Cyclic Oxidation of Diffusion Aluminide Coatings

Abstract

The diffusion aluminide coatings are used for high-temperature applications. Structural materials of particular components degrade during service due to fatigue, creep, oxidation, corrosion and erosion. The requirements of higher efficiency of modern industrial applications increase the development of new structural materials, technologies and protective coatings. Properties of many structural materials such ultimate tensile strength, creep strength and fatigue are generally optimized for maximum high-carrying loading with less emphasis on environmental resistance. For these applications, the performance characteristics are limited by the operating conditions, which can be tolerated by the used materials. The main structural materials for high mechanical and thermal loading are superalloys protected against aggressive environment by coatings. Cyclic oxidation is the superposition of thermal cycles in an oxidation environment. The main goal of the experimental work was to compare the cyclic oxidation of protective Al and AlSi coatings deposited on both Inconel 713 LC and MAR-M247 superalloys. The resulting graph revealed that samples from IN 713 LC without coating show good resistance and their mass change is maintained above zero limit. Samples from MAR 247 LC with both Al and AlSi coatings appear to be the most acceptable selection of combination relating to superalloys/coating.

1. Introduction

Increasing the technical level, utility values, quality and reliability of engineering components belong to the main directions of economic development. The surface treatment also plays a significant role in the final quality of the structural unit. An inappropriate choice of surface treatment, non-compliance with the technological procedure or just an unaesthetic appearance can devalue an otherwise excellent technical work.

Surface engineering includes those processes that make it possible to modify the surfaces of structural components with the aim of improving primarily their mechanical and utility properties and, last but not least, their esthetic appearance, or to reduce the economic costs of production. It is therefore a modification of the surface of the component, which does not affect the bulk properties. The main goal is to increase the service life of components and the operational reliability of equipment. It is often aimed at increasing corrosion resistance, surface sliding wear resistance, providing a diffusion barrier, thermal insulation, etc. while reducing production costs, especially energy savings with an impact on minimizing environmental pollution.

There are several ways to treat metal surfaces, which can be divided into three main categories:

1. Strengthening by phase transformation is the treatment of the surface without changing the chemical composition of the substrate, while the structure of the surface is changed by thermal or mechanical treatment. The most common request is to increase the hardness while maintaining the toughness of the bulk material.

2. Modification of surface layers by changing its chemical composition through diffusion processes at elevated temperatures without affecting the properties of the substrate.

3. Surface treatment by coating deposition. This group includes a wide group of coating processes, where the properties and chemical composition of the substrate differ significantly from the substrate. Unlike previous surface treatment methods, the coating/substrate interface is clearly distinguishable in this method. The most important aspect is the thickness of the coating, its porosity and adhesion.

1.1 Modifying the surface without altering chemical constitution.

1. Heat treatment

   Heat treatment is the practical application of physical metallurgy with the aim of modifying the material structure to obtain desired properties. Strengthening by phase transformation essentially belongs to thermal technology processing. Surface hardening is a common method of changing the mechanical properties of surface layers especially for steels. The purpose of this treatment is to increase hardness and wear resistance component surface. The conditions of transformation or breakdown of austenite are decisive depending on temperature, time and rate of cooling. To austenitize the surface layer, it is necessary to perform heating at a higher rate than the heat is dissipated into the part. Phase diagrams are used to understand the heat treatment process. For example, an Isothermal diagram (IT) describes the formation of austenite, Time Temperature Transformation (TTT) diagram.describes the decomposition of austenite, Continuous Cooling and Continuous Heating diagrams are commonly referenced to as CCT and CHT diagrams, which are also used to understand the effects of heating and cooling of steel during heat-treatment process [1]. In practice, heating is commonly used by flame or electrical induction. Recently, they have been focusing on themselves for changes induced by laser irradiation or electron beam techniques [2, 3, 4, 5, 6]. Surface heat treatments by laser are manufacturing technologies that are gaining industrial interests inthe last years.

   Laser radiation makes it possible to concentrate the power density up to 10¹⁶ W.m⁻² and deliver considerable amount of energy to the treated surface contactlessly and quickly. To the characteristic energy parameters of radiation, pulse length or speed of movement have a great influence on the temperature field temperature source (for a continuous laser), temperature-physical characteristics and geometric dimensions processed body. Depending on the power density and duration of action, either only heating and cooling in the solid phase takes place in the surface layer, or there is melting of the surface and the possibility of forming an amorphous state during particularly rapid cooling. Rapid temperature and structural changes cause thermal stresses, which can also result in the formation of cracks. Lasers are particularly suitable for curing relatively small or inaccessible surfaces. Laser is considered a useful technique since it offers control over case depth, uniform microstructure, uniform hardness, and minimum material distortion. Laser transformation hardening of steels is a diffusion-less transformation process and eliminate the formation of pearlite phases [7].

   In addition, the interestingness of this process lies in the possibility of direct integration of a very flexible laser heat source without the use of a quenching medium, as well as in the possibility of producing various microstructures with a hard surface and residual compressive stresses. Electron beam techniques have lower investment and operating costs.

2. Mechanical treatment

   Cold working the surface by shot peening or other specialized surface treatments were designed to enhance the surface integrity and service life of structural components. Most of these mechanisms are very sensitive to microstructure. Surface treatments like Ultra Sonic Peening, Laser Shock Peening, and Shot Peening are used to enhance the surface microstructural properties that help in improving the service life. These applications lead to deformed surface layers and increase the stored energy, form compressive residual stresses due to higher density of defects in the crystal structure, increase the hardness, fatigue life and stress corrosion resistance. Residual compressive stresses and increased surface hardness are the key outcomes of shot peening [8, 9, 10]. The simulation methodology and numerical prediction are used to optimize microstructure and shot peening parameters to reduce the relaxation of compressive residual stresses during service life [11].

3. Thermomechanical treatment

   This process consists of a combination of heat treatment and plastic deformation. Thermomechanical processing is based on a combination of several operations such as deformation, heating and cooling performed in different cycles, which further determine their classification. A high density of structural imperfections such as dislocations, vacancies or layer defects and their distribution due to deformation mechanisms affect the processes of structural changes and thus the resulting properties and integrity of the surface of the components. However, structural changes generate new ones and redistribute the original ones. Thus, the mechanism and kinetics of these structural changes depend on the nature and density of imperfections and consequently affect the number and distribution of such imperfections. Thermomechanical treatment is although used to improve the mechanical properties of alloys by introducing twins and nano-precipitation phase in the lamellar structure of the homogenized alloys [12]. Multi-stage thermomechanical processing with the stepwise temperature reduction can improve not only the mechanical but also the electrical properties of the alloy [13].

1.2 Altering the chemistry of surface regions of the substrate

1. Thermochemical diffusion treatments are widely applied to transport reactive interstitials elements (carbon, nitrogen, boron or their combinations) from a medium (gas, plasma etc.) to the surface of structures. Reaction of these elements with solid under higher temperatures lead to the phase transformation and forming of separate layers at the surface [14]. Thermomechanical processes such carburizing, nitriding, boriding or carbonitriding are different from processes chromizing, aluminizing and siliconizing in which the elements Cr, Al or Si are represented as a substitute in a solid solution. Diffusion of interstitial elements into steels can take place at low (ferritic range) or high temperatures (austenitic range). This processes are also used for decreasing friction coefficients and improving wear resistance of structures [15].

2. Ferritic processes include convenient gas nitriding to enhance the surface hardness of structures by uniform nitride layer [16]. Plasma nitriding diffusion technique improve tribological properties of steel surfaces depending on processing parameters and demanded layer characteristics [17]. Low thermal ferritic nitrocarburizing include various treatment media such salth bath, endothermic ammonia gas mixtures, and methane or propane/ammonia/oxygen mixtures. Nitrocarburized layers consist of an inner diffusion layer and outermost thin compound layer, which is detrimental due to its brittle nature. Transition layer can be formed between the above two layers in some cases [18].

3. Austenitic treatments broadly include carburizing process, which requires a high temperature and carbon donor. Various agents fulfill the function carbon donor for the different carburizing processes (powder, salt bath or gas atmospheric and low pressure carburizing). Carbonitriding and boronizing belong to the austenitic treatment and they are performed at elevated temperatures (near 900°C). Laser–beam carburizing process is used to enhance fatigue wear resistance of surface steel structures. In this process, the surface of the material is coated with graphite prior to laser irradiation. Carburizing can be achieved by two mechanisms: (a) the surface alloying mechanism, which incorporates melting of the substrate and dissolution of the graphite in the liquid phase; and (b) the solid-state diffusion mechanism, which incorporates austenitization of the substrate and carbon diffusion in austenite [19].

4. Electroplating is the process using electrodeposition of metals onto surface from an anode (part containing desired metal) to a cathode (part to be plated). The anode and the cathode are placed into electrolyte chemical bath and exposed to the electrical charge. This process has high productivity, simplicity of equipment and low operating costs. More metals (Cd, Cr, Cu, Au, Fe, Ni, Ag, Ti, Zn) and their combination can be used for electroplating. Plastic parts manufactured by 3D printing can be also electroplated [20, 21]. This combination offers unique tensile strength options for finished designes.

