18.2.1 General

The need for coating systems in the domain of high temperature arose by the 1960s when improved performance criteria could not be met adequately in spite of new materials with superior physical, mechanical, and metallurgical properties. Operating efficiency and production economy had to be considered and an improvement sought.

In this section, a brief overview of the main categories of coating systems is provided.

18.2.2 Diffusion Coatings

Nowadays the coatings for high temperature applications include mainly diffusion coatings, overlay coatings, and thermal barrier coatings (TBCs). The protective character of diffusion coatings attributes to the protective nature of the Al₂O₃, Cr₂O₃, and SiO₂ scale formed, respectively, on the aluminides, chrominides, or silicides at elevated temperatures. Diffusion coatings have first been developed and still are the most used coating. Aluminide diffusion coatings are proved to be the cost‐effective solution for high temperature oxidation, which are used widely for protecting turbine blades and vanes. The properties of the aluminide coating depend on the process methodologies used to deposit the coating, the substrate composition, and the subsequent treatment. The coating deposition rate and morphology depend on the process temperature and time. Processing temperature influences the rate of diffusion, at which alloy elements may diffuse and the metallurgy of the surface compound may form; thus, it is a critical parameter in the processing and manufacturing of diffusion coatings. The coating time at temperature defines the thickness of the coating formed during the diffusion step. The thickness of the coating is also a main factor of protecting property of the coatings. Two basic mechanisms typify the diffusion coatings, depending on whether the main diffusing species is aluminum diffusing from the coating to substrate of the base metal of the substrate alloy diffusing outward to the coating layer. These coatings are usually produced by pack cementation, out‐of‐pack cementation, and chemical vapor deposition (CVD), which involve the diffusion of a predominant element such as aluminum to form “diffusion coating” layers. It should be noted that coatings are typically fine grained with more grain boundaries compared with substrate materials. Diffusion rates in coatings are therefore expected to be faster than in substrate of similar composition. The three main diffusion coating mechanisms will be analyzed in the next section. Here, the diffusion coating mechanisms will be briefly considered, taking the diffusion aluminide coatings as an example.

Aluminizing is achieved by two different processes that are based on the activity of aluminum in the gas phase and the aluminizing temperature. The low activity high temperature (LAHT) process is a “one‐step” method with outward diffusion of Ni to form β‐NiAl coating. On the other hand, the high activity low temperature (HALT) process concerns the inward diffusion of aluminum, typically giving rise to a δ‐Ni₂Al₃ coating, which requires subsequent heat treatment to convert it to β‐NiAl. This is because δ‐Ni₂Al₃ is a brittle phase that has lower mechanical property than β‐NiAl. As might be expected, the mechanisms by which the coatings grow are also different. For the LAHT process, the aluminum activity is insufficient for it to be the predominant diffusing species, and accordingly, coatings form by diffusion of Ni from the alloy substrate into the region of the coating. For the HALT process, the aluminum activity in the gas phase and at the surface of the coating is high enough to facilitate the inward of Al into the alloy substrate. The growth mode is an important factor when considering coating integrity, since coatings formed by outward Ni diffusion can trap diluent particles (alumina) brought by aluminizing packs. While in the coating β‐NiAl phase layer of LTHA, much unwanted phases can precipitate dispersively, which decrease the mechanical property of the coating. Which coating process is selected depends upon a number of features, e.g. heat treatment specifications for the substrate alloy, applications, nature of packs available, integrity issues, etc.

In LAHT aluminizing (low aluminum content in the pack [1.2–1.5 wt% of Al], 900–1150 °C), the aluminum is deposited on the surface, but at reduced rate, and nickel simultaneously diffuses outward to the surface. Then a β‐NiAl surface layer forms. The interdiffusion coefficient in the β‐NiAl phase of the Al–Ni system varies strongly with composition. At 1050 °C, the minimum value of interdiffusion coefficient is not at the stoichiometric composition, but displaced to the low‐Al side by 1–2 at.% (that is 48–49 at.% of Al). Ni diffuses predominantly in low‐Al β‐NiAl.

The rapid nickel diffusion to the surface, coupled with the low aluminum activity at the surface, effectively holds the surface aluminum content close to 50 at.%. Hence, LAHT coatings are near stoichiometric at the surface and Ni rich within the substrate β‐NiAl phase. This structure ensures outward of Ni transport and the outward growth of the coating microstructure. Another important result is that slowly diffusing elements Ta, W, Re, and Mo from substrate are unable to diffuse to form significant concentration levels in the outwardly growing β‐NiAl. The outer zinc zone of this coating therefore appears much “cleaner” and is free from such precipitates as observed in the HALT coatings. The limited solubility of Cr, as well as Ta, Mo, and W, in the β‐NiAl phase produces precipitates of topologically close‐packed (TCP) phases, forming an interdiffusion zone (IDZ) under the “clean” β‐NiAl phase layer. Thus, the IDZ includes the complicated TCPs and β‐NiAl phases. This coating composition and structure are less dependent on the alloy composition. The further diffusion processes between such aluminide coatings and Ni‐based alloys are determined by the aluminum activity in the β‐NiAl phase. As long as this is lower than the equilibrium activity among β, IDZ, and γ + γ′ matrix, Ni will diffuse out of the base alloy into the β phase. The coating grows further outward.

The kinetics of the formation of aluminide coating on pure Ni by LAHT has also been studied, and this mechanism is illustrated in Figure 18.2.

In respect to the HALT theory interpreted above, if the activity of Al is low, Ni preferentially diffuses out through the coating and combines with Al to form NiAl. The coating grows outward. In this case, Ni diffuses faster than inward diffusion of Al. Pack particles are entrapped in the coating. These particles are assumed to be dragged out by the diffusion of Ni.

As previously mentioned, HALT coatings are processed at lower temperature (700–950 °C) as the first step, with a higher aluminum content (1.7–2.7 wt% of Al). The substrate material (γ‐Ni + γ′‐Ni₃Al) reacts with the depositing aluminum, forming a surface layer of δ‐Ni₂Al₃ over the layer of β‐NiAl. The diffusivity of Al through δ‐Ni₂Al₃ is very high, and substantial amounts of Al can be forced deeper and deeper into the material. In δ‐Ni₂Al₃ phase layer, the diffusivity of Ni is near zero, while Al diffuses rapidly; thus the formation of surface layer of δ‐Ni₂Al₃ results from the inward diffusion of aluminum. Substrate alloying elements such as W, Mo, Ta, and Re are selectively diffusing at the coating/substrate interface. Once the solubilities of alloying elements of W, Mo, Ta, Re, and Cr are saturated, many precipitates rich in these elements will occur for their limited solubility to β‐NiAl phase and distribute in the entire coating structure. After a second step of heat treatment at higher temperature (950–1100 °C), Ni is able to diffuse outward to transform the brittle δ‐Ni₂Al₃ phases into Al‐rich β‐NiAl. This step is also usually combined with the heat treatment required to recover substrate properties.

Now we have preliminary concepts of LAHT and HALT, and then we can primarily establish whether the coating belongs to one process or the other. However, in the actual manufacturing there are all kinds of operations by adding different amounts of Al and at different temperatures. Therefore, the resulting coating could be significantly different and have very complex structures.

Coating–superalloy interdiffusion is principally responsible for the phase transformations, oxidation behavior, and degradation of the mechanical properties of the coating. Moreover, the diffusion behavior in the multicomponent coating layers is very complicated due to the interactions among the components. Thus, diffusion is an important factor to be considered when designing β‐NiAl coatings.

