analysis showed that the serrated samples displayed wide Cr enrichment at the serrated GBs. The common feature of both samples was Cr enrichment, and strain on the GBs was considered as the trigger of GBS. The strain by first-principle calculation was ~0.97 kJ/mol. Thus, it was demonstrated that the minimum 0.97 kJ/mol of strain can lead to GBS by interstitial atoms at the GBs. [START_REF] Yoon | Influence of initial microstructure on creep deformation behaviors and facture characteristics of Haynes 230 superalloy at 900 °C[END_REF][START_REF] Lee | A new approach to strengthen grain boundaries for creep improvement of a Ni-Cr-Co-Mo superalloy at 950 °C[END_REF][START_REF] Hong | Significant decrease in interfacial energy of grain boundary through serrated grain boundary transition[END_REF], and significant serration was observed for the first time. M 23 C 6 carbides together with significant amount of MC carbides were observed at the GBs, whereas there was no evidence of γ˝ precipitation. The slow cooling treatment was found to effectively promote the formation of serrated grain boundaries, even in the absence of aging. Meanwhile, as the aging time is increased, precipitation of the δ phase was evident. Considering the recent study which described a tensile test performed at ambient temperature with L-PBF as-built and various heat-treated samples, modified GBS heat treatment (solid solution temperature increase from 1250→1300 °C) was applied to L-PBF Alloy 625 ('GBS HT' sample). This heat treatment was developed to improve the GB strengthening that typically occurs at high temperature with low-stress condition. Through this heat treatment, the columnar structure of As-built sample was destroyed, and all grains were equiaxed. Additionally, serrated GBs were successfully produced, and most of the GBs were decorated with NbC carbides. Furthermore, recrystallization heat treatment (RX HT) was also performed to induce the recrystallization and proper solid-solution state having a high hardness value. These two heat-treated samples were subjected to high-temperature tensile testing with L-PBF as-built (AB) and wrought Alloy 625 samples. The tensile tests were conducted at room temperature, 500 °C, 600 °C, and 700 °C under the same condition in air. A room-temperature tensile test was conducted to verify the reliability of each sample and its results can be compared with previous study of [4.3.2] in aspect of sample manufacturing conditions. The tensile graphs of the two heat-treated samples were located between those of the AB and wrought samples. Thus, their strengths were higher than that of the wrought sample, and they had higher ductility values than the AB sample. These results may be attributed to the differences in whether the initial microstructure is columnar or equiaxed, whether there are straight GBs or serrated GBs, and whether or not there is precipitation. As in a previous study in which tensile testing was conducted at room temperature, the three samples (wrought, RX HT, GBS HT) having equiaxed grain structure showed improved ductility. On the other hand, the other three samples (AB, RX HT, GBS HT) having precipitates showed improved strength values. Although heat treatments could not improve the ductility as much as the wrought sample, the DSA phenomenon, which shows an unexpected tensile property at a typical strain rate and temperature, was rarely observed in the GBS HT sample. Therefore, the GBS HT sample seemed to have a good effect on some part in the L-PBF manufactured Alloy 625. This section is based on an article which has been published in J. W. Lee, M. Terner, S. Y. Jun, E. Copin, P. Lours, H. -U. Hong, Mat. Sci. Eng. A, "Heat treatments design for superior hightemperature tensile properties of Alloy 625 produced by Selective Laser Melting" 790 (2020)

High temperature oxidation of NiCrAlY coated Alloy 625 manufactured by selective laser melting

L-PBF manufacturing is considered particularly appropriate for parts with complex geometry with a multilayered microstructure, including thermal barrier and/or environmental coating systems, to satisfy both structural and environmental specifications. From the point of environmental specifications, the high-temperature oxidation of fully L-PBF processed NiCrAlY by the previously tested and succeeded optimal parameter conditions (P = 250 W, v = 800 mm/s) [4.3.1] overlaying Alloy 625 substrate parts was investigated at temperatures between 800 °C and 1000 °C. Freestanding bulk NiCrAlY, Alloy 625, and bi-materials were fabricated using the L-PBF method. To be representative of coated components, two thin layers (about 100 μ m) of NiCrAlY powder were laser scanned at the surface of Alloy 625 to form the multi-material. The coated samples were then extracted using a gentle mechanical polishing technique to ensure repeatable thickness and surface finishing of the specimen. It is worth noting that the fabrication of L-PBF NiCrAlY was particularly difficult and resulted in macro cracks within the bulk NiCrAlY specimens due to the large volumetric residual stress, whereas the reason why micro cracks are observed at the laser bead level for coating is predicted by delayed cracking. However, it was not damaged enough to prevent the study to be done. As expected, the bulk NiCrAlY and bi-materials demonstrated an improved oxidation behavior compared to Alloy 625. Indeed, the oxidation of bulky freestanding NiCrAlY resulted in a dense and continuous Al 2 O 3 layer in this temperature range, which was what was expected for this composition of NiCrAlY. This means no evaporation effect of Al was observed (an issue observed sometimes in L-PBF). However, the formation of oxides at the surface of the bi-layer material differed from that of the bulky freestanding NiCrAlY. As a comparison, the oxidation of the bi-layer material resulted in regions with a dense and continuous Al 2 O 3 layer as well as regions composed of a mixture of external Cr 2 O 3 and internal Al 2 O 3 . Extended EDS maps at the surface of the bi-material highlighted the heterogeneous distribution of the constitutive elements of the NiCrAlY coating, resulting in some regions with Al activity lower than that required for the formation of a continuous and dense Al 2 O 3 layer. Low-Al-activity and high-Al-activity regions were related to the topography of the L-PBF surface and corresponded to peak and valley regions, respectively where it comes from the in homogeneous thickness of the powder layer deposited. Furthermore, cracks were mainly found in high-Alactivity regions and they could be attenuated by applying a stress relieving heat treatment or the diffusion. The fabrication of coated but small components with a

Thermal shock resistance of the NiCrAlY coated Alloy 625 multi material manufactured by L-PBF (paper in preparation)

This is the first time to our knowledge that a superalloy + bond coat system is produced by L-PBF at once. Investigation of the oxidation behavior, previously reported, highlighted the reliability of the bi-material. This bi-material is intended for use in high temperature exposure conditions such as aeroengine gas turbine applications. For this reason, resistance to thermal shock (rapid heating and cooling cycles) is also critical. A Burner Rig approaching the severe conditions which could be experienced during operation imposing very rapid heating and cooling was used. This study was conducted in collaboration with our colleagues in Italy from the Department of Applied Science and Technology of Politecnico di Torino. The Burner Rig in Politecnico di Torino (Fig. 18) is a custom made apparatus specifically designed to reproduce the severe temperature conditions experienced in a gas turbine engine. The machine is equipped in particular with a burning chamber consisting of a heat resistant ceramic plate machined to host up to 10 specimens, 20 air quenching tubes for fast cooling of the top and/or bottom surface of each specimens, 20 thermocouples for precise measurements of the top and bottom surface of each specimens. On top of the plate, a burner delivers an actual flame by burning either methane or a mixture of regulated methane+fuel. Only methane was used in the present study. A system of manual valves with manometer allows the regulation of the air flow rate for quenching. The Burner Rig is controlled manually by means of a software which allows to regulate the methane flow and the air/methane ratio, trigger air quenching and monitor/control the temperatures. The maximum designed temperature is 1300°C. In the present study, the methane flow was set to 1 m 3 /h, the air/methane ratio was set to 1.2, the maximum temperature was set to 950°C and the quenching temperature was set to 300°C. All experimental tests were carried out in the frame of a two month testing campaign in Turin. Also part of the international collaboration created around the present Ph.D. thesis, characterization of the materials was the subject of a 4-month The details of these different conditions are described in the following:

-L-PBF superalloy substrate: the substrate, Alloy 625, is produced in all the cases by L-PBF with an SLM Solution SLM125HL machine according to the optimum conditions given by the manufacturer and verified using the recommended process conditions, as in the rest of the study [4.3.1].