5. Oxide coatings have been developed for large variety of properties such as high chemical stability, crystal structures, compositions, physical, mechanical, optical or electrical properties with various thickness, porosity etc. In general, they are brittle at room temperature. They are largely used as insulating coatings. Thin layer of oxide coatings deposited on the surface serve as physical barrier between substrate and environment. Metals, nonmetals and ceramics can be used as protective coatings and various method has been applied to form functional coatings.

   Nowadays, the structural materials made for marine, aerospace engineering, transport vehicles must be protected against corrosion, abrasion, or high temperature oxidation and dense and adhesive oxide coatings improve the resistant of substrates against these degradation modes and enhance the surface chemical and mechanical properties [22]. Nanocoatings developed in thin layer having scale 1–100 nanometers thickness offer much better processing properties than conventional coatings. They provide effective solution against chemical or mechanical effects. Nanocoatings are not protective layers but their particles bind themselves both physically and chemically to the substrate surface and give it very good protection. Nanocoatings have huge application in different industries except for automobile, marine or aircraft industries, also in defense, electronic or medical industries [23, 24].

6. Sulfur treatments are used due to their good lubricating properties. Laser cladding and low temperature ion sulfurizing technology (300°C) are processes for deposition of sulfur layers with wear resistance and self-lubrication performance. The resulting product FeS layer have excellent anti-friction properties [25].

7. Phosphating is one of the most treatment technology in electrical or automotive industries because low production cost and simplicity. Processes include zinc, calcium, manganese or magnesium phosphate coatings especially with excellent lubricity. Chemicals involved in oxidizing accelerators such as nitrates, nitrites, chlorates, hydrogen peroxide, etc. and heavy metal ion accelerators, which are usually toxic to humans and lead to the environmental problems. Therefore, the main goal is to investigate ECO-friendly accelerators for phosphating process [26, 27].

8. Ion Implantation is the most common application in the semiconductor industry. Various ions can be implanted onto steel structures with the aim to control crack propagation and corrosion. The technology by which ions of gaseous or metallic elements are accelerated in an electrical field and impacted into a solid. Ions impinging on the component cause many chemical and physical changes of materials transferring energy and momentum to electrons causing a structural change. The ion energy and the target composition determine the depth of ion penetration. Some studies reported that some of ions have antimicrobial effect and its nanoparticles act as antimicrobials [28].

1.3 Surface treatment by coating deposition

The current age is the age of surface treatment technologies that allow countless possibilities of obtaining coatings with multiple properties. State-of-the-art manufacturing techniques are used in diverse areas, such as optical memories and filters, fiber lasers, LED displays, various implants or environmental devices.

1. Weld cladding is technique to achieve enhanced mechanical and tribological properties of several engineering components such as high strength and corrosion resistance. It is technique that allows to create relatively thick coating on substrate material (from 1 mm to several cm) using TIG welding, plasma transferred arc cladding or laser beam cladding [29].

2. Laser surface alloying have been used for improve mechanical properties and promote functional applications and of surfaces of advanced structural applications [30]. This technique makes it possible to obtain novel multiphase coatings to improve mechanical, wear properties and refine the microstructure [31].

3. Thermal spraying is an important surface modification method and plays crucial role in industrial production for depositing of metals, alloys, metal oxides, various ceramic and composite materials or polymers on a variety of substrates to form functional surfaces. Key functions include restoration, corrosion and erosion protection, heat insulation or conduction, electrical insulation or conduction etc. Thermal spraying can be formed either by deposition of ductile solid metals (referred as cold spaying) or by deposition through molten or semi-molten particles (flame spraying, high-velocity oxygen or air fuel, plasma spraying, wire arc spraying). Among advantages of these processes belong high deposition rate, low processing cost, wide thickness range and minimal degradation of substrates. However, this technology is a line-of-sight process and coatings have lower bond strength, are porous and have anisotropic properties. Porosity is sometimes beneficial in case the coating must have low thermal conductivity [32]. Liquid thermal spraying provides a novel method for preparation of micro-nano structured coating, which significantly reduces manufacturing cost. In this method suspended droplets in the plasma are first atomized and vaporized, then partially sintered and melted before hitting the substrate surface to form a coating [33].

4. Galvanizing is the process of forming coating with zinc which provides corrosion protection of steels. Galvanized coatings are widely used for steel reinforcement in concrete to increase the long-therm service life of concrete structures. It has been shown that galvanized bar slows down cracking during service and concrete spalling is less likely or delayed [34].

5. Chemical Vapor Deposition (CVD) is based on the dissociation of metal compound vapors at elevated temperatures to produce thin, diffusion-bonded, adherent coatings of metal carbides, nitrides, carbo nitrides and oxides; typically TiN, TiC, Ti(CN) and Al₂O₃. CVD coatings are used on carbide tool tips (indexable inserts) and on selected tribological items. Plasma assisted chemical vapor deposition (PACVD) uses electrical discharge to produce plasma and this method permit forming of thin solid film on the substrate. Recent studies have reported higher effectiveness of process and better properties of uniform coating using hollow cathode [35].

6. Physical Vapor Deposition (PVD) is a process based on physical principles. The essence of the method is thermal evaporation of the material from the target, or ion sputtering on the surface of the substrate, where a thin, highly adhesive layer is subsequently formed. PVD processes take place at relatively low temperatures of (150-500) °C in the presence of an inert gas or a vacuum, in which no toxic substances are released. Therefore, this method is ecological and characterized by high purity of the process. In most cases, the grain sizes in PVD coatings are below 100 nm and are therefore classified as nanostructured materials. The properties of PVD coatings are mainly influenced by their chemical composition, structure (size and shape of grains, number and distribution of defects, density of the coating) and also residual stresses, which depend on the application method and parameters. PVD processes generate compressive stresses in the coatings, and the coatings exhibit high modulus of elasticity, hardness and low coefficient of friction. The PVD coating process includes various techniques such as arc evaporation, magnetron sputtering and ion plating, which make it possible possible to apply thin layers of precisely determined thicknesses to the base material [36].

7. Painting is the technology of applying a thin organic coating to the surface of a material for the purpose of corrosion protection of substrates or improving the esthetic appearance. Compared to other processes for protection or decoration, the technology offers a number of advantages such as lower costs, a wide range of coatings with requirements for color, gloss and surface structure, resistance of the coating to various corrosion agents and their combinations, good dielectric properties, if necessary also ensuring suitable conductivity for grounding induced or static electricity. The organic coatings have been developed to meet the latest environmental regulations.

The preceding brief review shows an extraordinary extend and variability of methods utilized in the surface engineering. Hereafter, we will focuse on some important kinds of coatings on metallic surfaces and the processes of their degradation in service. Metals as the primary construction materials used in industry are exposed to oxidation and to active chemical environment. To ensure the quality and protection of metal products, different kinds of metal coatings are employed in the industry. There are two basic determinants of metal coatings.

The first class of coatings serves as a protection against the aggressive environment, i.e., they increase the resistance to thermal oxidation, corrosion and erosion.

The second class represents layers obtaining by surface hardening which increases surface strength for load carrying, surface hardness of a substrate material and introduce compressive residual stresses.

Thus, with respect to their function, the coatings can be divided into two main classes:

1. surface protection and.

2. surface hardening.

Diffusion aluminide coatings (DACs), overlay coatings and thermal barrier coatings (TBCs) are typical representatives of the first class. They improve the resistance to high temperature creep, fatigue, thermal oxidation, corrosion and erosion. Nickel and cobalt based superalloys are mostly utilized as substrate materials for turbine blades, blade rings, stator/rotor discs, etc. From the point of view of corrosive and oxidizing effects, to the monolayer protective coatings belong DAC and overlay coatings. DAC are based on the intermetallic compound β-NiAl that forms under the influence of the substrate (usually Ni superalloys). The technological procedures lie in various methods of deposition of an equally balanced suspension of aluminum- and silicon powders on the substrate surface together with activators (e.g., NH₄Cl) and organic binders (e.g. coloxylin). Diffusive tempering is applied at about 1000°C in an inert argon atmosphere or vacuum with the dwell time of several hours and subsequent slow cooling in the retort. As shown in Figure 1, the resulting microstructure of the layer consists of two sublayers – the outer layer (OL) and the diffusion zone (DZ). The microstructure of OL is composed of the β-NiAl phase with small number of complex phases and carbides Al-Cr-Ni, Mo-Cr-Nb. The DZ is formed by β-NiAl with many Cr-Mo, Mo-Cr-Nb rich particles. DACs are an integral part of the substrate material with just a slightly higher hardness. They primarily role is to protect the surface from high-temperature oxidation, corrosion and erosion but they should not reduce the resistance to creep, cyclic creep and thermo-mechanical fatigue [37]. On the other hand, the composition of the overlay coatings (known as MCrAlX) remains independent of the substrate alloy.