18.2.3 Overlay Coatings

Diffusion‐type coatings, used successfully on early gas turbines, were tied to the substrate composition, microstructure, and design. Later, some changes were introduced: (i) in superalloy composition, such as reduction in Cr and increase in other refractory metals; (ii) in microstructure, by castings with more segregation; and (iii) in design, by air cooling and with thin walls (which introduced higher thermal stresses). These changes required coatings that were much more independent of the substrate. Overlay coatings met this necessity.

Overlay coatings also overcome the process restrictions encountered in diffusion coatings, especially the variants, viz. Cr/Al, Ta + Cr, or the Pt aluminides, all of which give better stability and oxide hot corrosion resistance than Al alone. MCrAlY compositions (M = Ni, Co, Fe alone or in combination; Y represents oxygen‐reactive elements such as Zr, Hf, Si, and Y) are the principals in the series of overlay coatings developed by the electron beam evaporated physical vapor deposition (EBPVD) technique for multiple load use. More than 4 million aerofoils have been successfully processed by EBPVD.

MCrAlY overlay used in gas turbines are usually Ni and/or Co with high Cr, 5–18% Al, and Y addition around less than 1% for stability during cyclic oxidation. They are multiphase alloys with ductile matrix, e.g. gamma Co–Cr, containing a high fraction of brittle phase, e.g. beta CoAl. The Cr provides oxidation hot corrosion resistance, but too much Cr affects substrate phase stability. The success of most overlay coatings is the presence (and perhaps location) of oxygen active elements such as Y and Hf, which promote alumina layer adherence during thermal cycling, giving increased coating protectivity at lower Al levels. Y mostly appears along grain boundaries if a MCrAlY is cast but is homogeneous if plasma sprayed. Thus MCrAlY with 12% Al are more protective than the more brittle diffusion aluminides with 30% Al.

Overlay claddings deposited by hot isostatic pressing (HIP), electron beam evaporation, or sputtering methods are diffusion bonded at the substrate/coating interface, but the intention here is not to convert the whole coating thickness to NiAl or CoAl. There is thus more freedom in coating composition, whose properties can be maximized to the type required. Composition based on NiCr, CoCr, NiCrAl, CoCrAl, NiCrAlY, CoCrAlY, FeCrAlY, and NiCrSi have been successful in gas turbine engines. They are generally alumina formers with only 10% Al unlike the 30% in nickel aluminide coatings. Cr increases the Al activity, which allows this feature. Higher Al levels cause brittleness and a higher DBTT, and the Al levels are generally held below 12% (5–10% preferred). The coatings are also more ductile than NiAl and CoAl and can be rolled and bonded by HIP. In general, NiCrAlY give best results against high temperature oxidation, while CoCrAlY are best for hot corrosion.

18.2.4 Thermal Barrier Coatings

TBCs are ceramic coatings applied over metal substrates to insulate them from high temperatures. They consist of zirconium oxide, ZrO_2, stabilized by about 8 wt% yttria, Y₂O₃, or magnesia, MgO. They are applied over a bond coat that is usually an MCrAlY overlay coating but can be a diffusion coating. The ceramic layer is typically 5–15 mils thick, and the overlay bond coat is usually 3–8 mils thick.

The bond coat improves the adhesion of the ceramic top coat by reducing the oxide buildup underneath the ceramic. (The ceramic is porous to oxygen and will spall off when the oxide formed on the bond coat is sufficiently thick.) The bond coat is applied in the same manner as for overlay coatings.

The ceramic layer is normally plasma sprayed in an air environment. This layer can also be made by electron beam physical vapor deposition (EBPVD), which is used on aircraft engines. The ceramic layer made by EBPVD is columnar grained and is more resistant to spalling than the plasma‐sprayed TBC.

Another type of ceramic layer that has been developed is called a segmented TBC. It has cracks through the thickness of the ceramic, which are perpendicular to the substrate surface. These cracks provide increased resistance to spallation, similar to EBPVD ceramic layers. The segmented TBC is made by a proprietary plasma spray process, and very thick ceramic layers can be made.

The TBC protects the metal by acting as an insulator between the metal and the hot gases. The thermal conductivity of the ceramic is 1–2 orders of magnitude lower than the metal. Although thin, the TBC can significantly reduce the metal temperature provided that the metal component is air cooled. (The air cooling provides a heat sink.) Furthermore, the ceramic has a higher reflectivity than the metal. This means that more of the radiative heat is reflected away, which is important for combustion hardware.

During startup and shutdown, the TBC improves the thermal fatigue life by reducing the magnitude of the temperature transients the metal is exposed to. A 10 mil thick TBC on airfoils in experimental aircraft engines has achieved more than 300 °C reduction of metal temperature. During steady‐state operation, the TBC lowers the temperature of the underlying metal, thereby improving its durability. It also reduces the severity of hot spots.

The principal failure mechanism of TBCs is the spallation of the ceramic layer. Spallation is caused by the synergistic interaction of bond coat oxidation and thermal cycling. The oxidation occurs at the interface between the bond coat and the ceramic. As more oxide forms at this interface, the ceramic layer is more likely to spall.

18.3.1 General

There are three major processes by which the diffusion coatings can be formed: CVD, pack cementation, and out‐of‐pack cementation, which involve the diffusion element such as aluminum (chromium or silicon). Overlay coatings can be applied by CVD, physical vapor deposition (PVD), high velocity oxygen/flame (HVOF), or plasma spray. Plasma spray is one of the most widely used processes today. TBCs are usually deposited by PVD, EBPVD, and plasma spraying techniques. All these processes and a few more advanced techniques for producing high temperature coatings are described in this section.

18.3.2 Chemical Vapor Deposition

CVD is a process wherein a stable solid reaction product nucleates and grows on a substrate in an environment where a vapor‐phase chemical dissociation or chemical reaction occurs. It uses a variety of energy sources, viz. heat, plasma, ultraviolet light, etc., to enable the reaction and operates over a wide range of pressures and temperatures. CVD is a long‐established, economically viable industrial process in the field of extraction and pyrometallurgy. Some of the well‐known processes are:

• The Mond process:
  
  18.1
  
  
• The Van Arkel process:
  
  18.2
  
  
• The Kroll process:
  
  18.3

The method has also been used to form freestanding, simple, and complex shaped articles from metals that are not very amenable to fabrication, e.g. W:

18.4

The technology of CVD took on new dimensions with a fresh view taken at the deposition aspects of the process. This transition of emphasis from extraction to deposition of the CVD phenomenon has made CVD an important sector of coating technology for producing new materials and coatings with improved resistance to wear, erosion and corrosion, good thermal shock resistance, and neutron absorption characteristics. Components from micro‐ to macro‐sizes are now coated by CVD. In the deposition area of CVD, the lamp and electronics industry uses many of the earlier processes to obtain longer service life and better performance for a number of products. The CVD industry took a bigger leap when demands arose for new high temperature materials with greater high temperature corrosion resistance and strength in all the sectors using gas turbines and other propulsion hardware, in the military science, engineering, and aviation fields. In parallel, CVD technology has had rapid application in the semiconductor industry, the many branches of fuel and energy – nuclear, fluidized bed, oil, solar and chemical industries and the tool industry. Solid‐state devices; electrode materials; thermionic devices; erosion‐, corrosion‐, and wear‐resistant components; fission barriers; and many more applications can be listed in CVD technology.

It would be confusing and misleading to list a number of processes as CVD methods in as much as different high temperature coatings cannot be merely termed as diffusion coatings. The minor or major parameters that influence the actual achievement of the coating itself are those that facilitate categorization. Otherwise, conventional CVD, plasma‐assisted CVD (PACVD), laser‐assisted CVD (LACVD), pack, vacuum pack, pack‐slip, vacuum pack‐slip, CCRS, or RS may all be just CVD. Similarly, diffusion coatings can be obtained by first hot dipping, spraying, cladding, fused salt electrodeposition, electrophoresis, plasma, or all PVD, and all the above “CVD” methods. What gives the distinction to any process?