-L-PBF Bond Coat: The NiCrAlY bond coat L-PBF processing optimization campaign has been described previously [4.3.1] and adapted same way.

-GBS heat treatment: This heat treatment was adapted to assess if this heat treatment is effective for coated sample either. The GBS heat treatment applied was also previously described in [4.3.5].

-Sprayed Top Coat: At last, after all heat treatments had been carried out, a YSZ top coat was deposited using sol-gel process. This deposition was realized in IMT-Mines Albi. 8mol% Y 2 O 3 -ZrO 2 (Tosoh) was dispersed into a zirconium propoxidebased sol and was subsequently deposited on substrates by air spraying. Consolidation was then carried out by a heat treatment at moderate temperature (600 °C) for 4h in order to avoid any significant alteration of the previous microstructure-setting treatments of the bond coated substrates. only under 1 min and the reheating to 950°C took approximately 4 min. After each session, the machine had to be cooled down and opened. Each sample was photographed and weighed with a precision balance (0.1 mg), before being repositioned in the chamber for the next session on the next day. Slight maintenance of the machine was also necessary after each session to avoid in particular degradation of the quenching tubes. Only methane was used as burning means and the air/methane ratio was set 1.2 to provide an oxidizing environment. Nevertheless, it seems clear in Fig. 22 that applying the GBS treatment before deposition of the top coat (GBS+TC) was found detrimental to the coating's adherence and weight loss due to coating's spallation was significant after only few cycles as opposed to AB+TC. It is assumed that some oxidation may have taken place during the GBS heat treatment realized in a tubular furnace under protective argon flow, however no clear evidence could be detected. SEM-EDS analysis revealed the Hardness was measured every 20 µm from the top surface of the bond coat to 200 µm deep into the substrate for all conditions using a test load of 50g (Fig. 25). The hardness values increased significantly after thermal cycles for the un-heat treated materials (AB and AB+TC), as opposed to GBS+TC material, especially within the bond coat part. The higher hardness was attributed to precipitation during thermal cycling between 300°C and 950°C (γ′′ in particular), though evidence are not yet available. On the contrary, heat treated sample GBS+TC showed only a little difference at the bond coat and substrate interface. It could also be noted that the hardness values of the substrate of the GBS+TC material was significantly lower than those of the as-built substrate (AB and AB+TC) in Fig. 25, which could be expected due to recrystallization and grain growth resulting for the heat treatment. This is consistent with the results of the mechanical properties in tension characterization [4.3.5]. This recrystallization and grain growth was also confirmed by EBSD analysis in Fig. 26. While the characterization campaign of the tested specimens is still in progress, all materials exhibited good adhesion of the bond coat and the substrate was excellent with no signs of delamination due to thermal shock testing.

Thesis for the Degree of Doctor Philosophy in CWNT & IMT Mines-Albi

86

VI. Conclusions

This research investigated the feasibility of producing a NiCrAlY bond coat on an Alloy 625 substrate by L-PBF. The properties of both an Alloy 625 substrate and an NiCrAlY-coated Alloy 625 substrate were evaluated. In addition, the optimum parameter conditions for MCrAlY coating of an Alloy 625 substrate by L-PBF were verified for the first time based on the welding metallurgy theory. In addition, GBS was successfully introduced in an L-PBF manufactured Alloy 625 substrate to strengthen its grain boundary. Therefore, this study demonstrated that Alloy 625 fabricated by L-PBF has high potential for real industrial applications by improving its manufacturing efficiency together with its high-temperature properties by L-PBF coating, which provides resistance to oxidation/hot corrosion, and GBS heat treatment, which provides resistance to GB cracking. Key contents of each paper are summarized as follows.

1. The optimum parameter conditions for MCrAlY bond coating by L-PBF on an Alloy 625 substrate were selected as follows: layer thickness t = 50 μm, hatch spacing h = 120 μm, laser power P = 250 W, and laser scanning speed v = 800 mm/s. These settings were based on several criteria, such as an approximately 250-μm-thick NiCrAlY coating with a very low residual porosity of ≤0.1%, including a ∼35% dilution zone, a smooth hardness profile between 275 and 305 Hv, low residual stress levels, and a progressive Al concentration distribution.

2. For the mechanical properties' evolution of L-PBF processed Alloy 625, various heat treatments were adopted, and each sample was evaluated through microstructure analysis. A fine columnar/dendritic structure with a high density of dislocations composed the as-built sample, and it contributed to its high strength. Direct aging showed the highest strength effect through the precipitation of γ″ phases and the discontinuous formation of elongated Cr-

  Thesis for the Degree of Doctor Philosophy in CWNT & IMT Mines-Albi 2 été suggérés pour améliorer les propriétés de l'alliage. Il a été vérifié que l'échantillon tel que construit en L-PBF avait une résistance suffisamment élevée; cependant, il avait le potentiel d'être fracturé plus rapidement que souhaité en raison de la faible ténacité et des défauts dus à un refroidissement non homogène. Par conséquent, divers traitements post-thermiques ont été appliqués pour améliorer ses propriétés mécaniques à température ambiante et à haute température. Avec tous ces traitements thermiques, il était possible de produire des échantillons fabriqués par L-PBF qui avaient des microstructures similaires à celles d'un ouvragé, qui est la microstructure préférée, composée de grains polycristallins et de précipités. Ils ont montré une ductilité améliorée et ont maintenu une résistance raisonnable. Bien que la microstructure traitée thermiquement ressemblait beaucoup à la microstructure forgée, ses joints de grains (GB) étaient généralement plus faibles à des températures plus élevées. Par conséquent, un traitement thermique par dentelure aux limites des grains (GBS) a été suggéré. Le GBS est une réaction spontanée dans laquelle des GB ondulés sont formés; la réaction peut être déclenchée par des traitements thermiques spécialement conçus. Il a été démontré qu'il est efficace pour inhiber la propagation des fissures, en particulier lors du fluage des superalliages à base de Ni principalement solides renforcés en solution. Un traitement thermique spécifique, comprenant un refroidissement lent entre la solution et les températures de vieillissement, a conduit avec succès au SGB dans les études précédentes. De plus, le mécanisme du GBS a été étudié pour le comprendre en profondeur afin qu'il puisse être appliqué à d'autres systèmes d'alliages.

 2 .