Figure 1.

Corrosion and oxidation resistance of metallic coatings on superalloys.

The most popular representatives of the second class are carburized or nitrided surface layers (CNSLs) produced by gas diffusion or plasma deposition processes - the case hardening. This process increases the surface hardness of a substrate material and introduces compressive residual stresses which results in a higher resistance to fatigue and wear of low-carbon steels, the materials mostly used in various dynamically loaded components as gears and shafts at room temperature. All the above mentioned components are subjected to cyclic loading in bending (tension) or, most frequently, to a combined bending (tension)/torsion loading.

2. Diffusion aluminide coatings

The diffusion aluminide coatings belong to the wide category of coatings used for high-temperature applications. Structural materials of particular components degrade during service due to fatigue, creep, oxidation, corrosion and erosion. With requirements of higher efficiency of modern industrial applications increase the development of new structural materials, technologies and protective coatings. Properties of many structural materials such ultimate tensile strength, creep strength and fatigue are generally optimized for maximum high-carrying loading with less emphasis on environmental resistance. For these applications, the performance characteristics are limited by the operating conditions which can be tolerated by the used materials. The main structural materials for high mechanical and thermal loading are nickel and cobalt-base superalloys protected against aggressive environment by coatings [38, 39]. The coating systems currently in use particularly in gas turbine blade applications can be divided into three generic groups [40]:

1. Diffusion aluminide coatings (DACs) are based on the intermetallic compound β-NiAl that forms under the influence of the substrate (usually Ni superalloys). Application processes include pack cementation and gas-phase processes such as chemical vapor deposition.

2. The properties of MCrAlX substrate/coating system (where M represents elements of Ni, Co, Fe or their combination; X represents minor elements such as Y, Ce, Si, Ta, Hf, etc.) can be controlled and balanced for a specific application. In Figure 1 it is demonstrated how the applied coating can be selected on basis of the corrosion and oxidation performance.

3. Thermal Barrier Coatings (TBCs) is a complex system, which consist of three different layers: ceramic top layer (typically of YSZ), thermaly grown oxide (TGO formed at TBC/BC interface), metallic bond coat (BC) and substrate.

The coatings deposited on rotor turbine blades provide an optimal protect in the range of specified lifetime against destructive effects of high-temperature corrosion, oxidation and erosion provided the following requirements are satisfied [41]:

• High oxidation and corrosion resistance.

  The coating must be thermodynamically stable and creates protective thin oxide film of uniform thickness. Slow growth rate of protective and adherent scale is desirable. In accordance with these requirements the coating should contain higher content of element (Al, Cr) to able to form protective scales.

• Interface stability. Low diffusion rate across the interface coating/substrate at use temperatures, minimized brittle phase formation and no undesired changes within coating are required.

• Good adhesion. Good adherence of coating to substrate; optimized properties of coating surface before deposition (rough or smooth); matched coating/substrate properties to reduce thermal stresses; minimized growth of residual stresses during coating deposition.

• Mechanical strength. Coating must withstand service-related stresses (creep, fatigue and stresses generated by impact of others particles). Suitable combination of strength and toughness is demanded and similar match of both thermal expansion coefficients for minimization of thermal stresses and thermal fatigue.

Diffusion aluminide coatings have been designed to protect nickel-base superalloys at elevated temperatures and they are based on simple aluminide or modified aluminide compounds. DACs and modified DACs involve the surface layers of substrate alloy which is enriched with aluminum, chromium, silicon, platinum, zirconium, palladium, yttrium or hafnium to form protected oxides through diffusion and/or enhance TGO adherence [42]. Their thickness is in the range of (10–100) μm. These elements react with the main constituents of substrate and form intermetallic compounds (aluminides) with significant part of element comprising protected oxide. For nickel-base superalloys NiAl (Figure 2) is the main aluminide compound of coating which produces corrosion and oxidation resistant thermally grown oxide.

Figure 2.

Binary phase diagram Al-Ni.

The selection of the appropriate coating composition depends on requirements of the corrosion/oxidation resistance and the environment. The main alloying elements in metallic coatings are briefly described below [43, 44]:

Nickel is the base element in DACs in Ni base substrates because it depresses chemical activity of Al and minimizes the interdiffusion between coating and substrate.

Aluminum form on the top of coating oxide Al₂O₃ which protects against oxidation up to 1200°C.

Chromium reduces critical level of Al needed to form alumina and chromium oxides protect coating against hot corrosion and oxidation up to 900°C.

Silicon is effective against low temperature hot corrosion and promotes formation of Al₂O₃.

Platinum increases TGO adherence as well as reduces oxidation kinetics. The presence of Pt in a solid solution, or as two-phase coating, provide resistance to I. type of hot corrosion; if a continuous PtAl₂ phase is present, coating is resisting against II. type of hot corrosion.

Zirconium has beneficial effect on the adhesion of TGO, it hinders the outward diffusion of Al cations.

Palladium provides excellent high temperature oxidation resistance and scale adhesion by reducing or preventing void growth, mitigating the detrimental effects of sulfur and accelerating alumina scale healing after spallation. Pd replaces Pt in aluminide coatings due to high costs of Pt.

Yttrium is a typical oxygen-active element which has beneficial effect on corrosion behavior of coatings.

Hafnium additions to aluminides markedly improve the adhesion of thermally grown oxides thereby improving the high-temperature oxidation performance.

DACs are the most spread types of coatings for turbines blades. Their disadvantage is strong dependence on substrate composition. For the application in aerospace industry is very important to find the critical ductile-to-brittle transition temperature (DBTT) which limits the utilization. Material undergoes under DBTT to failure by brittle fracture. The DBTT of diffusion aluminide coatings on superalloys varies between (700 up to 900) °C [45]. Some of the factors influencing the DBTT are coating process, composition, microstructure and phase distribution. With decreasing content of Al, thickness and roughness of coating DBTT decreases. DBTT of NiAl phase is reduced by more than 100°C when Al content is lowered from 32 to 25 wt. %.

2.1 Technologies of deposition

Results of many examinations have confirmed that aluminide coatings are characterized with very good corrosion and oxidation resistance at high-temperatures. They have homogeneous microstructure with good thermomechanical properties. In order to protect critical internal surfaces from oxidation and hot corrosion, the state of the art involves the use off diffusion coatings, formed by pack cementation, chemical vapor deposition (CVD), above-the-pack diffusion processes and/or by slurry cementation.

2.1.1 Pack cementation

Halide-activated pack cementation is a CVD process in which “pack” is made up of a mixture of a pure metal or alloy source, a halide salt activator and an inert filler material, usually Al₂O₃.

A filler allows the active pack constituents (the master alloy and activator salt) to be well distributed. This allows the halide vapors to easily reach the substrate surface, resulting in an overall uniform coating. The substrate to be coated is either buried inside the pack (the ‘in-pack process’) or suspended outside the pack (the ‘out-of-pack’ process). A controllable atmosphere, usually Ar or H₂/Ar, surrounds the pack as it is heated at a definite process temperature ranging from 700 to 1050°C. The metallic powder reacts with the halide salt activator to form volatile metal halide species of significant partial pressures. In the case of Pt-modified aluminide coatings, components are electroplated with a thin Pt layer prior to the aluminization.

Suppose NH₄Cl is used as the activator, the reaction will be given below [46]:

$$\mathrm{M}\text{(s)}+\mathrm{x}{\mathrm{NH}}_4\mathrm{Cl}\text{(s},\text{g)}=\mathrm{MClx}\text{(g)}+\mathrm{x}{\mathrm{NH}}_4\text{(g)}$$E1

where M is the element to be coated.

For aluminizing process Al, AlCl, AlCl₂ and AlCl₃ vapor substances can be formed and for chromizing Cr, CrCl₂, CrCl₃ and CrCl₄.

Partial pressure gradients support vapor transport tio the metal surface where desired coating specified phase, microstructure and composition forms via dissociation or disproportionation of the halide molecules and reaction with substrate.

On this basis, the pack cementation can be described as follows [46]:

1. the reaction and a thermodynamic equilibrium between the activator and the powder, which determines the vapor pressures of the active gaseous substances in the pack;

2. the movement of these gases through the porous pack;

3. diffusion of the metal halides to the alloy matrix surface;

4. reactions on the surface to deposit the coating elements and form the products with vapor substances; and.

5. solid-state diffusion of the coating elements into the matrix to form the protective surface layer.

Illustration of these events are schematically visible in the Figure 3. All of these events dictate the nature of the final microstructure and thickness of the surface layer.

Figure 3.

Ilustration of steps occuring in the pack cementation process.

There are several pack cementation processes, i.e. pack aluminizing, chromizing, siliconizing, or co-deposition needed for suitable surface layers. Aluminized coatings and chromized coatings are two traditional diffusion coatings that have been extensively used for almost a century.

Halide-activated pack cementation is a simple and inexpensive process useful for different structure geometries and sizes. It has great potential for using by innovative ways.