The obvious factor that emerges is that a basic CVD process has a number of limitations, e.g. the types of reactions possible and available, and the minimum substrate temperature that has to be maintained. This results in restricting any eventual substrate/deposit configuration and control. Improved techniques stretch the reactant–product spectrum and can also (i) lower the temperature; (ii) reduce the hazards and mismatch in the morphological aspects, viz. nucleation and growth; (iii) alter the physical–mechanical properties favorably; and (iv) provide better bonding.

CVD technology offers the following favorable aspects to semiconductor industry:

1. A wide range of silicon epitaxy thicknesses and resistivity.
   
2. A low cost, highly productive method of depositing polysilicon, Si₃N₄, and high temperature SiO₂.
   
3. A proven technology for passivating silicon devices with low temperature deposition of silicon oxide and silicon nitride.

A wide range of metallic and inorganic coatings can be produced by CVD. In many instances CVD coatings exhibit unique properties, which are difficult to produce by other methods.

A CVD process was seen as a thermochemical process in which the substrate is regarded as a collector of the deposit, and the substrate/deposit interaction is both undesirable and unnecessary for deposit growth. However, for a CVD deposit, the post‐deposition state is a mandatory parameter to meet the service requirements; substrate/deposit interaction to some degree is expected. Further, heat treatment, and therefore diffusion, is often a step following the primary CVD process (or even during deposition, since the substrate is heated). Thus, substrate/deposit interaction always occurs and is also one of the ways deposit consolidation can be achieved.

A CVD reactor is the apparatus‐equipment configuration that carries out the CVD process. CVD reactors are categorized on the basis of the temperature and pressure under which they operate. A CVD reactor can have a cold wall or a hot wall depending on the experimental optimization parameters. When the reactor wall surrounding the heated substrate is comparatively cool relative to the temperature of the substrate, it is a cold‐wall reactor. The substrate is heated directly, and product deposition occurs mostly on the heated component. In a hot‐wall reactor, the deposition product can nucleate both on the reactor wall and the heated substrate, though to different degrees, if a temperature difference prevails across the wall substrate area and zone. The substrate may be independently heated or not, but the reactor itself is heated in order to facilitate the process.

The CVD reactor assembly consists of three components: (i) the reactant supply system, (ii) the deposition systems, and (iii) the reactant–product retrieval, recycle, and disposal system. The entire reactant supply (using solid and/or gaseous reactants) has to be carefully balanced before introduction into the reactor system, predicted to reach the required partial pressures at the temperatures of reaction and deposition. The primary features to consider concerning the substrate are the substrate/deposit adherence and growth. CVD reactions can be exothermic but mostly are endothermic. This could result in substrate cooling. On a production scale, large volumes of gases are handled, which can induce convective coating. Both these situations have to be compensated. Substrate heating optimization involves achievement of the correct temperature profile and coating exposure time. The substrate/reactor parameters in a deposition system should be designed to achieve a laminar flow. The input can be parallel, perpendicular, or angular to the substrate depending on the substrate holder design. The movement (rotation, lateral, etc.) of the substrate within this region also compensates to some extent any nonuniformity in the deposition.

CVD occurs through a chemical reaction; therefore, the thermodynamics of the system that can drive the chemical reaction, the basic chemical kinetics to provide information on the reaction rate, and mass and heat transfer profiles that are influenced by the reactor and substrate size and design, require consideration.

Mass transport of a CVD system links diffusion control, chemical kinetics, and fluid‐flow dynamics. Diffusion growth can be ascertained from chemical kinetics by rate constant of growth = A exp (−E/kT), where E is the activation energy, k is the Boltzmann’s constant, and T is the temperature. With increase in substrate temperature, the growth rate changes at a transition temperature from a predominant kinetics to a diffusion control, the latter being mostly under mass transport influence (Figure 18.3).

Many models of CVD reactors and their kinetics have appeared in the recent literature, but space restrictions impose our departure to the next coating process.

18.3.3 Pack Cementation

Pack cementation is an extended CVD method in which the substrate to be coated is introduced in a retort and surrounded by a powder mixture, which is composed of the source alloy(s), an activator, and an inert filler. The retort is then heated to the desired coating temperature usually under inert or reducing environment. In these conditions, the source element reacts with the activator, forming a gaseous transporting species. This diffuses through the pack to the substrate surface where it decomposes, allowing the metallic element to be deposited to penetrate into the substrate by solid‐state diffusion.

As for all diffusion coatings, the microstructure is highly dependent on the substrate, which is coated. For this reason, process conditions have to be optimized for each material.

The first “cementation process” was presented by Allison and Hawkins in 1914, who deposited Al on iron and on steel. But it is only since the 1960s that this process stimulated interest because of the development of coatings for the protection of gas turbine blades, especially those made of Ni‐based superalloys. The process has then been extended to Co and Fe‐based alloys. Concerning the elements deposited, these are mainly Al followed by Cr and Si. Nevertheless, codeposition of two or three elements became possible with the theoretical understanding of the process. The extensive works on Cr–Al codeposition published by Rapp on steels (1989, 1994, 1997) and by Young et al. on nickel‐based alloys (da Costa et al. 1994a,b, 1996; Gleeson et al. 1993) are the major examples, but development of Cr–Si (Rapp 1996), Al–Si (Wynns and Bayer 1999), Ti–Al (Weber 2004), and Ti–Si–Al (Rosado and Schütze 2003) coatings can also be mentioned. Eventually, more recent investigations seem to focus on the addition of reactive elements such as Y, Ce, or Hf by codiffusion.

The principles of the pack cementation technique changed remarkably little since 1914, but the composition and the quality of the substrate–deposit systems have been optimized in reply to the requirements of a large range of service applications.

From an industrial point of view, the pack cementation technique has been used very extensively for nearly 60 years for coating gas turbine blades (Mévrel et al. 1986). Nowadays it is estimated that on first‐stage gas turbine blades, more than 80% of all coated airfoils are coated by pack cementation (Goward 1998).

The coating of larger pieces has been the subject of some research investigations and even leads to the development of a few patents, showing that coating even inside a tube is possible by this technique (Baldi 1980; Wynns and Bayer 1998). During the 1990s the US company Alon developed several industrial processes involving the pack cementation technique. The purpose was to protect ethylene cracker furnace components from carburization and coke formation (Kurlekar and Bayer 2001). With Alon’s facilities, 15 m long furnace tubes can be coated with Al, Cr, and/or Si on an industrial scale, producing 100 000 linear meters of tubes per year (Alon). After one year service in an ethane cracking furnace, the coating reduced the coke formation by a factor of 2, and post‐exposure analysis revealed that the coating microstructure was essentially unchanged (Ganser et al. 1999).

Typically the powder pack mixture consists of three components:

• A metallic source: A fine powder of the element(s) to be deposited on the surface of the substrate (Al, Cr, Si, Ti, or combinations of these elements).
  
• A halide salt activator as, for example, NaCl, NH₄Cl, AlCl₃, NaF, NH₄F.
  
• An inert filler powder (e.g. Al₂O₃) preventing the powder mixture from sintering at high temperature.

Table 18.1 lists different packs that can coat various elements with the corresponding temperatures and heat treatment durations.

Figure 18.4 illustrates the major chemical reactions involved in the general pack cementation process. In order to allow a better understanding of the following paragraph, the reactions are written for the particular example of aluminizing of iron activated by NH₄Cl.