 2 Enfin, il était prêt à appliquer le GBS au L-PBF Alloy 625.  Le GBS a été induit avec succès dans l'alliage L-PBF 625 et ses propriétés de traction à haute température ont été évaluées avec du L-PBF recristallisé, du L-PBF tel que construit et de l'alliage 625 forgé. ductilité que le L-PBF Comme Thesis for the Degree of Doctor Philosophy in CWNT & IMT Mines-Albi 6 construit à haute température. En particulier, l'échantillon induit par GBS a une résistance à l'occurrence de DSA.  Une étude plus approfondie de la résistance à l'oxydation et au choc thermique de l'alliage revêtu de L-PBF NiCrAlY est nécessaire.  Un test TGA pour l'alliage 625 en vrac, l'alliage 625 revêtu de NiCrAlY en vrac et NiCrAlY, tous fabriqués par L-PBF, a été réalisé et l'échantillon revêtu a montré une résistance à l'oxydation intermédiaire entre deux échantillons en vrac. Par conséquent, la fabrication de composants revêtus par L-PBF n'est pas anodine. Le test de choc thermique a été testé pour un échantillon revêtu d'As-NiCrAlY avec et sans couche de finition. De plus, un échantillon revêtu d'As-NiCrAlY traité thermiquement avec une couche de finition a également été appliqué. Il n'y avait pas de délaminage entre le substrat et la couche de revêtement de liaison NiCrAlY, ce qui signifie que sa propriété de dilatation / contraction thermique est stable même dans les conditions de refroidissement / chauffage rapide. 1. A novel approach to the production of NiCrAlY bond coat onto IN625 superalloy by selective laser melting Cette étude portait sur la faisabilité de produire un matériau de couche de liaison MCrAlY directement sur un substrat en alliage 625 produit par L-PBF. Dans la présente étude, le L-PBF a été utilisé pour fabriquer à la fois le substrat de superalliage à base de Ni et la couche de liaison MCrAlY, ce qui est la première fois que cela a été étudié, au meilleur de nos connaissances, en relation avec la fabrication AM. Quinze paramètres de traitement différents ont été testés, et tous les cas ont montré une dilution substantielle entre NiCrAlY et le substrat de superalliage en raison de la 7 caractéristique de refusion typique du processus L-PBF, assurant une excellente liaison métallurgique. Les paramètres optimaux du procédé, épaisseur de couche t = 50 μm, espacement des hachures h = 120 μm, puissance laser P = 250 W et vitesse de balayage laser v = 800 mm / s, ont été sélectionnés en fonction de plusieurs critères. Un revêtement NiCrAlY d'environ 250 μm d'épaisseur a été obtenu avec une très faible porosité résiduelle ≤0,1%, comprenant une zone de dilution ∼35%, un profil de dureté lisse entre 275 et 305 Hv, de faibles niveaux de contraintes résiduelles et une distribution progressive de la concentration en Al . Ces résultats ont confirmé la faisabilité du L-PBF comme méthode alternative pour le dépôt de revêtements dans les systèmes TBC ou le revêtement MCrAlY lui-même. Influence of heat treatments on microstructure evolution and mechanical properties of Inconel 625 processed by laser powder bed fusion Cette étude portait sur le comportement mécanique à température ambiante et la microstructure d'échantillons d'alliage 625 tels que construits et traités thermiquement fabriqués par L-PBF. Contrairement au reste du travail, dans ce cas, les échantillons ont été réalisés avec des machines différentes (machine EOSINT M270 Dual Mode, qui est équipée d'un laser à fibre Yb 200W avec un diamètre de spot de 100 μm) avec de la poudre (distribution de taille entre d30 de 16 μm et d90 de 48 μm) à partir d'une source différente (puissance laser de 195 W, vitesse de balayage de 1200 mm / s, distance d'éclosion de 0,09 mm et épaisseur de couche de 0,02 mm) dans le cadre d'une étude menée en collaboration avec Politecnico di Torino en Italie . Bien que l'échantillon de L-PBF fabriqué tel quel a montré de bonnes propriétés de traction, un traitement post-thermique est recommandé pour obtenir des propriétés mécaniques spécifiques afin de satisfaire les exigences industrielles. Différents types de traitements thermiques ont été effectués pour produire chaque microstructure et ont sélectionné chaque traitement thermique représentatif par comparaison de la dureté Thesis for the Degree of Doctor Philosophy in CWNT & IMT Mines-Albi 8 Brinell. Les carbures γ˝ et M23C6 à grains recristallisés sont apparus dans l'échantillon traité par vieillissement direct (DA) (700 ° C / 24 h), produit à partir de la microstructure telle que construite, qui présentait des structures colonnaires à haute densité de dislocation et des carbures NbC. Un traitement en solution (S) (1150 ° C / 2 h) a donné à l'échantillon une structure de grains austénitiques équiaxes recristallisés avec des carbures MC primaires et secondaires. En outre, l'échantillon traité par solution / vieillissement (SA) (1150 ° C / 2 h -700 ° C 24 h) a montré la même microstructure que celle de l'échantillon S, mais avec des carbures γ˝ et M23C6. La température optimale pour les traitements thermiques a été choisie en comparant les valeurs de dureté Brinell et l'observation de la microstructure. Quatre types d'échantillons ont été soumis à des essais de traction à température ambiante avec l'échantillon tel que construit comme contrepartie. L'échantillon tel que construit a montré des valeurs de limite d'élasticité (YS) et de résistance à la traction ultime (UTS) de 783 MPa et 1041 MPa respectivement, avec une ductilité de 33%. L'échantillon DA présentait des valeurs YS et UTS de 1 012 MPa et 1 222 MPa, avec la ductilité la plus faible de 23%. L'échantillon S a montré les valeurs YS et UTS les plus basses de 396 MPa et 883 MPa avec l'incrément de ductilité le plus élevé de 55%. L'échantillon SA a montré une augmentation significative de la résistance: YS = 722 MPa, UTS = 1 116 MPa et une ductilité corrélée de 35%. Selon les résultats de traction, le traitement en solution contribue à l'augmentation de ductilité et le traitement de vieillissement améliore la résistance. Par conséquent, il a été démontré que l'évolution de la microstructure de l'alliage L-PBF 625 sous divers traitements thermiques offre un large spectre de propriétés mécaniques possibles.

  heat treated wrought Alloy 625 Ni-based superalloy Récemment, l'introduction du GBS a été étudiée principalement pour les superalliages à base de Ni renforcés en solution solide à faible fraction volumique γ ′, tels que l'alliage 230 et l'alliage 617. Le présent travail a étudié l'induction du GBS dans superalliage à base de Ni renforcé en solution ayant une minorité de caractérisation de renforcement par précipitation par le carbure MC et γ˝. Le traitement thermique spécialement conçu (1250 ° C / 30 min -5 ° C / min de refroidissement lent -800 ° C / 0-1000 h) a été appliqué à l'alliage corroyé 625 sur la base des travaux précédents, et une dentelure significative a été observée pour la première fois. Des carbures M23C6 ainsi qu'une quantité significative de carbures MC ont été observés aux GB, alors qu'il n'y avait aucun signe de précipitation γ˝. Le traitement de refroidissement lent s'est avéré favoriser efficacement la formation de joints de grains dentelés, même en l'absence de vieillissement. Pendant ce temps, à mesure que le temps de vieillissement augmente, la précipitation de la phase δ était évidente.