2.1.2 Chemical vapor deposition

In the CVD method two basic processes, i.e. the low-activity and high-activity process of aluminizing are distinguished. During the aluminide coating deposition process at high temperature, about 1050°C, low aluminum content NiAl phases are created (low activity process), whereas at about 700°C, NiAl phases containing more aluminum are formed (high activity process).

In “low-temperature high activity” (LTHA) process the coating grows by the inward diffusion of Al and Ni₂Al₃ (δ-phase) forms at the matrix surface. Al further diffuses inward and the initial surface of the substrate remains at the outermost surface upon aluminizing. Alloying elements that diffuse more slowly and carbides in the substrate remain in δ-phase. Ni₂Al₃ is a brittle phase and furher vacuum annealing is necessary to transform it into desired thermodynamically stable NiAl (β-phase). A disadvantage of coatings obtained by LTHA process an overall thickness of structures, in some cases is not possible to use LTHA technology.

In “high-temperature low activity” (HTLA) process, thermodynamic activity of Al is lower and the coating grows by the outward diffusion of Ni, which leads to direct formation of additive β-NiAl phase above the initial surface. Interdiffusion zone forms in the substrate due to depletion of Ni and slowly diffusing elements.

Mixed coating growth mechanisms are also known and involve the diffusion of both Al and Ni [47]. This is schematically illustrated below (Figure 4).

Figure 4.

Schematically illustrated microstructures in the low and high activity processes after heat treatment.

An effective role of diffusion barrier between the coating and superalloy would slow the loss of Al to the substrate and limit or delay diffusion of refractory elements into the coating. Microstructure of superalloys IN 713 LC after aluminizing and heat treatment is visible in the Figure 5.

Figure 5.

Microstructure of coating on superalloy IN 713 LC after high activity CVD aluminizing and heat treatment.

The CVD method gives the possibility of control of the AlXn concentration and the temperatures of the processes and through this control of aluminizing process. The HTLA CVD aluminizing process indicate superior oxidation resistance compared with the LTHA process because clean β-NiAl phase is free of precipitates and carbides.

2.1.3 “Out of pack” or “over-pack” method

This method operates in a manner very similar to pack cementation; except the parts to be coated are suspended either above the pack or downstream from the pack (vapor generating) retort. The technology is based on placing the parts to be coated in a retort or vacuum furnace oven without coming into direct contact with the granulated mixture. The transfer of gases, which ensure the formation of the coating, is made possible by a neutral gas, guided by the container. Currently, several variations of the applied methods are already known (e.g. SNECMA). This process relates to the deposition of an aluminum coating on a metal part, especially on a hollow metal part comprising an internal liner [48].

The key step of the “out of pack” method is the formation of Al(Cl, F, Br)_n (n < 3) halides at temperatures above 800°C which react with the surface of superalloy by the following reaction [49]:

$$\mathit{Al}+{\mathit{AlX}}_3\to {\mathit{AlX}}_n$$E2

The AlX3 compound is called “activator”. At the temperature 900-1150°C halides in the gaseous phase are reacting with substrate according to the reaction:

$${\mathit{AlX}}_n+\mathit{Ni}\to {\mathit{\text{AlNi}}}_y+{\mathit{AlX}}_3$$E3

where: 1/3 < n < 3.

Activator is one of reaction products, is penetrating into the source of aluminum and is supporting the process according to the scheme:

$$\mathit{Al}+{\mathit{AlX}}_3\to {\mathit{AlX}}_n$$E4

$${\mathit{\text{AlNi}}}_y+{\mathit{AlX}}_3\leftarrow {\mathit{AlX}}_n+\mathit{Ni}$$E5

All reactions are in the equilibrium and they are leading for creating the intermetallic phase AlNiy.

CVD and “out-of-pack” methods are the same from the view of chemical reactions. The difference between both methods results from the fact, that reaction:$\mathit{Al}+{\mathit{AlX}}_3\to {\mathit{AlX}}_n$ is occurring in the special generator (external) where from products are transported in the gaseous phase to the retort in which coated elements are.

2.1.4 Slurry process

Slurry aluminide coatings belong to the high temperature diffusion processes. They have been well known for a number of years and are widely used to protect metallic surfaces from oxidation and hot corrosion. Composition and microstructure of coatings formed by slurry-aluminizing process are similar to those obtained by pack cementation. Coatings made by slurry-aluminizing have high quality and the technology is characterized by smallest consuming applied materials. This allows to obtain coatings of the lowest cost and slurry method is one of the cheapest available methods of manufacture of aluminide coatings. Because of the environmental constraints of pack cementation and of out-of-pack processes related to the use of halides and other hazardous chemicals, significant efforts led to the use of water-based technologies [50, 51]. The suspension designed for deposition process can be easily modified and may be stored for a long time. In the deposition process an aluminum-containing slurry is being applied onto the component to be coated, and during heat treatment the slurry react with the substrate forming an aluminide. Slurry is composed from a powder mixture of aluminum or an aluminum alloy plus an activator along with a binder either by spraying or brushing, which is then submitted to diffusion heat treatment at temperatures usually in the 1000–1200°C range. From the point of view of the coating microstructure properties, this method can form precipitates that inhibit the adhesion of the coating. Another problem lies on inaccessibility of some area as cooling internal passages of turbine blades. Modified aluminide coatings have been industrially developed to overcome the composition limitation that conventional slurries have. Al-Si eutectic slurries are commercially often used for elements working in hot-corrosion environments. No other components have been successfully diffused simultaneously by the slurry route. The microstructure of AlSi layer is shown in the Figure 6 and it is composed of two sublayers. The outer Al-sublayer is created by NiAl with content of Al < 50 at. % and lower amount of phases containing refractory elements. The inner Si-sublayer creates Ni-Al matrix and higher amounts of phases of elements as Si, Cr, Mo, Ti, W and Co. The thickness of the layer varies in the range of 20–50 μm. Al-sublayer is formed by inward diffusion of aluminum what is typical for high-activity coatings. The inner Si-layer can forms during aluminizing process or operation in the high temperature environment needle-like precipitates in a β-NiAl matrix, which are oriented perpendicular to the interface substrate/diffusion zone (Figure 7) where is although visible coating degradation.

Figure 6.

Microstructure of Al-Si coating on nickel-based superalloy JS6K.

Figure 7.

Microstructure of Al-Si coating on JS6K superalloy after operation in the high temperature environment.

2.2 Oxidation resistance of DACs

Protective coatings used on turbine blades were developed to serve as physical barriers between aggressive environment and the substrate. In addition, TBCs are used as thermal barriers, retard creep degradation and reduce the severity of thermal gradients. Up to now, however, no coating that would fully survive the aggressive turbine environment has been found.

The coating degrades during service at two fronts: at the coating/gas-path interface when service temperatures are below the melting point of the coating and at the coating/substrate interface at higher temperatures when diffusion mechanisms play the main role in degradation of the system substrate/coating.

The most serious degradation modes are as follows [38]:

• High-temperature oxidation.

• Hot corrosion.

• Damage by thermal and thermo-mechanical fatigue.

• Mechanical damage by erosion.

• Creep degradation during overheating.

Inter-diffusion of elements at the interface with the substrate that results in a creation of undesirable phases is, sometimes, also mentioned as an independent degradation mode.

DACs coatings have been designed to withstand three types of environmental attack: high temperature oxidation, high-temperature hot corrosion (type I) and low-temperature hot corrosion (type II). Oxidation is a special form of corrosion degradation mode which occurs when metals and alloys are exposed to the oxygen environment. However, it should be paid regard to that the oxidation of the DACs at the coating/gas-path interface results in the formation of a protective oxide scale and in respect thereof high-temperature oxidation is not explicitly a degradation mode. If the formed oxide scale is thin, slow-growing, and adherent, it protects the substrate from further oxidation and form barrier to further oxidation. If it be to the contrary oxide scales spall and substrate is exposed to the oxygen environment and suffers from consumption of metals. “Pilling–Bedworth ratio” (PBR) is an important parameter for prediction of the oxide protection properties. It was found out that if the volume of oxide is less than the volume of metal consumed in the reaction, then it is likely that a porous oxide layer will result. This criterion is effective for most metals and alloys of practical importance and by PBR we can assume if the oxide is protective or not. Based on this theory PBR for oxidation of alloys can be expressed as [52]:

$${\mathit{PBR}}_{\mathit{\text{alloy}}}=\frac{\mathit{\text{Volume of}}a\mathit{\text{mole of}}{B}_x{O}_y}{\mathit{\text{Volume of}}x\mathit{\text{moles of}}B\mathit{\text{in the alloy}}}.$$E6

It is believed that if $\mathit{PBR}>1$ compressive stresses are developed in the oxide scales while tensile stresses occur when $\mathit{PBR}<1.$The larger the difference of PBR from 1, the larger stresses. Although is well known that direct relation between PBR value and the level of stresses does not hold because mechanisms for growth stresses and their relaxation are complicated. PBR calculated for aluminum gave value 1.29. However, in the practice alloys are widely used for high temperature applications and values of PBR are different from those of metals. Authors in [52] calculated PBR for Ni₃Al, NiAl and NiAl₃ and their results indicated that the PBR for oxidation of Ni-Al alloys are larger in regard of PBR for Al metal. PBR for Ni₃Al was calculated from 1.71 to 1.88; for NiAl was in the range from 1.64 to 1.78 and PBR for NiAl3 from 1.48 to 1.57.