The whole process relies on the formation of gaseous halides according to the general reaction:

18.5

AlCl₃, AlCl₂, AlCl, Al₂Cl₆, and Al₂Cl₄ are the chlorides involved in a chloride‐activated aluminizing process.

The partial pressures of each of the gaseous halides formed are established by their thermodynamic stability, which varies with the process conditions: composition of the pack, type of activator, temperature, pressure, and type of inert or reducing environment.

In the case of aluminum deposition, AlCl₃ is the major halide formed at low temperature, whereas at higher temperature the activity of AlCl becomes higher.

Once they are formed, the halide molecules diffuse through the gas phase to the substrate (e.g. iron) surface, where they adsorb and decompose owing to the general reactions:

• Disproportionation:
  
  18.6
  
  
• Displacement/exchange:
  
  18.7
  
  
• Direct reduction:
  
  18.8
  
  
• Dissociation:
  
  18.9

By using alkaline or earth alkaline halide activators as, for instance, NaCl, KCl, or CaCl₂, the precipitation of the activator salt has also to be considered (Ravi et al. 1989):

• Precipitation:
  
  18.10

The aluminum formed at the surface of the substrate of Reactions 18.6–18.10 can then diffuse into the solid substrate, forming the desired coating.

The predominance of the Reactions 18.6–18.10 depends, first of all, on the stability of the gaseous halides involved. The deposition particularly occurs by disproportionation Reaction 18.6 when, first, the vapor pressure of the substrate halide is low and when, secondly, the coating element has higher and lower halides of comparable vapor pressures in order to set atomic aluminum‐free. When the vapor pressure of the substrate halide becomes comparable to the gaseous species of the coating element, the contribution of the exchange Reaction 18.7 becomes important. The latter reaction even gets undesirable if the vapor pressure of the substrate halide becomes higher than the source supplier halide, as it would lead to a significant metal loss and to porous coatings. In a NaCl‐activated pack, this phenomenon can occur during chromizing or siliconizing of iron, whereas Reaction 18.7 stays minimal during aluminizing (Kung and Rapp 1989).

The reader may, at this point, easily catch the compromise that has to be found. Because halide molecules must diffuse through the gas phase from the pack to the substrate, the coating composition depends on the gaseous halide activity and their stability or their ability to decompose at the substrate surface. Hence, the formation and the decomposition of the gaseous halide have to be optimized at the same time in the pack and at the substrate surface, respectively.

Furthermore, it is a prerequisite that the thermodynamic activity of the incorporated element is always lower at the surface than in the pack (Gupta and Seigle 1980). This activity gradient drives the gas‐phase diffusion of the halide molecules from the pack to the substrate surface. As a consequence, a desired coating composition cannot be obtained by simply using a master alloy of the same composition (Rapp 1989).

Moreover, the concept of “major depositing species” has been defined. This corresponds to the gaseous species that is responsible for the major part of the deposition. In the case of a Cr–Al codeposition process by a chloride‐activated pack, Rapp (1989) and da Costa et al. (1994a,b) showed that although the vapor pressure of AlCl₃ is several orders of magnitude higher than that of Cr halides and other Al halides, the codeposition is possible by optimizing the process conditions so as to get comparable vapor pressures of AlCl and CrCl₂. Indeed, AlCl₃ is too stable and does not decompose enough at the substrate surface. The Al transport occurs via AlCl. This is thus considered as the major transporting species for Al, whereas CrCl₂ is the major transporting species for the deposition of Cr (Rapp 1989).

The formation of the coating can be described in the light of diffusion couples (Wachtell 1974). The coating actually corresponds to the IDZ between the substrate element(s) and the deposited element(s). The driving force for the solid‐state diffusion is the activity gradient between the pack/coating and the coating/substrate interface.

The aluminizing of iron can thus be treated as a binary diffusion couple, involving pure Al and pure Fe. As a consequence, the phases formed in the coating should follow the FeAl binary phase diagram (Figure 18.5).

As can be seen from this diagram, the aluminizing of iron can form several intermetallic phases, FeAl₃, Fe₂Al₅, FeAl₂, FeAl, or Fe₃Al as well as solid solution of Al in Fe. The effective formation of these phases is however controlled by the powder composition, the temperature, and the duration of the process. These parameters especially control the diffusion fluxes, which determine the final coating structure.

Goward and Boone (1971) distinguished the low and the high activity pack structures (Figure 18.6), as already discussed in Section 18.2.2. Their investigations considered the aluminizing of a nickel‐based superalloy. Two powder pack mixtures differing in the nature of the Al source (pure Al or Ni₂Al₃) were investigated at about 800 and 1100 °C, respectively. The observations lead the authors to consider that the high activity pack (with the powder containing pure Al) structure is characterized by a high aluminum content at the surface with the coating phase Ni₂Al₃. The underlying phases down to the coating–substrate interface followed the Ni–Al phase diagram with decreasing Al content. For the low activity pack (with the powder containing Ni₂Al₃), the coating is single‐phased β‐NiAl, which is directly in contact with the substrate.

The difference was attributed to the diffusion fluxes:

• In the high activity pack conditions, the coating formation was attributed to exclusive inward Al diffusion, the diffusion of Ni being minimal at low temperature.
  
• In low activity packs, the coating formation was interpreted as single Ni outward diffusion. This last mechanism also requires higher temperatures to allow diffusion of Ni.

One example of nickel aluminide diffusion coating produced by LAHT pack cementation process is illustrated in Figure 18.7. The coating consisted of a well‐defined to “clean” layer with NiAl as its major phase and a large diffusion zone underneath.

An example of SEM‐BSE (backscatter electron) image of LTHA coating on Ni‐based superalloy CMSX4 (this is a typical second generation Ni‐based single‐crystal superalloy characterized by the replacement of most of the Ti with Ta, by relatively high Co and low Mo content) is shown in Figure 18.8. This coating was deposited using an aluminizing pack containing 2 wt% Al at 900 °C and was heat‐treated for 2 hours at 1120 °C and then for 24 hours at 845 °C. The β‐NiAl phase layer is uniform in thickness, and many precipitates (white small particles) distributed dispersively in the entire β‐NiAl phase layer. Squillace et al. (1999) has also analyzed clearly the coating structure after the first step of LTHA. Three layers are visible: the inner layer has a striated appearance, the outer layer is equiaxed with many inclusions, and the middle layer is featureless. The coating layer has deposited very high amount of Al. In the book entitled High Temperature Coatings by Bose (2007), it was presumed that higher inward diffusion of Al with lower outward Ni diffusion has avoided the formation of Kirkendall porosity and also eliminated the embedded pack particles.

Although mixed types of mechanisms can occur as a result of varying pack activities, substrate compositions, and temperature, during practical application of the coating process, it is still of great use nowadays to relate the observed structures to two archetypes (Figure 18.6).

The aluminizing of iron shows the same kind of behavior. At 900 °C, a three‐layer structure with Fe₂Al₅, FeAl₂, and FeAl is formed in high activity packs, and a single‐phased FeAl coating in low activity packs (Levin et al. 1998).

With the help of thermodynamic calculations performed with a computer software, the activities of the gaseous halides can be determined considering all pack reactions together. Such considerations allow the optimization of the pack concentration and process conditions needed to form the desired coating phase. More precisely, the activities of the gaseous halide depend on the reactants involved in the halide formation and thus on the type of metallic source mixed to the pack powder. Therefore, the direct dependence of the coating structure with the master alloy determines whether the coating deposition occurs in a high or low activity process.