à

  la fois les spécifications structurelles et environnementales. Du point de vue des spécifications environnementales, l'oxydation à haute température de NiCrAlY entièrement traité au L-PBF par les conditions de paramètres optimales précédemment testées et réussies (P = 250 W, v = 800 mm / s) recouvrant le substrat en alliage 625 les pièces ont été étudiées à des températures comprises entre 800 ° C et 1000 ° C. NiCrAlY, alliage 625 et bi-matériaux autoportants en vrac ont été fabriqués en utilisant la méthode L-PBF. Pour être représentatif des composants revêtus, deux couches minces (environ 100 um) de poudre NiCrAlY ont été balayées au laser à la surface de l'alliage 625 pour former le multi-matériau. Les échantillons revêtus ont ensuite été extraits en utilisant une technique de polissage mécanique douce pour assurer une épaisseur et une finition de surface répétables de l'échantillon. Il convient de noter que la fabrication du L-PBF NiCrAlY était particulièrement difficile et a entraîné des macro-fissures dans les échantillons de NiCrAlY en vrac en raison de la forte contrainte résiduelle volumétrique, alors que la raison pour laquelle des micro-fissures sont observées au niveau de la perle laser pour le revêtement est prédite par fissuration retardée. Cependant, il n'a pas été suffisamment endommagé pour empêcher la réalisation de l'étude. Comme prévu, le NiCrAlY en vrac et les bi-matériaux ont démontré un comportement d'oxydation amélioré par rapport à l'alliage 625. En effet, l'oxydation de NiCrAlY autoportant volumineux a donné une couche d'Al2O3 dense et continue dans cette plage de température, ce qui était attendu pour cette composition de NiCrAlY. Cela signifie qu'aucun effet d'évaporation de l'Al n'a été observé (un problème parfois observé dans le L-PBF). Cependant, la formation d'oxydes à la surface du matériau bicouche différait de celle du NiCrAlY autoportant volumineux. A titre de comparaison, l'oxydation du matériau bicouche a conduit à des régions à couche d'Al2O3 dense et continue ainsi qu'à des régions composées d'un mélange de Cr2O3 externe et d'Al2O3 interne. Des cartes EDS étendues à la surface du bimatériau ont mis en évidence la distribution hétérogène des éléments constitutifs du revêtement NiCrAlY, entraînant certaines régions avec une activité Al inférieure à celle requise pour la formation d'une couche continue et dense d'Al2O3. Les régions à 13 faible activité en Al et à haute activité en Al étaient liées à la topographie de la surface du L-PBF et correspondaient aux régions de pic et de vallée, respectivement d'où il provient de l'épaisseur homogène de la couche de poudre déposée. De plus, les fissures ont été principalement trouvées dans les régions à haute activité Al et elles ont pu être atténuées en appliquant un traitement thermique de détente ou la diffusion. La fabrication de composants revêtus mais petits avec un revêtement cassant par L-PBF n'est pas anodine et nécessite une étude plus approfondie. 7. Thermal shock resistance of the NiCrAlY coated Alloy 625 multi material manufactured by L-PBF (paper in preparation) C'est la première fois à notre connaissance qu'un système de superalliage + couche de liaison est produit à la fois par L-PBF. L'étude du comportement d'oxydation, précédemment rapportée, a mis en évidence la fiabilité du bi-matériau. Ce bi-matériau est destiné à être utilisé dans des conditions d'exposition à haute température telles que les applications de turbines à gaz de moteurs d'avion. Pour cette raison, la résistance aux chocs thermiques (cycles rapides de chauffage et de refroidissement) est également critique. Un gréement de brûleur approchant les conditions sévères qui pourraient être rencontrées pendant le fonctionnement imposant un chauffage et un refroidissement très rapides a été utilisé. Cette étude a été menée en collaboration avec nos collègues italiens du Département des sciences appliquées et de la technologie du Politecnico di Torino. Le Burner Rig in Politecnico di Torino (Fig. 18) est un appareil sur mesure spécialement conçu pour reproduire les conditions de température sévères rencontrées dans un moteur à turbine à gaz. La machine est équipée notamment d'une chambre de combustion constituée d'une plaque en céramique résistante à la chaleur usinée pour accueillir jusqu'à 10 échantillons, 20 tubes de trempe à l'air pour un refroidissement rapide de la surface supérieure et / ou inférieure de chaque échantillon, 14 chaque spécimen. Sur le dessus de la plaque, un brûleur délivre une flamme réelle en brûlant soit du méthane, soit un mélange de méthane régulé + carburant. Seul le méthane a été utilisé dans la présente étude. Un système de vannes manuelles avec manomètre permet la régulation du débit d'air pour la trempe. Le Burner Rig est contrôlé manuellement au moyen d'un logiciel qui permet de réguler le débit de méthane et le rapport air / méthane, déclencher la trempe de l'air et surveiller / contrôler les températures. La température maximale prévue est de 1300 ° C. Dans la présente étude, le débit de méthane a été réglé à 1 m3 / h, le rapport air / méthane a été fixé à 1,2, la température maximale a été fixée à 950 ° C et la température de trempe a été fixée à 300 ° C. Tous les essais expérimentaux ont été réalisés dans le cadre d'une campagne d'essais de deux mois à Turin. Fait également partie de la collaboration internationale créée autour du présent doctorat. thèse, la caractérisation des matériaux a fait l'objet d'un programme de stage de 4 mois. Un étudiant d'IMT Mines-Albi, sous la direction du présent auteur Jiwon Lee, était responsable de la caractérisation.

Figure 1 .

 1 Figure 1. Total additive manufacturing market size (billion $) by opportunity segment and its prospect by 2027 [17].

Figure 2 .

 2 Figure 2. Pie-chart of generally used AM materials (searched in google scholar, total 100 papers).

Figure 3 .

 3 Figure 3. Pie-chart of study fields in AM related to Ni-based superalloy

Figure 4 .

 4 Figure 4. Various AM methods and their available materials [24].

Figure 5 .

 5 Figure 5. Schematic diagram of DED process [26].

Figure 6 .

 6 Figure 6. Schematic comparison of pre-scanning heated EBM and L-PBF with preheated base plate [39].

  Superalloys are classified as Fe-based, Ni-based, and Co-based depending on their base metal, and Ni-based superalloys are the most often used.Fe-based alloys are widely applied in discs and rotors with relatively low operating temperatures for economic reasons, although they have lower mechanical properties at 27 temperatures higher than 750 °C. Co-based superalloys are used at extremely high temperatures above 1000 °C, but they have the disadvantage of relative high cost of cobalt. Finally, Ni-based superalloys use nickel as a base metal in combination with other elements, such as Cr, Co, Al, Ta, W, Re, Mo, and C to improve their hightemperature mechanical properties[START_REF] Sims | A history of superalloy metallurgy for superalloy metallurgiests[END_REF].

Fig. 7 [Figure 7 .

 77 Fig. 7 [44].

Table 4 .

 4 Roles of alloying elements in the Ni-base superalloys [42, 44]. Phase relation / property Roles Solid solutions strengthening of γ W. Mo, Ti, Al, Cr, Re strengthen Fe, Co, Cu Strengthen slightly Solid solution strengthening of γ´ Mo, W, Ti, Cr strengthen V, Co weaken Mn, Fe, Cu strengthen slightly Mole fraction of γ´ Cr, Ti, Al, Nb, Mo, Co, Ta, V, Fe increase Anti-phase boundary energy Ti, Co, Mo, Fe increase Al, Cr decrease γ-γ´ lattice mismatch Ta, Nb, C, Ti increase Cr, Mo, W, Si, V, Cu, Mn decrease Al, Fe, Co, negligible effect Coarsening rate of γ´ Ti, Mo, Nb, Co, Fe decrease Cr increase B, Zr no effect Oxidation and hot corrosion resistance Cr, Hf, Si, Y, La, Pt improve other element variable Precipitation strengthening of a Ni-based superalloy means strengthening by precipitation of γ´ and γ˝ phases. Here, γ´ is an A 3 B-type intermetallic compound having an ordered FCC structure (L1 2 ), which is coherent with the matrix. In an A 3 Btype compound, A is mainly composed of electronegative elements, such as Ni and Co, and B is composed of electropositive elements, such as Al, Ti, Nb, Hf, and Ta. Elements like Cr, Mo, W, and Fe can be substituted for either A or B. Accordingly, the general chemical formula of γ´ is (Ni, Co) 3 (Al, Ti) [45]. The γ´ phase is an L1 2ordered structure in which the Ni atom is located at the center of the crystal, and the Al atoms are located at each vertex.There are two main mechanisms of strengthening in a precipitation-strengthenedNi-based superalloy. First, when the γ´ phase is precipitated, dislocations movement is hindered, and they stack upon each other due to a coherence difference from the matrix. Secondly, if a dislocation goes into the γ´ phase, it shows an anti-phase boundary (APB) and is separated into partial dislocations. The formation of an APB or a stacking fault (SF) between two partial dislocations hinders dislocation's motion, which means that the structure is 'strengthened'.

Figure 8 .

 8 Figure 8. The variation of strength increment with the particle size. The strengthening mechanism is changed by critical size (r 0 ) of a precipitate [46, 47].