2.2.1 Cyclic oxidation

Cyclic oxidation is the superposition of thermal cycles in an oxidation environment. Alloys used at high temperatures are subjected to the operative cycles which vary widely depending of operating conditions.

More or less rapid temperature changes in oxidizing atmosphere result in thermo-mechanical stresses in the oxide scales which fail due to spallation. Turbine blades of aircraft engines are a typical example for working in such conditions. Turbine blades of aircraft engines are a typical example for working in these conditions.

The main parameters that determine the operating conditions of aircraft turbine blades are the gas temperature at the turbine inlet, pressure, velocity and composition of the gas flow. The gas temperature is the most important parameter determining the specific thrust and performance of an aircraft engine. The thermal composition of the gas flow at the turbine inlet is inhomogeneous; is caused by factors that determine the dynamics of the gas flow, such as the construction of the combustion chamber, the layout of the burners, or the combustion process. The stator disks and blades of the high-pressure turbine are the hottest part of the turbine, their temperature is (200-300) °C lower than the temperature of the gases, while the difference in temperature fields is based on the height of the blades and their circumference. The temperature field in the rotor part is more homogeneous and results from the high rotational speed of the impeller.

The critical elements of an aircraft turbine are the rotor blades. They determine the maximum permissible gas temperature and the lifetime of the entire engine. They are loaded with centrifugal and dynamic forces, which cause them to be strained by tension, bending and twisting. Tensile stresses are usually higher than (200-250) MPa and are different on the convex and concave sides, being the highest near the root of the blade. High temperatures and voltages, unstable load conditions and the possibility of resonant vibrations make vanes among the most complicated parts of the engine. The highest temperature load is in the upper third of the blade, where the centrifugal force is the lowest. The temperature of the gas stream can increase by up to 500°C in just a few seconds in aircraft turbines during transients or sudden regime changes. Static stresses and high temperature of the blades lead to their deformation as a result of creep. The combination of the effects of high temperature, dynamic forces, and thermal stresses causes blade failure due to thermal fatigue.

The surface of the blades is exposed to the effects of combustion products, which cause their degradation by oxidation, corrosion and damage by solid particles carried by the gas flow. The fuel combustion process takes place at the temperature of (2000–2200) °C. Hydrocarbon combustion is the process of its oxidation by oxygen from the air, the main products of combustion are CO₂ and water vapor. The gas flow at the entrance to the turbine contains, in addition to the main ones, also secondary products of combustion, such as certain amounts of CO, H₂, CH₄, solid particles of carbon and sulfur compounds and other elements entering the chamber from fuel or air. The difference in the sample of combustion products depends on the chemical composition of the fuel and on the environment (overseas, industrial, desert).

Cyclic oxidation tests are used to monitor scale adherence and the ability of scale to successfully regenerate after repetitive scale failure. The performance of materials is generally monitored by gravimetric method although others methods have been used [53]. Experimental techniques are needed to obtain both, data on the kinetics of the oxidation reactions and characterization of the scales formed during oxidation.

The most important method used to study oxidation kinetics and oxidation rate is the gravimetric method. The principle of this method is to examine weight change due to oxidation as a function of time. Heating and cooling can be quite rapid and short cycles can be used. A simple gravimetric technique involves the exposition of the sample with known area in a furnace followed by measuring the weight change at definite intervals of time, using a sensitive balance. In this method the experiment have to be interrupted every time the weight change is measured. The sample is frequently heated and cooled what causes changes in the scale behavior such as onset of scale failure, buckling, spallation and finally mass loss. Figure 8 represents typical mass changes during thermal cycling for three types of coatings. Non-protective behavior and negative weight change during thermal cycling represents gross scale spallation occurring from the onset of the cycling test. Protective behavior of a scale for a definite time is characteristic for materials that have a limited reserve of elements (Cr, Al or Si) to form stable oxides. After certain number of cycles scale composition changes to less protective spinels which spall rapidly and mass loss is recorded. Spallation is the loss of protective oxides at the coating/oxide interface. Strains induced by stresses during repeated thermal cycling due to thermal mismatch between oxide and metal, result in crack initiation and eventual spalling of the scales. In the case that the scale is too thin to sustain a temperature gradient, not thermal shock to the oxide. If the oxide thickens sufficiently, the strain energy stored in the oxide becomes greater than that for fracture of interface, the scale spalls [54].

Figure 8.

Schematic diagram of typical mass changes of coatings.

2.2.2 Experimental work to study of DACs oxidation behavior

The main goal of the experimental work was to compare the cyclic oxidation of protective aluminide coatings deposited on two types of nickel superalloys, Inconel 713 LC and MAR-M247 (Table 1). All samples with and without aluminide coatings were exposed to cyclic oxidation. Two type of superalloys were deposited by aluminide coating and Si modified aluminide coating using CVD out-of-pack process.

	Cr	Al	Mo	W	Nb	Ti	Ta	Zr	Hf	C	B	Co
IN 713 LC	11.85	5.8	4.54		2.27	0.72		0.11		0.04	0.015
MAR 247 LC	8.4	5.5	0.7	10.0		1.0	3.0	0.05	1.5	0.15	0.015	10.0

Table 1.

Chemical composition of tested superalloy (wt. %, bal. Ni).

Experimental samples casted from IN 713LC and MAR 247 LC for cyclic oxidation test had cylindrical shape with dimensions of 14 × 5 mm. Their microstructure was consisted of the γ matrix strengthened by γ´ phase with the shape of cuboidal particles of Ni₃(Al, Ti) as coherent precipitates and complex carbides. Three types of samples for two kinds of superalloys were used for cyclic oxidation tests [39]:

• samples of IN 713 LC and MAR 247 LC without coating,

• samples of IN 713 LC and MAR 247 LC with aluminide coating and,

• samples of IN 713 LC and MAR 247 LC with Si modified aluminide coating.

Aluminide coatings were applied by the “out-of-pack” method and Si modified aluminide coatings were made by the method of “pack-cementation”.

Disc samples in ceramic bowl (Figures 9 and 10) were placed into induction furnace with the temperature 1100°C. After 23 h, the tested samples were taken from furnace and they were immediately exposed to the cooling process in the air at room temperature for 1 h and then the mass changes were measured. This description represented one oxidation cycle. Mass changes of samples were written down after each second cycle and the photos were done.

Figure 9.

Samples of IN 713 LC with and without coatings before testing.

Figure 10.

Samples of MAR 247 LC with and without coatings before testing.

Gravimetric weight changes of cyclically tested superalloys with coatings followed weight changes during oxide formation and the results were used to compare the resistance of superalloy/coating systems to cyclic oxidic loading in a high temperature environment as well as to choose the most suitable combination for application in practice [39].

Macroscopic views of all samples after 3rd, 10th and 20th cycles are visible in the Figures 11–13. As for uncoated superalloy MAR 247 LC test of cyclic oxidation was interrupted after 13th cycle due to continuing decrease of weight.

Figure 11.

Overall view on all samples after 3rd cycle of testing.

Figure 12.

Overall view on all samples after 10th cycle of testing.

Figure 13.

Overall view on all samples after 20th cycle of testing.

The resulting graph of samples made from IN 713 LC with and without coatings is shown in Figure 14. From this picture, it is clear that IN 713 LC with unmodified aluminide coating had lower life-time than without coating and it is very surprising finding. Weight changes after each oxidation cycle for MAR 247 LC samples with and without coatings are visible in the Figure 15. We can see that life time of MAR 247 LC without coating is very low and structural components made from this type of superalloy without coating could not be used for practice application in the oxidation environment. On the basis of results in the Figure 15 we can see that life time of samples from MAR 247 LC superalloy with aluminide Al and AlSi coatings is practically the same.

Figure 14.

Mass change data for IN 713 LC with and without coatings achieved from thermogravimetric analyses.

Figure 15.

Mass change data for MAR 247 LC with and without coatings achieved from thermogravimetric analyses.

The resulting graph of all tested samples is shown in Figure 16. This picture revealed that samples from IN 713 LC without coating after 24-h cycles in environment of 1100°C show good resistance and their mass change maintain above zero limit. Samples from MAR 247 LC superalloy with both aluminide Al and AlSi coatings appear to be the most acceptable selection of combination relating to superalloys/coating.

Figure 16.

Weight changes after each oxidation cycle for all tested samples.