The activator is also a determining reactant involved in Reaction 18.5 too, from which it can be said that an unstable activator results in high halide partial pressures. Moreover, concerning the choice of cation for the activator, some researchers recommend the use of ammonium salts, because of their decomposition at high temperature. Indeed, the NH₃ produced in Reaction 18.5 is not stable above 350 °C decomposing by the following reaction:

18.11

which shows a double advantage:

• The formation of hydrogen enhances the deposition of coating element(s) through the reduction Reaction 18.8.
  
• The precipitation Reaction 18.10 does not occur, leaving the chlorides behind for reacting with the master alloy.

Concerning the thermodynamic aspects of pack cementation, it is also necessary to note that as a consequence of the temperature dependence of the Reactions 18.5–18.11, the process temperature directly influences the formation of gaseous halides and their decomposition, as well as the values of the gaseous and solid‐state diffusivities that also influence the kinetics of the cementation process.

Like all CVD processes, pack cementation involves gas–solid reactions. For many cementation coatings, the coating process can be considered as six steps in series (Cockeram and Rapp 1995). Usually at the beginning of the process, the flux of master alloy element delivered by halide molecules on the surface is slower than the limiting flux of this element dictated by the solid‐state diffusion. Consequently, the beginning of the process follows the gas diffusion kinetics, and the growth of the depleted zone is rate limiting. At the end of the transition regime (regime after which equilibrium in the pack is achieved), the aluminum flux delivered by the gas phase equals the maximum uptake into the solid phase. At that moment, the kinetics limitation of the process changes from gas diffusion to solid‐state diffusion.

18.3.4 Out‐of‐Pack Cementation

The out‐of‐pack or over‐pack process operates in a manner similar to pack cementation except that the components to be coated are suspended either above the pack or below the pack (vapor generator) retort.

The transport of aluminum species from the vapor phase to the substrate occurs by gas‐phase diffusion and by solid‐phase diffusion of aluminum into the substrate to form the aluminide phases. The former increases the surface concentration of aluminum in the coating, while the latter decreases it. The surface composition of the coating tends to reach a steady‐state value in a short time after the commencement of the process. In the vapor‐phase aluminizing, the rate of transport of aluminum to the substrate is much faster than the solid‐phase diffusion of aluminum into the substrate. Thus the composition of aluminide coating is decided by the kinetics of the solid‐phase diffusion. The coating process is divided into a number of steps. They are:

1. Formation of the aluminum subchlorides by the reaction of the aluminum metal or alloy and the aluminum chloride vapor.
   
2. Transport of the subchlorides to the substrate by gas‐phase diffusion.
   
3. Reaction leading to the deposition of aluminum at the substrate surface.
   
4. Diffusion of aluminum into substrate with the formation of the coating consisting of different intermetallic phases.
   
5. Diffusion of the reaction products from the substrate back to the reactor.

Steps (a) and (c) are very fast at the operating temperature; therefore, the thickness of the coating process is controlled by step (b) the vapor transport and (d) the solid‐phase diffusion. Step (e) decides the purity of the coating.

The schematic diagram of the out‐of‐pack process is given in Figure 18.9. The coating vapors are transported to the components by an inert carrier gas. Plumbing is designed that the vapors can access to both external and internal surfaces of the components. The retort is inserted into a furnace and held at the desired temperature for the selected duration.

This approach results in a much cleaner and uniform coating for very complicated geometry components, with no entrapped pack particles.

In industry, for example, components are aluminized using a proprietary SIFCO vapor‐phase aluminizing process. A picture of the facility is shown in Figure 18.10a. At the beginning of the vapor‐phase coating process, the components, activators, and fillers are loaded into a coating box, and then the coating box is covered with a retort. During the vapor‐phase process, the coating box is heated to an elevated temperature of about 1080 °C, held for three to four hours, and then cooled down to room temperature. The temperature profile during the aluminizing process is shown in Figure 18.10b.

18.3.5 Physical Vapor Deposition

PVD is a coating process in which the atoms of the coating material are evaporated and then deposited onto the surface of the part (Boone 1986). The process must be carried out in a vacuum chamber. Coating materials can be pure metals, metallic alloys, or ceramics. Most materials can be coated by the PVD process.

Electron beams are usually used to vaporize the coating material, in which case the process is called “electron beam physical vapor deposition,” or EBPVD. A schematic of the EBPVD process is shown in Figure 18.11. The electron beams are focused by magnets on the ingot of the coating material, causing it to melt. Due to the vacuum environment, a vapor cloud of the coating material forms over the molten ingot pool. This vapor rises and deposits on the part to be coated. The evaporated atoms move in a straight line from the molten ingot to the part. Complete coating coverage is obtained by rotating and tilting the part in the vapor cloud.

Diffusion normally occurs during the PVD process at the substrate–coating interface. This diffusion forms a metallurgical bond of the coating with the substrate. However, post‐deposition diffusion treatments are commonly performed to improve the bond and the coating. Glass bead peening is also frequently used to close areas of unbonded columnar grain structure to produce a fully dense structure.

The PVD process is a line‐of‐sight process, so internal and recessed surfaces cannot be coated. A good vacuum is required to prevent oxidation of the reactive vapor cloud and the substrate surface, since oxide layers can result in a nonadherent coating.

Because of the large part size of industrial gas turbine components and limited chamber size of PVD coaters, the cost of PVD coatings for industrial gas turbine parts can be quite high. These costs may decrease as larger PVD coaters become commercial.

18.3.6 Plasma Spraying

Plasma spraying is a process in which a powder of the coating material is heated and accelerated toward the part being coated by a plasma gas stream. Since it is a line‐of‐sight coating process, it is necessary to rotate and tilt the part and plasma gun during coating. Internal passages, such as cooling holes, cannot be coated by plasma spray. There are several books and reviews on plasma spraying, and some of them are listed in Futher Reading.

The plasma is generated from a high current arc with inert gases and powder particles, which is then accelerated toward the part, as shown in Figure 18.12. The particles strike the substrate surface at a high velocity. (Prior to spraying, the substrate surface is roughened by an abrasive process, such as grit blasting.) The particles cool very quickly on the roughened surface. The coating is attached to the substrate by the mechanical interlocking of the coating and the surface irregularities. In addition, chemical bonds can form, and post spraying heat treatments can be used for further interdiffusion of the coating and substrate.

Particles continue to be deposited at a high rate, such as a million particles per second. Each particle forms into the irregularities of the previous particles and the coating is built up. These solidified particles are called splats and cool very quickly. Voids form in the coating due to gas that is trapped by the plasma spray, and oxides are deposited due to particles that are oxidized. Sometimes particles are not heated enough and become embedded in the coating as spherical particles.

A post‐coating heat treatment is often carried out to ensure coating–substrate bonding through interdiffusion and to homogenize the structure of the coating. This is normally performed at 980–1090 °C in a vacuum or inert atmosphere for periods of 1–10 hours.

Plasma spraying can be performed in air at atmospheric pressure (air plasma spraying), in air with an inert shielding gas such as argon (shielded plasma spraying), or in a vacuum chamber. The vacuum processes are called vacuum plasma spraying (VPS) or low pressure plasma spraying (LPPS). A schematic of the air plasma spray and LPPS processes is shown in Figure 18.13. The LPPS process decreases contamination of the powder and base metal from oxidation and, therefore, substantially improves coating quality. The process also produces high density coatings with virtually no unmelted particles, and it allows the substrate to be heated without oxidizing. Shielded plasma spray can produce excellent coatings but requires special procedures and equipment to be used (Taylor et al. 1985).