[Ni 3

 3 Nb] with an orthorhombic DO a structure[START_REF] Sundararaman | Some aspects of the precipitation of metastable intermetallic phases in INCONEL 718[END_REF][START_REF] Shankar | Microstructure and mechanical properties of Inconel 625 superalloy[END_REF][START_REF] Decker | The Superalloys[END_REF]. Large amounts of solute elements are added to improve the high-temperature mechanical properties of Ni-based superalloys, which are combined with the C. The carbides present in Ni-based superalloys are mainly observed at the grain boundaries; these include MC, M 23 C 6 , and M 6 C. MC carbide is generated at high temperature combining Ti, Ta, Nb, W, Co, Cr, and Hf with C, which is produced in the form of a blocky and script shape during solidification. The M 23 C 6 carbides are produced in the temperature range of 790 to 816 °C, especially when Cr binds to C, which is deposited Novel Fabrication of Alloy 625 and MCrAlY Bond Coat by L-PBF and Microstructure Control 33 in the grain boundaries and improves creep resistance. M 6 C carbides are produced in a temperature range of 816 to 982 °C in alloys containing large amounts of Mo and W, and they have excellent high-temperature stability. As shown above, these carbides have a stable temperature range and react with different phases to form other types of precipitates. They undergo the following reactions [51]: Here M´ is gamma phase depleted in certain alloying elements and M˝ gamma phase of slightly different compositionThe other phases which can be found in Ni superalloys are topological close packed (TCP) phases, and they include laves phases, μ, η, and σ. They are precipitated as plate or acicular shapes in Ni-based superalloys. They are brittle, so when they are subjected to stress, they are susceptible to cracking around them and their morphology make them preferential crack initiation sites. Therefore, TCP phases should be suppressed as much as possible. In particular, the σ phase, which is a metallic compound formed in the range of 650 to 925 °C, is the most problematic one. It is characterized by a structure similar to M 23 C 6 carbide. Hence, it is mainly nucleated in M 23 C 6 carbide, and

Figure 9 .

 9 Figure 9. Block diagram of Ni-based superalloy development [54].

Figure 10 .

 10 Figure 10. Time-temperature transformation (TTT) diagram of Alloy 625 [56].

Figure 11 .

 11 Figure 11. (left) High density of dislocations forming columnar structure in L-PBF Alloy 625

  TBCs are composed of a bond coating part (MCrAlY) and a ceramic top coat part (ZrO 2 -7 wt% Y 2 O 3 ). The bond coat part is primarily deposited on the substrate alloy surface between the substrate and top coat. It balances the thermal expansion mismatch between two different layers because failures often occur due to delamination of the coating part due to the thermal expansion difference. In addition, it provides oxidation resistance as a reservoir of oxide film formation elements and provides roughness, which improves top coat adherence. Gas turbine engines show improved efficiency when this coating method is adopted. It is used in combustion cans, transition pieces, nozzle guide vanes, and blade platforms[START_REF] Tawancy | Failure mechanism of a thermal barrier coating system on a nickel-base superalloy[END_REF][START_REF] Bennett | A philosophy for thermal barrier coating design and its corroboration by 10000h service experience on RB211 nozzle guide vanes[END_REF].

Figure 12 .

 12 Figure 12. Schematic of general coatings used for superalloy applying at high temperature; diffusion coating, overlay coating, thermal barrier coating (from left to right) [67].

  Thesis for the Degree of Doctor Philosophy in CWNT & IMT Mines-Albi42MCrAlY coating is often applied to structural materials used at high temperatures, such as turbine blades and vanes. Oxidation/corrosion damage of Ni-based superalloys can be generally prevented by overlay MCrAlY coating. The element M depending on the substrate matrix, is usually composed of Ni, Co, or NiCo. Cr is added to provide resistance to hot-corrosion, and Al is required to form the surface oxide layer on the surface of the superalloy substrate. Most of the MCrAlY is composed of FCC-γ matrix due to its high solubility of Cr and Al, 25% and 11%, respectively, and preferable mechanical properties like high temperature characteristics and good toughness. The Y elements improves resistance to oxygen diffusion taking place along grain boundaries and a thermally grown oxide (TGO) layer by substituting into the alumina or segregating at the grain boundaries (GBs). Moreover, Y-oxide pegs can act as anchors on the coating surface to enhance the TGO adhesion. It plays different roles depending on the oxide type; it promotes the formation of chromia scale, decreases the scale growth rate, improves scale adherence in chromia-forming alloys, whereas its major role is improving the adherence of the scale in alumina-forming alloys[START_REF] Nijdam | Effect of reactive element oxide inclusions on the growth kinetics of protective oxide scales[END_REF][START_REF] Guo | Effect of co-doping of two reactive elements on alumina scale growth of β-NiAl at 1200 °C[END_REF][START_REF] Nijdam | Effect of Y Distribution on the Oxidation Kinetics of NiCoCrAlY Bond Coat Alloys[END_REF].

Figure 13 .

 13 Figure 13. Predicted phase map of NiCrAlY alloy (typically 22 wt% Cr and 10 wt% Al, bal. Ni, the very low amount of Y was not taken into account) [77].

Figure 14 .

 14 Figure 14. One example of heat-exposed (from (a) to (e)) MCrAlY coated GTD-111DS sample at 1100 °C. Volume fraction of β-NiAl phase in the bond coat part is decreased and depleted zone is increase with longer exposure time [78].

Figure 15 .

 15 Figure 15. Comparison of polycrystalline, directional, and single-crystal alloys [86].

Figure 16 .

 16 Figure 16. Comparison of powder size distribution.



  Further study about resistance to oxidation and thermal shock of L-PBF NiCrAlY coated alloy is required. TGA test for the bulk Alloy 625, bulk NiCrAlY and NiCrAlY coated Alloy 625 all manufactured by L-PBF was conducted and coated sample showed intermediate oxidation resistance between two bulk samples. Therefore, the fabrication of coated components by L-PBF is not trivial. Thermal shock test was tested for As-NiCrAlY coated sample with and without top coat. Additionally GBS heat treated As-NiCrAlY coated sample with top coat was also applied. There was no delamination between substrate and NiCrAlY bond coat layer, which means its thermal expansion/contraction property is stable even in the rapid cooling/heating condition.The optimized parameters of L-PBF fabrication of a NiCrAlY bond coat onto an Alloy 625 substrate are described based on several criteria. Significant remelting of the underlying substrate led to the formation of an intermediate dilution zone, which indicates excellent metallurgical bonding. This result suggests the feasibility of L-PBF as an alternative method for the deposition of coatings in TBC systems or MCrAlY coating itself. Prior to evaluation of the properties of each coated sample, the reliability of the mechanical properties of the Alloy 625 substrate was firstly evaluated.Because of this reason, tensile testing at ambient temperature of the L-PBF fabricated Alloy 625 was firstly conducted. Three different heat treatments were performed to obtain the specific mechanical properties of the L-PBF Alloy 625. Through direct aging treatment, solid-solution treatment, solid-solution followed by aging treatments, the samples showed three different microstructures and tensile properties, which were compared to those of the as-built sample. The as-built sample showed high tensile strength due to the high dislocation density and fine MC carbides with very fine dendritic structures. Solution treatment alone at 1150 °C for 2 h led to recrystallization due to high density of dislocations through the repetition of rapid cooling and heating, which increased the ductility. Aging treatment at 700 °C for 24 hours contributed to γ˝ and M 23 C 6 formation. Each heat treatment and their combination affected each sample's microstructure, which was the basic cause of their mechanical properties. Furthermore, from this study, it was verified that simple solutionizing heat treatment can induce recrystallization in Alloy 625 fabricated by L-PBF. This room temperature tensile test results led to the interest in mechanical property improvement at high temperature.The high temperature tensile test was conducted with the L-PBF manufactured Alloy 625 as-built and microstructure-controlled samples. The recrystallized/solisolution sample showed a high possibility to improve the mechanical property, especially with regard to ductility by microstructure control. We now focused on the ductility improvement of L-PBF Alloy 625. Therefore, GBS was considered as a one way to improve ductility of the sample by enhancing GBs. GBS is intended to strengthen the GBs, which are likely to be weakened at high temperature by improving the GB energy. For the successful application of GBS to L-PBF manufactured Alloy 625, an in-depth analysis and understanding of GBS is essential. The principle of GBstrengthening by GBS can be demonstrated by phase transformation from straight GB to serrated GB due to the decrement of interfacial energy through serrated GB transition. Its mechanism was explained by investigation of a Ni-20Cr model alloy in which enrichment of Cr near and at GBs exerted strain on the GBs for the onset of GB serration. The actual minimum critical strain energy needed for GBS onset was