2.2.3 Microstructural research of DACs in the oxidation environment

2.2.3.1 Microstructural examination of IN 713 LC without coating

Microstructure of IN 713 LC before cyclic oxidation test (Figure 17) consisted of the γ-solid solution strengthened by γ´ phase with the shape of cuboidal particles of Ni3(Al, Ti) as coherent precipitates (dark particles in Figure 18) and complex carbides (bright particles in Figure 17).

Figure 17.

Microstructure of IN 713 LC without coating before testing.

Figure 18.

γ-solid solution of IN 713 LC with γ´ phase.

After 8 cycles of cyclic oxidation testing the microstructure of IN 713 LC without protective coating degraded as we can see in Figure 19. It was found that γ´ particles near the surface region dissolved to the solid solution and bellow is the area of larger γ´ coarsened ones. The thickness of oxide scales and area of dissolved precipitates under the surface reached on average 5.3 μm and 32.6 μm, respectively.

Figure 19.

(a) Microstructure of IN 713 LC (b) Detail focused on the after 8 cycles of testing microstructure near surface region.

From Figures 19 and 20 is clear that degradation of uncoated samples from 8 to 18 cycles continued in the sense of formed the thicker area of oxide scale (9.9 μm) and dissolved precipitates (46.4 μm). Mechanical properties of investigated alloy mainly depend on morphology, size and volume fraction of gamma prime strengthening particles. Increasing of cycle number resulted in decreasing of the γ´ phase volume fraction what is not acceptable considering strength of the alloy.

Figure 20.

(a) Microstructure of IN 713 LC (b) Microstructural appearance of after 18 cycles of testing the region under oxide surface.

Growth of oxide scale and area of dissolved precipitates continued with increase in test cycles. After 28 cycles of cyclic oxidation was observed growth of oxide scale deeper into superalloy and the oxide thickness reached 12.9 μm. The thickness of dissolved precipitates was 73.4 μm (see Figure 21a and b). 38 test cycles caused growth of oxide scales up to 17.8 μm and area of dissolved precipitates to 100.1 μm (Figure 22a and b).

Figure 21.

(a) Microstructure of IN 713 LC (b) Oxide scales after 28 cycles after 28 cycles of testing of testing.

Figure 22.

(a) Microstructure of IN 713 LC (b) Oxide scales under surface of after 38 cycles of testing alloy after 38 cycles of testing.

Chemical contents of the γ-solid solutions of elements in nickel (dissolved precipitates areas) in all samples after cyclic oxidation tests were roughly identical. Contents of Al were from 7.5 at. % to 8.5 at. % in the samples after 8, 18, 28 and 38 cycles; Ni ranged between 73.1 at. % and 73.4 at. %; Cr was from 15.1 at. % to 15.3 at. %. Bellow this zones were areas of coarse precipitates which contained a bit higher amount of Al (9.1 at. % – 9.8 at. %), a lower content of Cr (12.8 at. % – 13.4 at. %) and an identical content of Ni. Cross sections of oxide scales of all samples after testing (Figure 23) revealed two zones; inner oxide layer and outer oxide layer. It was found that the outer oxide zones were composed of lower contents of Al (about 59 at. %) compared to the inner zone (about 96 at. % Al), higher contents of Cr (about 4.9 at. %) with respect to that in the inner ones (about 0.6 at. %) and a considerably higher content of Ni in the outer layers (about 32 at. %) compared to that in the inner layers (about 2.5 at. %). From these examinations results that the degradation of uncoated samples from IN 713 LC superalloy started with a creation of oxide scales on the surface by exhausting of aluminum and chromium from the alloy inside and formation of an area of dissolved γ´ precipitates.

Figure 23.

(a) Oxide scales after 8 cycles (b) Oxide scales after 18 cycles.

Next cyclic oxidation cycles resulted in linearly increasing thicknesses of oxide scales and area of dissolved precipitates, i.e., solid solution γ without strengthening γ´-phase. The γ-solid solution grown about 5 times faster than oxide scales (Figure 24). The continuously growth of oxides and their mass gain on the surface of samples without coating has shown that although the oxides underwent spallation processes during each cooling cycle of testing, the surface of alloy was doped by aluminum from the bulk after each cycle (see Figure 14), whereas strength and creep properties of the alloy decreased. Decreasing mass gain started from the 29th cycle when the forming of oxides continued mainly in the inside of the alloy whereas the surface was subjected to the spalling process. X-ray diffraction analysis of phase of the sample after 38 cycles of cyclic oxidation revealed a presence of protective oxide phases on the base of Al and Ti in the mixture with NiAl₂O₄ spinel phase (Figure 25). NiAl₂O₄ could be formed in the process of cyclic oxidation by the reaction of NiO and Al2O₃.

Figure 24.

Thicknesses of oxides and gama-solid solution after 8, 18, 28 and 38 cycles of testing.

Figure 25.

X-ray analysis of phases on the surface of sample after 38 cycles of cyclic oxidation.

2.2.3.2 Microstructural examination of IN 713 LC with Al coating

Figure 26a and b represent the microstructure of superalloy IN 713 LC with Al coating deposited by the “out-of-pack” method before the cyclic oxidation test. The upper layer of the coating was composed from the NiAl solid solution and the inner one contained Ni₃Al solid solution with particles of Cr, Nb, Ti and Mo. The thicknesses of upper and inner sublayers were 32.8 μm and 17.7 μm, respectively. Figures 27–30 show changes of microstructures of samples after 8, 18 and 28 cycle, respectively. They revealed that the degradation of the protective coating continued very fast and a zero gain was reached after 28 cycles of testing (see Figure 14). This mass change of 0 mg/cm² compared to the original state of sample indicated that the reserve of aluminum was exhausted and the protected function of the coating was stopped. The oxide thicknesses on the sample surfaces after 8, 18 and 28 cycles were 6.68, 4.97 and 4.68 μm, respectively.

Figure 26.

(a) Microstructure of IN 713 LC with Al coating before testing (b) Detail on the coating before testing.

Figure 27.

Microstructure of IN 713 LC with Al coating after 8 cycles of testing.

Figure 28.

Microstructure of IN 713 LC with Al coating after 18 cycles of testing.

Figure 29.

Microstructure of IN 713 LC with Al coating after 28 cycles of testing.

Figure 30.

Intermediate zone between the inner layer of coating and the superalloy of sample after 28 cycles of testing.

The oxidation behavior at the beginning of the test was controlled by the formation of a fast-growing layer of a mixed oxide. At cyclic oxidation times longer than 12 cycles this rather thick layer started to spall off. In accordance with the unit mass change (Figure 14) up to 12th cycle the protective oxide scales formed on the surface of samples prevailed over the effects of spallation and starting from the 13th cycle the spalling processes become dominating. Both the weight gain in Figure 14 and the measured oxide layer thickness demonstrate this behavior. A crack in the coating of sample after 8 cycles was found as a result of thermal stresses (Figure 31). A major problem of such coatings is that the coefficient of thermal expansion of the alumina layer differs from the coefficient of expansion of the base material.

Figure 31.

Crack in the Al coating after 8 cycles of testing.

During thermal cycling, stresses arise between the aluminum oxide top layer and the coating material. The resulting oxide layer is relatively brittle and tends to crack and peel off, exposing the fresh surface to a damaging atmosphere. This repeating process consumes the aluminum in the coating. When the aluminum level in the coating drops below a certain point, the coating becomes ineffective as an alumina generator and the protective benefits of the coating material are lost.

Process of spalling can be examined by the view on the surfaces of samples after 8, 18 and 28 cycles (Figure 32a–d).

Figure 32.

(a) The surface of IN 713 LC with Al coating before testing (b) The surface of IN 713 LC with Al coating after 8 cycles of testing (c) The surface of IN 713 LC with Al coating after 18 cycles of testing and (d) The surface of IN 713 LC with Al coating after 28 cycles of testing.

Dark areas in Figure 32b–d responded to protective alumina scales and white places to the Ni₃Al phase with an amount of Cr, Nb, Ti and Mo. This is in accordance with EDS analyses of elements from surfaces and cross sections of each of samples. An example of this examination from the sample after 18 cycles we can see in Figure 33 where phase 1 corresponds to the oxides and phase 2 to the solid solution Ni₃Al.

Figure 33.

EDS analyses of phases from surface and cross section of samples after 18 cycles of testing.