When the plasma spraying is carried out in air or in an inert gas, the particles cool and slow down as they leave the plasma jet and collide with the gas molecules. If air is used, oxidation of the particles occurs. Inert gas greatly reduces, but does not completely eliminate, this oxidation.

18.3.7 High Velocity Oxygen/Fuel

Another method of spraying coatings, called high velocity oxygen/fuel (HVOF), has been developed. In this method, the powder is both heated and propelled by a high velocity flame onto the substrate. The high velocity allows denser coatings to be made than can be achieved by conventional plasma spray processes. Plasma spray produces a hotter flame than HVOF but has a lower particle velocity. The higher particle velocity of HVOF produces a denser coating, which improves the properties of the coating. HVOF is in commercial use for wear‐resistant coatings, which are made from ceramic powders. It is in developmental use for high temperature coatings.

The HVOF coating also has much less oxidation of the particles than air plasma spray. Thus, HVOF coatings can be produced that are far superior to air plasma‐sprayed coatings, despite the HVOF process being carried out in an air environment.

HVOF is being considered as an alternative to low pressure plasma spraying (PPS or VPS) (Kong et al. 2003) because of its lower cost. While LPPS coatings are still of higher quality than HVOF coatings, the quality of HVOF coatings may be sufficient for many applications.

18.3.8 Other Coating Processes

The most common methods of forming protective coatings are described in the previous subsections. But nowadays there are many more advanced processes with several advantages and applications that require, at least, a brief discussion. In general, these coating processes provide excellent properties such as strong bonding and low porosity and are smooth and homogeneous. Examples are as follows.

18.3.8.2 Ion Plating

A comprehensive review and bibliography (1963–1980) on ion plating is given by Mattox (1994), who invented the technique in 1964. The system is similar to the inert gas discharge used in sputtering, but the substrate is the cathode. Better adhesion results. High defect concentrations are put in the surface, and physical mixing of the coating and substrate surface occurs. A higher “throwing power” than vacuum evaporation methods is due to gas scattering, entrainment, and redeposition on the charged substrate. Ion plating is the most used method for steel tools coating. Irregular shapes are coated with little variation in adhesion, thickness, composition, and structure.

Ion plating with Al to protect U from oxidation has been successful in contrast to vacuum deposition that has been very unsuccessful due to the thin oxide that forms on U immediately after cleaning. In Ar glow discharge Al ion plating (Figure 18.14), the specimen surface is sputter cleaned by positive ion bombardment and is then held contamination‐free until and during the ion plating process. The high energy of ion collision with the surface heats it and promotes chemical reaction and interfacial diffusion (Bland 1968). Before ion plating U, the specimen was first deoxidized in 50% HNO₃ and electropolished in phosphoric acid. It was then suspended from a water cooled cold finger in a <5 × 10⁻⁶ Torr vacuum chamber, above a W filament of the same area as the specimen. Al evaporation clips were put on the W. The filament was grounded, and the specimen held at 0.5 KV potential (drawing typically 0.5 mA cm⁻²). Ar was leaked into about 10 μm Hg pressure and the high voltage applied, to clean the substrate with Ar⁺ ions. About 10 minutes later, the filament was slowly heated to deposit the Al. Then the filament, Ar, and high voltage were turned off in sequence. An axial magnetic field can be used to increase the electron path length and thus the ionization (Mattox 1994; Bland 1968; Thornton 1983, 1986).

Ion plating is good for wear‐ and erosion‐resistant coatings, due to its good adhesion characteristics, e.g. HfN for low friction and TiC for deformation and bonding resistance. For an incompatible coating and substrate, ion plating can be used as a “strike” for electroplating. The very good adhesion of ion‐plated coatings (much better than vacuum evaporated coatings) is attributed to the high energy of deposition. Two other advantages of ion plating are the good coverage of complicated shapes and the good structures and structure‐dependent properties, at low substrate temperatures.

18.3.8.8 Laser Surface Alloying

The application of laser in coating technology has developed rapidly since 1980 although the first attempt at laser surface alloying (LSA) was reported in 1964 (Cunningham 1964). Laser cladding by powder fusion progressed from research to development during 1974–1978 (Powell and Steen 1981). Now lasers are used as one of the main means of application of coatings, as a manipulated auxiliary in other deposition processes such as CVD and PVD techniques, as well as in a number of surface‐modifying methods such as cladding, welding, and spraying. The practical utility of lasers has stretched far beyond the area of coating technology itself.

The laser is a singular source of heat produced by means of highly energetic, coherent, monochromatic beam of photons that can be manipulated to operate in focused or defocused states. It is capable of attaining very rapid rates of heating and is clean, intense, and localized. It can be produced either in a continuous form or pulsed. A laser beam incident on a substrate can heat, melt, vaporize, or produce plasma depending on the manipulation of the controlling parameters. The principal sections of a laser are given in Table 18.2.

Section	Factors in design
Amplification of light	Optical quality; types of windows; means of heat removal; repair and efficiency; cost
Excitation	Power supply; I stage amplifying state auxiliary supply for cooling, pumps, etc.; weight, performance, efficiency; cost
Optical feedback system	Reflector components: mirrors, lenses, coating modulators; windows; mechanical and optical stability

The overall portability and efficiency of the equipment and its repair and maintenance costs have to be considered in its eventual selection for use. The wider application of lasers used in LSA coatings became possible when multi‐kilowatt CO₂ lasers came into production in 1970. Now CW CO₂ lasers are available from a few watts to many kilowatts power; CW Nd YAG lasers of 400 W and above are also being produced and used for laser surface processing.

LSA uses a laser beam as the heat source that is excellent due to the directionality, high intensity, and high spatial resolution of the laser beam. It heats a specific area very fast followed by faster cooling, resulting in a novel microstructure. The desired material can either be added simultaneously along with laser irradiation, or laser irradiation is carried out on the surface where the material has already been placed by some of the coatings already discussed. The schematic of the direct powder feed method, as well as the two step coating method, is shown in Figure 18.15.

The very rapid heating and cooling cycle allows a great extension of alloying possibilities, as conventional limits, such as solubility limits, seem to be overcome. Typical structures are extremely fine dendritic that contain the added elements or compounds in the interdendritic regions. The metallurgical stability of such structures is quite good and has been demonstrated by several applications.

18.3.8.10 Ion Beam‐Assisted Deposition

The versatility of the ion implantation process for the protection of metals has been further extended over recent years by combining a vacuum coating process with ion implantation. This can usually be carried out by straightforward adaptation of the implantation equipment. By the continuous or sequential processes of deposition and ion bombardment, it is possible to build up the coating to a thickness of a micrometer or even more. Furthermore, it is possible to choose an ion species that may react with this coating to convert it, partially or completely to a desired compound. At the same time, the interface between the coating and the substrate, always a plane of weakness, can be improved by ion beam mixing. Energetic ions, caused to pass through such an interface, can induce collisional mixing together with radiation‐enhanced forms of diffusive interpenetration to create a graded interface with excellent adhesion.

Among the field of ion‐assisted coatings (IAC), the term ion beam‐assisted deposition (IBAD) is reserved for the situation in which deposition and ion bombardment are carried out simultaneously. The extra energy imparted to the deposited atoms causes atomic displacements at the surface and the bulk as well as enhanced migration of atoms along the surface. These resulting atomic motions are responsible for improved film properties, including better adhesion and cohesion of the film, modified residual stress, and higher density, when compared with a simple PVD‐deposited coating. Refractory coatings such as (Si₃N₄) can be synthesized at very low temperatures when reactive ion beams or an evaporant is used.