  temperature was also modified (800 °C→870 °C) considering the MCrAlY coating stabilization/diffusion heat treatment for the future study to adapt this heat treatment to the Alloy 625/MCrAlY bi-material. The application of modified GBS heat treatment (1300 °C/90 min-direct aging/slow cooling-870 °C/2 h) to the L-PBF alloy was the first successful attempt ever. This GBS heat-treated sample and recrystallized/solidsolution treated (1200 °C /1 h) sample were subjected to high-temperature tensile testing with the L-PBF as-built and wrought Alloy 625 as the counterparts. The two heat-treated samples seem to show superior tensile properties at high temperature inNovel Fabrication of Alloy 625 and MCrAlY Bond Coat by L-PBF and Microstructure Control59comparison to the as-built sample. Even the GB serrated sample showed a higher elongation value than the recrystallized/solid-solution sample at a much higher temperature. However, they could not match the outstanding elongation value of the wrought alloy sample. The reason for the ductility loss of the L-PBF samples was investigated through analysis of their typical microstructures. In addition, the positive effect on the dynamic strain aging (DSA) of each heat treatment was also assessed. Now, in order to make sure that the reliability of NiCrAlY coating on Alloy 625 substrate by L-PBF, the oxidation/hot corrosion behavior and stability in rapid cooling/heating condition was investigated by TGA testing and thermal shock testing.The TGA test was performed for the L-PBF manufactured bulk Alloy 625, L-PBF manufactured bulk MCrAlY and L-PBF coated sample. The coated sample showed an intermediate result of oxidation kinetics compared to the bulk Alloy 625 and bulk NiCrAlY samples from the TGA tested at 800 °C, 900 °C , and 1000 °C . It shows that improvement in high temperature oxidation properties of L-PBF processed Alloy 625 is possible using a L-PBF processed NiCrAlY coating. Moreover, the adhesion of the bond coat and the substrate was excellent with no signs of delamination in thermal shock conditions. The L-PBF coating properties were validated by TGA testing and thermal shock testing, and consequently, it was demonstrated that the feasibility of the application of L-PBF for TBC coating on the Ni-based superalloy substrate was certified to be commercialized.

4 .

 4 This section is based on an article which has been published in J. W. Lee, M. Terner, H. -U. Hong, S. H. Na, J. B. Seol, J. H. Jang, T. H. Lee, Mater. Charact., "A new observation of strain-induced grain boundary serration and its underlying mechanism in a Ni-20Cr binary model alloy", 135 (2018) 146-153, DOI: https://doi.org/10.1016/j.matchar.2017.11.047. Thesis for the Degree of Doctor Philosophy in CWNT & IMT Mines-Albi 68 First evidence of grain boundary serration in a specifically heat treated wrought Alloy 625 Ni-based superalloy Recently, the introduction of GBS has been studied mainly for solid-solutionstrengthened Ni-based superalloys with low γ′ volume fraction, such as Alloy 230 and Alloy 617. The present work studied inducing GBS in Alloy 625, which is a solidsolution-strengthened Ni-based superalloy having a minority of precipitationstrengthening characterization by MC carbide and γ˝. The specially designed heat treatment (1250 °C /30 min -5°C/min of slow cooling -800°C/0-1000hrs) was applied to wrought Alloy 625 based on the previous work

70 5 .

 5 Heat treatment design for superior high-temperature tensile properties of Alloy 625 produced by SLM L-PBF manufactured Alloy 625 was subjected to high-temperature tensile testing.

Figure 19 .

 19 Figure 19. Geometry of each specimen and composition of materials in each layer.

Fig. 21

 21 Fig.21shows the samples as-prepared (t = 0 cycle) and after the entire thermal shock testing campaign (t = 450 cycles). Degradation of the sample (the ceramic Top Coat in particular) due to thermal shock was clear, and it seems that the samples subjected to the GBS heat treatment experienced the most severe degradation. There was no obvious sign of oxidation despite the oxidizing environment and the high temperature.

Figure 21 .

 21 Figure 21. The comparison of the samples' surface change before and after thermal shock tests.

Figure 22 .

 22 Figure 22. Mass change during thermal shock testing.

Figure 24 .

 24 Figure 24. OM micrographs showing cracking in the as-prepared AB specimens (top) and after 450 cycles (bottom).

84

Figure 25 .

 25 Figure 25. Vickers hardness profiles from the surface into the substrate for all three materials conditions before (0 cycle, blue) and after (450 cycles, red) thermal cycling between 300°C and 950°C.

Figure 26 .

 26 Figure 26. EBSD grain orientation map showing recrystallization and grain growth of GBS+TC (right) compared to AB (left) and AB+TC (center) as a result of heat treatment.

Table 2 . Comparison of manufacturing conditions of EBM and L-PBF

 2

		EBM	L-PBF
	Beam source	Electron beam	Laser
	Atmosphere	Vacuum	Ar atmosphere (<oxygen 0.1 %)
	Preheating	>600 °C	200-500 °C
	Powder size	45-105 µm	15-45 µm
	Average powder layer thickness	20-100 µm	50-200 µm
	Beam scan speed	0.3-1.0 m/s	>1000 m/s

the Degree of Doctor Philosophy in CWNT & IMT Mines-Albi 28Table 3 . Compositions of typical wrought Ni-based superalloys

 3 Table 3 show typical wrought Ni-base superalloy compositions. Ni-based superalloys have become more prominent in recent years due to the limitations of Fe-and Co-based superalloys. In particular, they are used in the main engines of aircraft and gas turbines for power generation, where hightemperature corrosion and oxidation resistance is required. Ni-based superalloys are

		C	Cr	Ni	Mn	Si	Mo	W	Co	Al	Ti	Nb	Cu	Ta	Fe
	Hastelloy X	0.1 5	22. 0	Bal	0.6	0.5	9.0	0.6	1.5						18. 5
	Alloy 600	0.0 4	15. 4	Ba.	0.2	0.1									9.7
	Alloy 601	0.0 3	21. 9	Ba.	0.2	0.2	0.1			1.4	0.3	0.1			14. 5	Solid solution
	Alloy 617	0.0 8	21. 6	Ba.	0.1	0.1	9.5	0.9	12. 5	1.2	0.3				hardening type
	Alloy 625	0.0 5	21. 5	Ba.	0.3	0.3	9.0	0.9	0.5	0.4	0.4	2.0		1.7
	Haynes 230	0.1 1	21. 7	Bal	0.5	0.4	1.4	1.4		0.3				1.2
	IN718	0.0 8	20	54	0.1	0.1	2.9			0.5	1.0	5			Bal	Precipita-
	Alloy 263	0.0 6	20	Bal	0.6	0.4	6		20	0.6	2.4		0.2		0.7	tion hardening
	IN740	0.0 3	25	Ba.	1.2	0.5	0.5		20	0.9	1.8	2			0.7	type (10~50%
	Waspaloy	0.0 3	19	Bal	0.2	0.2	4		13	1.5	3.0		0.2		0.7	γ´)

divided into 'wrought alloys', 'powder metallurgy (PM) alloys' and 'cast alloys' according to the method of manufacture. They are usually manufactured by PM or the casting method if cold working or hot working (e.g. forging, rolling, etc.) is difficult, which is the case if the ratio of gamma prime reinforcement (see section 2.2.2) is between 30% and 40%. Meanwhile, the gamma prime phase increases the creep strength as the fraction increases.