2.2.3.3 Microstructural examination of IN 713 LC with AlSi coating

Si modified aluminide coatings were made on the surface of IN 713 LC samples by the method of “pack-cementation”. Silicon was added to the aluminide coatings to improve their oxidation resistance and the oxide scale adherence and, as a consequence, the oxidation rate was lower. After the pack cementation process some of Ni atoms were replaced by Si in the solid solution of aluminide phase. This resulted from the X–ray diffraction phase analysis (Figure 34) since no Si-containing phases were found. During pack aluminizing, the superalloy samples to be coated were placed in an air-tight retort containing a mixture powder of aluminum and silicon activated with ammonium chloride and an inert Al₂O₃ filler which prevented the sources form sintering. The box was then inserted into a furnace and heated in a protective atmosphere. The pack cementation process is essentially an in situ chemical vapor deposition (CVD) coating process. The coating achieved by this method and subsequent heat treatment was composed from outer and small inner sublayers (Figures 35 and 36). According to the results of the EDS analyses, the element distribution of outer and inner coating layers and the substrate is displayed in Table 2. The element distribution confirmed that the outer layer contained a large amount of small precipitates containing atoms with higher atomic numbers such as Cr, Ti and Mo which are present in the substrate. This is a typical feature for inward diffusion of Al and Si. The small diffusion underneath suggests a limited degree of the outward diffusion of Ni. The total thickness of the coating including the diffusion zone was approximately 77 μm. The X-ray diffraction phase analysis by JEOL JDX-7S from the sample surface revealed that the matrix of the coating was the Ni₂Al₃ phase in which particles of niobium aluminide were present (Figure 34).

Figure 34.

The X-ray diffraction phase analysis from the surface of the sample IN 713 LC with the AlSi coating.

Figure 35.

The cross section of sample IN 713 LC with AlSi coating before cycling oxidation test.

Figure 36.

Individual parts of AlSi coating on superalloy.

	Al-K	Si-K	Ti-K	Cr-K	Fe-K	Ni-K	Nb-L	Mo-L
Outer layer	50.1	1.9	0.3	8.0	0.5	37.8		1.3
Inner layer	31.4		1.4	17.6	0.2	45.0	1.2	3.1
Substrate IN 713 LC	11.9		0.7	12.9		72.7		1.8

Table 2.

Element distribution in the IN 713 LC superalloy with the AlSi coating in at. %.

Microstructures of IN 713 LC samples with AlSi coating after testing of cyclic oxidation are visible in Figures 37–41. Specimens after 8 cycles of testing changed their microstructural view. The white particles visible in Figure 37a are composed from elements as Si, Cr, Ti, Nb and Mo as found on the basis of the EDX analyses. The content of Al in the coating was about 50 at. % and after 8 cycles decreased to 29 at. % and then remained constant up to 38 cycles. The content of Ni exhibited an opposite behavior and increased from the original 38 at. % to 64 at. % after 8 cycles and then remained constant up to 38 cycles.

Figure 37.

(a) Microstructure after 8 cycles (b) Crack in the coating after 8 cycles.

Figure 38.

Microstructure after 18 cycles.

Figure 39.

(a) Microstructure after 28 cycles (b) Cracks after 28 cycles.

Figure 40.

Microstructure after 38 cycles.

Figure 41.

View on the thicknesses of coatings after 8, 18, 28 and 38 cycles of testing.

The content of particles based on Si, Ti, Cr, Nb and Mo maintained practically identical during all the testing cycles. As one can see from Figure 41, the coating thickness grew larger after 8 cycles of testing and then remained practically unaffected up to 38 cycles.

There were practically no changes in the thickness of surface oxide scales during the cycling. Figure 42a–h represent surfaces of coatings before and after 8, 18, 28 and 38 cycles of testing. Contents of Al and Ni before testing changed from about 51 at. % and 40 at. % to 92,5 at. % and 5,8 at. % after testing, respectively. Alumina phase on the surface was subjected to the processes of buckling and scaling off and white particles were exposed. A large amount of small cracks on the top of coating (Figure 42c,e,g) represents an evidence of such a mechanism. Since the concentration of white phases during the cycling did not change too much it means that new alumina phases were created simultaneously with the scales. This behavior along with the small mass gain to the 11th cycle of testing and a very small weight loss up to the end of the cycling (Figure 14) confirm a good resistance of IN 713 LC with AlSi coating to the thermal cycling.

1. Surface of IN 713 LC with AlSi – 0 cycles

2. Sample after 8 cycles.

3. Detail of the sample after 8 cycles

4. Sample after 18 cycles.

5. Detail of the sample after 18 cycles

6. Sample after 28 cycles.

7. Detail of the sample after 28 cycles

8. Sample after 38 cycles.

Figure 42.

Surfaces of Inconel 713LC samples before cyclic oxidation and after 8, 18, 28 and 38 cycles of testing.

2.2.3.4 Microstructural examination of MAR 247 LC without coating

MAR 247 LC belongs to widely used high temperature nickel-base superalloy. It is a cast polycrystalline material used especially for turbine blades and discs. The chemical composition (Table 1) of the alloy has, in comparison with IN 713 LC, a larger content of heavy elements as tungsten and tantalum which act as the most efficient matrix hardeners and promote creep strength. Cobalt in MAR 247 LC has only a small direct influence on strengthening but it affects the solubility of elements in the matrix solid solution and often raises the temperature of solidus which can lead to the larger amount of precipitates at low to medium temperatures [28]. Carbon is present at higher concentration (0.15 wt.%) than in the IN 713 LC (0.04 wt.%) and it combines with reactive elements such as titanium, tantalum, hafnium and tungsten to form MC carbides. During processing or service these carbides can decompose to other forms as M₂₃C₆ and M₆C which are rich in chromium, molybdenum and tungsten. Excessive amount of tungsten together with molybdenum and chromium leads to the formation of so-called topologically close-packed (TCP) phases. Various semi-empirical models are used for balancing composition of superalloys to avoid forming of these undesirable phases. The material before testing and without coating was subjected to a hot isostatic pressing and, subsequently, it was heat treated by two steps to reach the microstructure shown in Figure 43a. This initial microstructure of the MAR 247 LC alloy consisted of γ Ni-rich solid solution containing a dispersion of γ′ precipitates, carbide particles, and γ/γ′ eutectics as depicted in Figure 43b.

Figure 43.

(a)View on the microstructure of MAR 247 LC without coating before testing (b) Phases in the microstructure of the MAR 247 LC alloy.

As one could see in Figure 15, the superalloy MAR 247 LC without coating revealed a very poor resistance to thermal cycling. Mass loss started immediately after the 1st cycle and continued during subsequent 12 cycles when test was stopped. Microstructures of this alloy after 8 cycles of testing well reflect that behavior – see Figures 44 and 45.

Figure 44.

View on the structure of MAR 247 LC after 8 cycles.

Figure 45.

Subsurface zone of the sample after 8 cycles.

Precipitation-hardenable superalloys usually have a good oxidation resistance in oxidizing atmospheres within their normal range of service temperatures. Exposure to high-temperature environments can cause changes in the alloy composition near the surface. Figure 45 shows the changes in the subsurface microstructure formed under a high temperature environment. As we can see in Figure 45, the surface of the alloy is covered by oxidation scales based on aluminum (area 5 in Figure 45) but an internal oxidation is also visible. Preferential oxygen attacks on carbide phases were observed. Since certain elements as aluminum or chromium are consumed by the scale layer, the bulk composition can become depleted. Subsurface zone (area 2 in Figure 45) is composed from a solid solution of elements in nickel with small amounts of aluminum and without γ´ coherent precipitates. Chemical content of phases highlighted in the Figure 45 is in Table 3.

MAR 247 LC	Al-K	Ti-K	Cr-K	Co-K	Ni-K	Zr-L	Mo-L	Hf-L	Ta-L	W-M
Area 1	7.5	1.3	9.7	10.7	66.6	0.2				4.0
Area 2		8.5	1.8	3.3	16.9			41.7	27.8
Area 3	5.4	0.9	11.8	11.8	65.2	0.1	0.5			4.3
Area 4	51.0		0.7	1.6	7.6			39.1
Area 5	95.8				4.2

Table 3.

Element distribution in MAR 247 LC without coating after 8 cycles of testing in at. %.

A comparison of both surfaces, i.e. before and after cycling oxidation, is presented in Figures 46–48. The EDX analyses confirmed that the chemical composition of the alloy on the sample surface (Figure 46) without cycling is identical with that declared in Table 1. The specimen surface after 8 cycles is covered by oxidation scales as can be seen in Figures 47 and 48. The surface is mainly rich in aluminum, chromium and hafnium. Tungsten and tantalum are present in carbides (white particles in Figures 47and48). The detail depicted in Figure 48 shows an attack of carbides along grain boundaries even more clearly than one could see from Figure 45.

Figure 46.

Surface of MAR 247 LC.

Figure 47.

Surface of MAR 247 LC without coating without coating before testing after 8 cycles of testing.

Figure 48.

Oxide scales on the surface of MAR 247 LC without coating after 8 cycles of testing.

2.2.3.5 Microstructural examination of MAR 247 LC with Al coating

Figures 49 and 50 show cross-sectional images of the sample MAR 247 LC with Al coating before testing of cycle oxidation at 1100°C in the air environment. The two-layered coating (Figure 50) includes the upper NiAl layer which serves as aluminum reservoir and the inner diffusion layer of solid solution of Ni with a smaller content of Al and carbide particles on the basis of Cr, W, Hf and Mo. The chemical content of elements in the highlighted areas (1-4) in Figure 50 achieved from the EDX analysis is in Table 4. The average thicknesses of upper and inner parts are 24.9 μm and 18.2 μm, respectively.