Ion implantation and laser processing of surfaces can improve the cracking tolerance of structural alloys. Alumina remains crystalline if ion is implanted with Cr, Ti, or Zr and its hardness increases. It becomes less sensitive to fracture when scratched (McHargue and Yust 1984). Figure 18.16 shows the abrupt transition from tensile to compressive stress in evaporated Cr films on Xe ion implantation.

IBAD is a coating process highly versatile for metals, ceramics, semiconductors, and dielectrics. At present, it is also largely applied to the fields of lithography and etching.

18.4.1 Aluminides

The overall durability of any coating can be assessed as its specific resistance to degradation to corrosion, rupture, and thermal fatigue. Degradation of aluminide coatings is largely dependent on the coherence of the Al₂O₃ layer, which is largely influenced by the available Al activity from the substrate, and the transition metal it is alloyed with in the diffusion band.

Diffusion processes from the substrate always happen in conjunction with those occurring in the coating itself. At an early stage, Al in the aluminide coating is expected to diffuse into the substrate. But as degradation proceeds, the diffusion direction reverses with Ni from the substrate entering the coating that is being consumed. This situation is ideal for Kirkendall void formation (Sequeira and Amaral 2014) and causes physical dislodgement of coating if the void coalesces at the substrate/coating interface. Although Al₂O₃ is generally considered as an n‐type oxide, it is argued that it is so only at low . Thus, at 600 °C, crystalline alpha Al₂O₃ grows beneath an amorphous film of gamma Al₂O₃ by inward diffusion of O₂. At 1300–1750 °C, Al₂O₃ is n‐type up to = 10⁻⁵ and becomes p‐type at higher oxygen pressures.

Nickel aluminide is still the best aluminide system in the gas turbine field although aluminiding was first used to benefit steels. At the commencement of oxidation, the non‐protective gamma alumina forms and spalls easily. A more dense and adherent alpha alumina then forms, and as long as this remains undamaged, it confers protective kinetics on the substrate. However, a damage and repair cycle soon causes Al depletion, which affects the diffusion mode and composition of the NiAl below the scale. Below the critical Al concentration, the single‐phase NiAl now rendered Ni rich tends to precipitate the gamma primer Ni₃Al phase that causes NiO to react with Al₂O₃ to form spinels and also NiO on its own develops porosity during growth, thus increasing mass transport. Appearance of the gamma prime phase thus marks the onset of the rapid oxidation stage. In a mixed gamma and gamma prime, the gamma phase is more readily oxidizable and forms a voluminous, porous multiphased non‐protective scale.

Chromium is the principal metal linked with aluminiding, and most high temperature alloys rely on the oxides of Cr and Al for a consolidated resistance to degradation. The degradation in a chrome‐aluminide layer is influenced closely by the major element diffusion and degradation mode and the temperature effect on them. In general, nickel‐based alloys form Cr‐rich intermediate layers during aluminizing unlike Fe‐based alloys. Fe–25Cr was found to be the most resistant to sulfide‐forming and sulfur‐containing environments at temperatures less than 800 °C where Ni‐ and Co‐based alloys fail, the latter especially in Na₂SO₄ deposition conditions. Oxidation with thermal cycling and isothermal oxidation in exposure to fused Na₂SO₄ showed that the beneficial effect of Cr was to arrest spallation, and change in the coating microstructure was considered instrumental in restricting pitting of the chrome‐aluminide diffusion coating. It is postulated that the Cr‐enriched zone acts as a barrier to refractory metal diffusion, e.g. Mo, W, V, from the substrate and their subsequent oxidation. An internal sulfidation zone where the spinel (Cr,Al)₃S₄ can exist with a broad range of Cr/Al ratio has also been characterized.

A mathematical model has been developed to explore carbon diffusion through aluminided layers to form Cr₇C₃ (Marijnissen and Klostermann 1980).

The extension of aluminided coating life for degradation is rarely achieved by Al alone. At its simplest, it is a dual metal system like Ni–Al, Co–Al, Cr–Al, or Pt–Al, and its most complex a minimum of four and a maximum of six elements are formulated, usually of the general form MCrAl–X1 or MCrAlY–X2. These formulae are the most investigated: M = Ni, Co, Fe; X1 and X2 = Ce, Hf, Pt group, Si, Ta, Ti, Y, Zr, etc.; unusual additives such as Mn, Zn, and Cu have also been investigated. Investigations on gas turbine coating systems have mostly concerned NiCrAlX1–X2 and CoCrAlX1–X2 compositions, while the coal gas conversion systems have concentrated on FeCrAlX1–X2 systems.

Yttrium is the only additive that has been very widely investigated in all the three systems, viz. FeCrAl, NiCrAl, and CoCrAl, in both laboratory and rig evaluation tests. Other additives have been tested in addition or as an alternative. Degradation of M–Cr–Al–X systems is mostly the degradation of the M–Cr–Al scale. The mechanisms for the remarkable beneficial effect yttrium has on scale retention especially under thermal cycling conditions have been reviewed (Lacombe 1987; Stringer 1987).

The directions investigations need to take would be to see how metal–oxide bonds are established, how stress generation modes change, how scale plasticity gets affected at the metal/scale interface both for the substrate metal and the growing scale, how vacancy and void formations are affected, how preferential growth in directions toward and inward into the substrate occur, and how do cracks propagate (Stringer 1987).

Additives such as Y (and Ce, Hf, Zr, etc.) are referred to as active additions. The observed effects as summarized from various investigations are (Lacombe 1987):

1. Y influences the scale plasticity due to the alteration in scale microstructure.
   
2. Oxide pegs are grown, which anchor scale to substrate.
   
3. Mechanical property differences are graded between the scale and the substrate.
   
4. Scale growth mechanisms are influenced by influencing the diffusion parameters, thus reflecting on the transport and growth of the barrier layer.
   
5. Pore formations are eliminated by vacancy coalescence effects.
   
6. Chemical links between the scale and substrate are improved.

The beneficial effect by Y is not realized unless its levels are kept low as an additive. Y is shown to produce a series of oxides, viz. Y₃Al₅O₁₂, YAlO₃, Y₄Al₂O₉ apart from two spinels – Ni(Cr_0.7Al_0.3)₂O₄, Ni(Cr_0.45Al_0.55)₂O₄, and oxides of Y, Al, and Ni when its concentration is varied over 0–5.5 wt% in triode sputtered NICrAlY coatings on NiCrAl substrates, when oxidized in one hour cycles in air at 1100 °C. Convoluted surface oxides, with fractures, voids entrapped in the mid‐layers, and internal yttrium oxides, were also identified.

Y and Hf decrease the corrosion rate of NiCrAl‐, NiCrAlTi‐, and NiCrAlZr‐ alloys, but this beneficial effect does not extend in full over the hot corrosion environment in the range of 600–800 °C because of liquid product or liquid product deposit formations. The medium is particularly damaging to CoCrAlY coatings. Earlier theories postulated appear to be on lines similar to those discussed above (Bolshakov and Fedorov 1956). Co₂₂Cr₁₁Al_0.3Y deposited on IN‐738 superalloy by EBPVD and implanted Y and Hf on Co₂₂Cr₁₁Al were both found to degrade badly at 700 °C under hot corrosion conditions.

Using CoCrAlY or NiCrAlY as the main coating system, many permutations and combinations have been studied: Hf has been a prime additive element (Hocking and Vasantasree 1986), Zr and Si follow close, and several others have been tested, e.g. W, Mo, Mn, Ce, Zn, and Cu. Overlay and underlay coating systems are well known. The degradation in all these involves interlayer – adhesion – compatibility to thermal cycling and stress and of corrosion interdiffusion effects during hot corrosion, high temperature oxidation/carburization, nitridation, etc.