The wrought alloys can be divided into 'solid solution-strengthened alloys', which show high-temperature strength by solutionizing added elements, and 'precipitationstrengthened alloys', which show greatly increased strength due to precipitation by aging treatment.

Thesis for

Table 5 . Summary of composition and crystal structure for the various phase in Ni-based superalloys [42-44].

 5

									Lattice
	Phase	Chemical formula	Ti	Nb	Mo	Cr	Fe	Lattice	parameters
									(Å )

Table 6 . Element composition of Alloy 625

 6

	Ni	Cr	Fe	Mo	Nb	C	Mn	Si	P	S	Al	Ti
	Bal.	20-23	5.0*	8.0-10.0	3.15-4.15	0.1*	0.50*	0.50*	0.015*	0.015*	0.40*	0.40*
										* means maximum value
	3.2.4 Precipitation formation of Alloy 625
	Three different types of carbides, namely, MC, M 6 C, and M 23 C 6 as well as
	intermetallic phases, such as γ˝ and δ can precipitate in Alloy 625 as indicated in
	Figure 10. Each precipitate is formed during heat treatment depending on the
	precipitation temperature and duration. Three kinds of carbides are usually
	precipitated at the grain boundaries. The MC carbide (M is Nb and Ti) is precipitated
	as thin grain boundary films at the temperature range from 871 to 1037 °C, but it is
	usually present during solidification as a large primary carbide. The primary MC
	carbide consumes a large amount of Nb, which is one of the main elements for
	strengthening; hence, homogenization and solid-solution treatment are accordingly
	required. M 6 C carbide (M is Ni, Nb, and Mo) and M 23 C 6 carbide (M is almost entirely

Cr) commonly have a blocky shape and stable up to 1080 °C and 700 to 800 °C

Novel Fabrication of Alloy 625 and MCrAlY Bond Coat by L-PBF and Microstructure Control 37

Table 7 . Comparison of room-temperature tensile properties of Alloy 625 manufactured by the PBF method and performed heat treatments.

 7

	Manufacturing / condition	Tensile property (YS(MPa)/UTS(MPa)/El.(%))	reference
	ASTM (standard) F3056-14	275/485/30
	L-PBF/as-built	800 ± 20/1030 ± 50/8-10	[60]
	L-PBF/as-built (modified 625 composite)	641.5 ± 23.5/878.5 ± 1.5/30 ± 2	[61]
	EBM/ annealing treatment (1120C/100MPa/4h)	360/-/58	[62]
	L-PBF/ Stress-relieved + HIP (1120C/100MPa/4h)	507/ 827/69..3	[63]

Table 8 . Phase information of NiCrAlY alloy.

 8

	Phase	Crystal structure	Lattice parameter (Å )
	γ (NiCr)	fcc	3.52
	γ´	ordered fcc, L1 2	3.572
	β-NiAl	ordered fcc, B2	2.88
	α-Cr	BCC	2.88

Table 9 . Chemical composition in wt.% of the Alloy 625 and NiCrAlY powders used.

 9

		Ni	Cr	Mo	Fe	Co	C	Nb	Al	Y
	Alloy 625	Bal.	20.80	8.27	3.49	0.6	0.08	3.12	0.35	-

This section is based on an article which has been published in G. Marchese, M. Lorusso, S, Parizia, E. Bassini, J. W. Lee, F. Calignano, D. Manfredi, M. Terner, H. -U. Hong, D. Ugues,

This section is based on an article which has been published in M. Terner, J. W. Lee, J. H. Kim, H. -U. Hong, Int. J. Mater. Res., "First evidence of grain boundary serration in a specifically heat treated wrought Alloy 625 Ni-based superalloy" 109 (2018) 1-8, DOI: https://doi.org/10.3139/146.111674.

brittle coating by L-PBF is not trivial and need further investigation. This section is based on an article which has been published in D. Texier, E. Copin, A. Flores, J.W. Lee, M. Terner, H. -U. Hong, P. Lours, Surf. Coat., "High temperature oxidation of NiCrAlY coated Alloy 625 manufactured by selective laser melting" Tech., 398 (2020) 126041, DOI: https://doi.org/10.1016/j.surfcoat.2020.126041.

52

L-PBF fabrication of a NiCrAlY bond coat onto an Alloy 625 substrate, itself produced by L-PBF has been studied as one part of this thesis [4.3.1]. Therefore, laser power P=250 W, laser scanning speed v=800 mm/s were selected based on several criteria. The process environment conditions, such as platform preheating (T = 200 °C) and O 2 content (<0.10%) remained unchanged in all research series. A constant layer thickness of 50 μm was set by lowering the build platform. A low carbon steel plate that could handle heat conduction of the Alloy 625 substrate was placed on the building platform (x-y table) and levelled.

Microstructure analysis

Sample preparation: All samples were mounted and polished with SiC abrasive paper down to 1200 grit followed by sequential polishing with 1-μm diamond paste.

They were then etched using Aqua Regia (as-built, heat treated, oxidation tested and thermal shock tested samples) and 100 ml of HCl with 0.5 g of CrO 3 mixed etchant

Novel Fabrication of Alloy 625 and MCrAlY Bond Coat by L-PBF and Microstructure Control

53

(tensile tested samples) for optical microscopy (OM) and scanning electron microscopy (SEM) analysis.

Transmission electron microscopy (TEM) samples were prepared by polishing to a thickness of ~200 µm. Then 3-mm diameter discs were punched from these thinned samples, which were electropolished to electron transparency in a TenuPol-5 jet polisher operated at a polishing current of 50 mA at a voltage of ~25 V, using a solution consisting of 200 ml of perchloric acid and 800 ml of methanol at -26 °C to -28 °C. Two samples were tested for each condition and placed on diametrically opposite sides of the ceramic plate in the Burner Rig chamber.

Observation

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The test conditions for the materials were ideally the most representative of the severe conditions envisaged in-use. As many as 450 cycles were carried out under very aggressive conditions for the tested materials. Further cycling was avoided to prevent complete degradation of the ceramic top coat for further analysis. With regards to the nature of the superalloy substrate, a maximum temperature of 950°C (2 min plateau) and quenching to 300°C was considered. Note that the 950°C is the average temperature in the chamber. The actual temperature for each individual specimen varied significantly and was conveniently recorded by 20 thermocouples for precise measurements of the top and bottom surface of each specimen. Nevertheless, the gradient of temperature was lower than 100 °C. Fig. 20 shows the first 30 cycles, totally 15 times of repetition were carried out for each sample. All specimens were hot mounted in phenolic resin and carefully subjected to conventional metallographic preparation of grinding with abrasive SiC papers and fine polishing with diamond pastes down to 1 µm. Fig. 24 shows OM micrographs of the top surface of as-received AB specimens (top) and after 450 cycles (bottom). Significant cracking was observed within the bond coat in both states. It was expected that it is occurred due to delayed cracking, which is already discussed in [4.3.6]. rich M23C6 carbides in the remaining columnar/dendritic structure. A solidsolution treatment caused recrystallization. This sample showed increased strength due to the added post-aging treatment. To achieve the required mechanical properties, appropriate post-heat treatment is required for L-PBF manufactured Alloy 625 to improve its mechanical properties.

3. Before application of the GBS to the L-PBF fabricated Alloy 625, the formation mechanism of GBS was investigated. Only in the case of Cr atom enrichment near and at the GBs, serrated GBs were observed from the experiment with a model alloy. This means that Cr atom enrichment caused the minimum strain of 0.97 kJ/mol on the GBs, which led to the onset of GBS.