Figure 49.

SEM micrograph of the microstructure of the sample MAR 247 LC with Al coating before testing at low magnification.

Figure 50.

Cross-section of sample MAR 247 LC with Al coating before cycle oxidation testing at 1100°C.

MAR 247 LC	Al-K	Ti-K	Cr-K	Fe-K	Co-K	Ni-L	Mo-L	Hf-L	Ta-L	W-M
Area 1	49.5		2.0	0.3	5.3	43.0
Area 2	25.9	1.9	11.9	0.3	10.0	43.8	0.6	1.1		4.6
Area 3	10.5	1.0	10.4	0.3	10.6	61.9			0.9	4.6
Area 4		28.0	2.7		2.2	9.8	1.6	8.6	34.4	12.6

Table 4.

Element distribution in MAR 247 LC with Al coating before testing in at. %.

The view on the microstructure of the specimen after 8 cycles of testing at 1100°C in the air for 8x23 hours is in Figure 51. The surface is covered by an oxide film and, with respect to this diffusion process, the upper part of the coating is depleted of aluminum. The content of aluminum in the matrix of upper layer dropped from 49.5 at. % before testing to 27.4 at. % after 8 cycles of oxidation. In this zone one can observe grains with a higher concentration of nickel (65 at. %) than in the matrix (in Figures 51–53 described as “Grains of Ni”). These grains are bigger in samples after 18 and 28 cycles of oxidation. The microstructure view on the sample after 38 cycles (Figure 54) is different from the previous ones.

Figure 51.

Microstructure of MAR 247 LC.

Figure 52.

Microstructure of MAR 247 LC with Al coating after 8 cycles with Al coating after 18 cycles.

Figure 53.

Microstructure of MAR 247 LC.

Figure 54.

Microstructure of MAR 247 LC with Al coating after 28 cycles with Al coating after 38 cycles.

Coating upper layer of the sample after 38 cycles has two phases as in previous samples (Figures 51–53) but proportion of these phases is changed. Grains of Ni are extended to a large extent that they can be assumed as matrix of the sample after 38 cycles and second dark phases mean unchanged grains of the upper layer matrix of samples after 8, 18 and 28 cycles.

Moreover, needles of topologically close packed phases (TCP) were observed. Presence of refractory elements in the superalloys provide strength benefits from solid solution hardening but a tendency for an alloy instability due to a formation of TCP phases (or secondary reaction zone phenomena) is high. The formation of these phases has a detrimental effect because of their brittle nature and depletion of Ni-rich matrix from strengthening elements [29]. The thickness of coating after 8 cycles rose up twice in comparison with original state and then it kept changeless up to the end of thermal cycling.

Appearance of the sample surface after 8 cycles of oxidation is in Figure 55 where a spallation of small oxide segments can be seen. Similar images were observed on samples after 18, 28 and 38 cycles (Figure 56).

Figure 55.

General view on the surface of sample after 8 cycles of oxidation and detail from spalled area.

Figure 56.

Spalled area and cracks on the oxide surface of sample after 38 cycles.

The spallation started already at very early stages of oxidation. Stress generation within the oxide during its growth and its release due to cracking in the scale or creep of the substrate metal led to the spallation process and exposure of the substrate. In spite of the loss of protective scales during the exposure by a substantial number of 38 cycles some new scales were reformed. On the basis of data achieved from the thermogravimetric analyses (Figure 15) the MAR 247 LC with Al coating after 38 cycles of cyclic oxidation can be still considered to be a good protective system.

2.2.3.6 Microstructural examination of MAR 247 LC with AlSi coating

Pack powder mixture for codepositing Al and Si on MAR 247 LC superalloys by the pack cementation process was used to form AlSi diffusion coating. Silicon (similarly to chromium) was added to the aluminide coating to enhance the resistance against oxidation and sulfur-carrying gases i.e. hot corrosion. However, application of Si is limited to small amounts because silicon bears the risk of forming low-melting phases in nickel-base superalloys [23]. Overall microstructure view on the sample MAR 247 LC with AlSi coating is visible in Figure 57 where coating is two-layered. The upper part has average thickness of 94.4 μm and the inner diffusion part is much smaller. Detailed chemical compositions achieved by EDX analyses of individual marked zones (see Figure 58) is in Table 5.

Figure 57.

General cross-section view on the sample MAR 247 LC with AlSi coating before testing.

Figure 58.

The coating at higher magnification.

MAR-247 + AlSi	Al-K	Si-K	Ti-K	Cr-K	Co-K	Ni-K	Mo-L	Hf-L	Ta-L	W-L
Area 1	44.7	4.2	0.4	7.6	5.9	34.8	0.2			2.3
Area 2	47.2	0.6	0.6	6.7	5.8	36.6	0.3			2.3
Area 3	45.2	0.0	0.2	4.9	6.4	41.2	0.2			1.9
Area 4	29.2		1.7	16.5	9.2	37.1	0.7			5.2
Area 5	10.9		1.1	10.4	10.5	61.3			0.7	4.3
Area 6	11.0	9.3	20.6	3.5	2.7	13.4	1.0	4.6	22.6	11.2

Table 5.

Element distribution in MAR 247 LC with AlSi coating before testing in at. %.

Silicon is the most placed in region 1 and it has thickness about 12.5 μm. Lower regions 2 and 3 have hardly any silicon and chemical content of others elements is similar to the region 1. The upper part of this coating has a few of precipitates based on Si and refractory elements (point 6 in Figure 58). Thin diffusion coating part (area 4 in Figure 58) is formed from NiAl solid solution with precipitates on the base of refractory elements and its chemical content is very similar to the diffusion zone of Mar 247 LC with Al coating (see Table 4). EDX analyses show that representation of silicon gradually falls from the surface to the diffusion zone of the coating. Under this coating, we can observe dendritic microstructure of the superalloy. Ni-rich γ matrix with γ´ semi-coherent precipitates contains MC carbides characterized by different shapes (from discrete blocky precipitates of a diversified shape and size to the shape known as Chinese script) and borides (Figure 57). MC carbides precipitated in the final stage of solidification via eutectic reaction with γ-matrix in the interdendritic areas. The secondary M₂₃C₆ carbides are very fine and exist at the grain boundaries.

Testing of samples on cyclic oxidation at the temperature 1100°C in the air environment showed that this system MAR 247 LC superalloy/AlSi coating had excellent resistance to cyclic oxidation. The amount of aluminum in the upper part of coating decreased after first cycles on the average from 47 at. % to 26 at. % and then up to 38 cycles of testing remained at the same value. This means that aluminum reservoir is sufficient to form protective oxides on the surface and the oxidation rate is slow. From the point of view of engineering design, kinetics of oxidation is very important because it gives an estimate of design life of system superalloy/coating. The microstructures of samples after 8, 18, 28 and 38 cycles are visible in Figures 59–62. The thickness of the coating of sample after 8 cycles rose sharply 3.3 times and then the thicknesses of all tested samples remained the same. Topologically close packed phases emerged after 8 cycles of testing, their amounts gradually increased to 38 cycles. This behavior was although observed in samples of MAR M247 LC with Al coating. From the point of view of other phases, the microstructure remained without significant changes up to 38 cycles of testing of the cycle oxidation. Chemical contents of phases found in MAR 247 LC with AlSi coating after 38 cycles were similar to the phases in MAR 247 LC with Al coating.

Figure 59.

Microstructure of MAR 247 LC.

Figure 60.

Microstructure of MAR 247 LC with AlSi coating after 8 cycles with AlSi coating after 18 cycles.

Figure 61.

Microstructure of MAR 247 LC.

Figure 62.

Microstructure of MAR 247 LC with AlSi coating after 28 cycles with AlSi coating after 38 cycles.

Surfaces of MAR 247 LC with AlSi coating samples after 8, 18, 28 and 38 cycles of testing represent Figures 63–69. We can observe that protective alumina oxide after 8 cycles spalled out and the layer composed from solid solution of Al in Ni and precipitates based on refractory elements revealed (Figure 64). Some of these spalled parts had a bigger size (the white particle in Figure 64). Next cyclic oxidation caused new alumina scales formation and white particles seen in Figure 65 were smaller. Creation of protective alumina scales continued (see Figure 66) and after 38 cycles of testing protective oxides (Figure 68) covered almost all the surface of coated superalloy.

Figure 63.

Surface of MAR 247 LC with AlSi coating after 8 cycles (left BSE image, right SE image).

Figure 64.

Detail from the surface displayed in Figure 63 (left BSE image, right SE image).

Figure 65.

Surface of MAR 247 LC with AlSi coating after 18 cycles (left BSE image, right SE image).

Figure 66.

Surface of MAR 247 LC with AlSi coating after 28 cycles.

Figure 67.

Part of the surface displayed in Figure 66 in higher resolution.

Figure 68.

Surface of MAR 247 LC with AlSi coating after 38 cycles.

Figure 69.

Part of the surface displayed in Figure 68 in higher resolution.