4. Based on the GBS mechanism, GBS was introduced to Alloy 625 for the first time. A specialized heat treatment (1300°C/90min-5°C /slow cooling-870°C /2hrs) was designed to induce GBS for Alloy 625, and it worked well. Unlike the results of a previous study, the serrated GBs were decorated not only with M23C6 but also MC carbides.

5. For the high-temperature tensile test application, GBS heat treatment was newly introduced to the L-PBF fabricated Alloy 625 ('GBS HT' sample). The as-built sample has significantly higher strength; however, its ductility was too low. Therefore, it was modified by GBS HT, which showed the increment of ductility from 600 °C, while its strength remained higher than that of wrought Alloy 625. Recrystallized L-PBF Alloy 625 also showed higher strength than that of wrought Alloy 625, while its ductility was lower than that of the GBS HT sample at temperatures higher than 600 °C. The recrystallization and fine precipitates contributed to relatively high strength and high ductility. The GBS HT sample showed results similar to those of the RX HT sample in terms of both strength and ductility, but it showed a small difference in the DSA phenomenon. The temperature range at which the undesirable DSA 7. GBS heat treatment was applied to the MCrAlY-coated Alloy 625 fabricated by L-PBF, and its thermal shock properties were compared to the non-heattreated sample. After 450 cycles of thermal shock testing in the temperature range from 300 to 950 °C, the weight gain was increased in the non-heattreated sample, while the heat-treated sample showed weight loss. The result of weight gain in the non-heat-treated sample can be attributed to the surface oxide layer forming as the sample endures each cycle. The weight loss is due to delamination of the top coating due to the surface oxide layer after heat treatment while the delayed cracking phenomenon in bond coat part was decreased after heat treatment. In all cases, samples exhibited no signs of oxidation and adhesion of the bond coat and the substrate was excellent with no signs of delamination due to thermal shock testing.

8. From present study, the feasibility of L-PBF manufacturing of TBC coating on the Ni-based substrate was certified to be commercialized. In addition to that, the effect of GBS heat treatment for Alloy 625 to high temperature properties showed positive effect and it will be continued to applied and studied for the

Novel Fabrication of Alloy 625 and MCrAlY Bond Coat by L-PBF and Microstructure Control

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other high temperature properties.

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Nouvelle fabrication de l'alliage 625 et du revêtement de liaison MCrAlY par la fusion laser sélective sur lit de poudre et contrôle des microstructures

Dans cette étude, l'alliage 625 a été élaboré par l'une des méthodes de fabrication additive (AM) les plus couramment utilisées, la fusion laser sélective sur lit de poudre (L-PBF), et ses propriétés mécaniques ont été évaluées à différentes températures. L'alliage 625 fabriqué par L-PBF a montré une résistance élevée et un allongement à la rupture médiocre. Ainsi, des traitements thermiques ont été appliqués pour améliorer ses performances. Un traitement thermique de mise en solution à une température supérieure à 1000 ° C a été appliquée à l'alliage L-PBF 625, ce qui a entraîné une recristallisation en raison de l'énergie stockée par le matériau attribuée à la forte densité de dislocations présentes dans l'alliage. Cette microstructure modifiée a montré une résistance satisfaisante lors des essais de traction à température ambiante (résistance plus élevée que l'alliage corroyé 625 et allongement supérieur à l'alliage L-PBF brut de fabrication). À l'étape suivante, l'alliage L-PBF 625 recristallisé a été soumis à un traitement thermique de Grain Boundary Serration (GBS, dentelure des joints de grains) pour améliorer ses propriétés mécaniques à haute température. Comme il s'agissait de la première tentative pour générer la GBS d'un alliage à haute teneur en Nb, il était d'abord important de comprendre le mécanisme GBS. Pour induire le GBS, il est nécessaire que les larges atomes de soluté se déplacent près des joints de grains (GB). Par conséquent, le traitement thermique GBS a été modifié pour être appliqué à l'alliage L-PBF 625. Le traitement thermique de GBS spécialement conçu a réussi à induire pour la première fois les motifs en zigzag des GB dentelés. Cet alliage GB-L-PBF dentelé 625 a montré des propriétés mécaniques améliorées à haute température en termes de ductilité accrue et d'élimination de l'effet de vieillissement dynamique (DSA). Pour améliorer davantage les propriétés à haute température de l'alliage L-PBF 625, un revêtement de liaison NiCrAlY a été nouvellement appliqué au substrat en alliage 625 par la même méthode (L-PBF) pour améliorer l'efficacité du processus de production et augmenter la résistance à l'oxydation/corrosion à haute température. Bien que leurs différentes propriétés thermiques aient conduit à de nombreux essais et erreurs dans la fabrication du matériau, les paramètres optimaux ont été définis et vérifiés pour évaluer le potentiel de commercialisation du procédé. La caractéristique de refusion du L-PBF a induit une bonne liaison entre le substrat et le revêtement, ce qui indique une bonne stabilité. Le comportement à l'oxydation de l'alliage 625 revêtu de NiCrAlY a été caractérisé par analyse thermo-gravimétrique (TGA) et des tests de chocs thermiques ; les résultats ont indiqué que le nouveau matériau avait une résistance à l'oxydation plus élevée que l'alliage 625 brut de fabrication. Par conséquent, le traitement thermique de GBS associé à un revêtement NiCrAlY efficace peut améliorer considérablement les propriétés mécaniques à haute température de l'alliage 625 fabriqué par L-PBF.

Novel Fabrication of Alloy 625 and MCrAlY Bond Coat by Laser Powder Bed Fusion and Microstructure Control

In this study, Alloy 625 was fabricated by one of the most commonly used additive manufacturing (AM) methods, laser powder bed fusion (L-PBF), and its mechanical properties were evaluated at various temperatures. The L-PBF fabricated Alloy 625 showed high strength and relatively poor elongation. Thus, some heat treatments were applied to improve its performance. A solid-solution heat treatment with a temperature of more than 1000 °C was applied to the L-PBF Alloy 625, resulting in recrystallization because of high energy stored within the alloy attributed by high density of dislocations. This modified microstructure of the L-PBF Alloy 625 sample showed the required strength under tensile testing at room temperature (higher strength than wrought Alloy 625 and greater elongation than L-PBF asbuilt alloy). In view of enhancing mechanical properties at high temperature, a grain boundary serration (GBS) heat treatment was specifically designed for L-PBF Alloy 625. Because this was the first attempt to produce GBS in a high-Nb-content alloy, it was necessary to understand its mechanism first. To induce GBS, it is necessary for large solute atoms to move near the grain boundaries (GBs). Therefore, the GBS heat treatment was modified for application to the L-PBF Alloy 625. The specially designed GBS heat treatment successfully induced the zigzag patterns of serrated GBs for the first time. This GBS L-PBF Alloy 625 showed improved hightemperature mechanical properties in terms of increased ductility and elimination of the dynamic strain aging (DSA) effect at elevated temperatures. To further improve the hightemperature property of the L-PBF Alloy 625, NiCrAlY bond coat was applied to the Alloy 625 substrate by the same method (L-PBF) for the first time to improve the efficiency of the production process and increase the resistance to oxidation. Although their different thermal properties led to many trials and errors in the manufacturing of the material, the optimal parameters for applying NiCrAlY bond coat deposition by L-PBF were set and verified to assess the potential for the process to be commercialized. The remelting characteristic of L-PBF induced good metallurgical bonding between the substrate and coating, which indicates good stability. The oxidation behavior of the NiCrAlY-coated Alloy 625 was characterized by thermal gravimetric analysis (TGA) and thermal shock testing; the results indicated that the novel coated material had higher resistance to oxidation than bulk Alloy 625. Therefore, the GBS heat treatment together with efficient NiCrAlY coating can greatly improve the hightemperature mechanical properties of L-PBF manufactured Alloy 625.

Keywords

Laser Powder Bed Fusion, Alloy 625, Heat treatment, NiCrAlY, Oxidation, Mechanical properties