18.4.2 Silicides

Silicon is the third element chosen in hot corrosion systems for its ability to form a barrier layer to resist degradation, the protectivity relying on the formation on SiO₂. It offers the best melt fluxing resistance in hot corrosion environments at high . The pattern of growth and retention of SiO₂ alone or as an oxide associate is influenced by the manner in which Si is introduced into the system, as an element additive to form an alloy by slurry or CVD methods or by deposition or controlled oxidation to form SiO₂ or deposition as ceramics such as SiO_xN_y or Si₃N₄ on SiAlON. Although normally very stable, the degradation of SiO₂ can occur through cracking, peeling, spallation, or vaporization loss as SiO (Pettit and Goward 1983).

Si deposited on titanium substrates via silane (SiH₄) formed Ti₅Si₃, which was found to decrease the oxidation rate over the temperature range of 700–1020 °C in O₂. Scale formation was found to be a mixed oxide (TiO₂ + SiO₂) layer at temperatures below 875 °C, while above 900 °C, an outer TiO₂ and an inner, inward growing TiO₂SiO₂ developed.

Silicon‐containing protective scales can be formed by Ni–5Si, Ni–11.5Si, a chromia–silica scale on Ni–40Cr–5Si and Ni–40Cr–11.5Si, and an alumina–silica scale on CoCrAlY–5Si and CoCrAlY–5S and CoCrAlY–12.5Si. Co‐silicides melt above 1200 °C and no liquid silicates form below 1380 °C. But a solid‐state reaction between CoO and SiO₂ is said to occur by an unusual surface diffusion process (Figure 18.18) (Schmalzried 1974). None of the refractory Cr silicides melt below 1300 °C, but a Ni–50Cr with a protective Cr₂O₃ scale is pitted extensively when pressed in contact with a Si/SiC composite for 120 hours at 1150 °C in air. A NiCr silicide underlayer in the molten state at 1150 °C is envisaged (Ni‐5 wt% Si solid solution and silicides Ni₃Si and Ni₅Si₂).

Very good resistance to vanadic attack has been reported for silicon‐coated Ni–20Cr alloys at 900 °C with a reduction of as much as 80% in the corrosion rate. No breakaway effects were observed for more than 600 hours at 900 °C in 80V₂O₅–20Na₂SO₄ melts. The increased protection appears to be due to the development and retention of a Cr‐rich barrier layer beneath a Si‐rich surface with the slag showing little reaction with Si. Although Si coated by pack, vapor deposition, plasma spraying, or ion‐plating methods was uniformly effective, ion‐plated coatings were found to be the most reliable with good adhesion, uniform composition, and even thickness. The combined effect of molten sodium sulfate and vanadate is catastrophic in most cases as shown in studies on binary and ternary Ni–Cr(10–30 wt%)–X alloys (X = Al, Si, V) and IN‐738 (Sidky and Hocking 1987).

18.4.3 Thermal Barrier Coatings

Failure of TBCs could be initiated by poor bonding to the substrate matrix or the coating parameters such as porosity, microstructure–microcrack distribution, thickness, phase distribution and cohesive strength, or the relative thermal expansion mismatch and residual stress of the substrate–coating system. Ceramic sintering and bond coat inelasticity could also contribute to TBC degradation. Bond coat pre‐oxidation causes degradation by adhesion failure. The actual failure generally occurs within the ceramic layer near the bond coat and is believed to be due to slow crack growth and microcrack linkup with the ceramic.

Ceramic coatings degrade by spallation due to transient thermal stresses aggravated by salts that induce hot corrosion. For non‐oxide‐based ceramics such as silicon carbide and nitride, resistance depends strictly on the nature of secondary phases at grain boundaries due to the densification of additives, e.g. MgO, Y₂O₃, Al₂O₃, etc. Reaction‐bonded Si₃N₄ and self‐bonded SiC coatings were severely attacked in vanadic sulfate melts at 820–1100 °C when exposed to a residual oil‐fired environment in burner rigs. Tested in both crucible and burner rigs with melt compositions based on Na₂SO₄ with additions of NaCl, NaCl + V₂O₅, and NaCl + Li₂O (950 °C, 200 hours), reaction‐bonded Si₃N₄ was found to be highly dependent on the overall environmental conditions for its corrosion. Excellent resistance was observed in oxidizing and acidic media but very poor resistance in reducing and mildly alkaline environments. Several crystalline phases also occur on the ceramic surface during the corrosion reactions. Sintered SiC also suffered a similar attack in sulfate melts at 900 °C, less in oxidizing conditions but more in basic melts or slags especially with carbonaceous material. However SiC was inert in pure N₂, H₂, or H₂–H₂S mixtures at 900 °C.

A good bonding interlayer between the ceramic coating and the substrate superalloys is vital for TBC survival. Titanium carbide and nitride are well‐bonded to superalloys if the Ni₃Ti intermetallic phase is controlled. An excess produces a rough deposit and thus poor bonding. A marginal carburization is also a requirement for good bonding.

Without an interlayer, ceramic substrate interactions are inevitable. MgO, Al₂O₃, SiO₂ (as fused quartz), SiC, and Si₃N₄ were all found to react with a Ni‐based superalloy substrate, with the most severe being SiC over 700–1150 °C, with conditions at the interface held at minimized . Si and C diffused inward, forming silicides and carbides with the degree of reaction on the ceramic side being similar to that of the substrate below 900 °C and less above 900 °C.

High temperature systems are expanding and cheaper fuels are being used. More studies appear to be needed on hybrid substrate systems and in fused salt media with nitrates, carbonates, and phosphates in conjunction with sulfates and chlorides of Ca, K, Mg, and Na; additive effects also need further clarification. Degradation studies of multilayer coatings are complex and need further investigations. Poor reproducibility in processing TBC systems is felt to be reflected in their degradation assessments (Begley and Wadley 2012). Metastable phase systems and their degradation are an open field.

1. 18.1 What is the necessity of coatings for high temperature operating components?
   
2. 18.2 Which is the best classification for various types of coatings?
   
3. 18.3 What are columnar coatings, and under what conditions are such coatings formed?
   
4. 18.4 Why are CVD coatings called diffusion coatings? Explain with a specific example.
   
5. 18.5 What are the main limitations of CVD coatings?
   
6. 18.6 Under what conditions does one get low activity coatings and high activity coatings? Which of the two is better, and why?
   
7. 18.7 Differentiate between the properties of flame spray and HVOF coatings.
   
8. 18.8 Do atmospheric plasma coatings have sufficient porosity and bond strength? If not, how is it rectified?
   
9. 18.9 What are the advantages of the transferred arc process over atmospheric plasma spray?
   
10. 18.10 How is the IBAD method different from other ion beam‐based techniques?

Further Reading

1. Belmonte, M. (2006). Adv. Eng. Mater. 8: 693.
   
2. Bushan, B. (1999). Diamond Relat. Mater. 8: 1985.
   
3. Cha, S.C., Gudenau, H.W., and Bayer, G.T. (2002). Mater. Corros. 53: 195.
   
4. Fahr, A., Roge, B., and Thornton, J. (2006). J. Therm. Spray Technol. 15: 46.
   
5. Franchi, D. and Rostagno, M. (2005). Metall. Ital. 97: 21.
   
6. Gao, W. and Li, Z. (eds.) (2008). Developments in High‐Temperature Corrosion and Protection of Materials. Cambridge: Woodhead Publ. Ltd.
   
7. Gerdeman, D.A. and Hetch, N.H. (1972). Arc Plasma Technology in Materials Sciences. New York: Springer‐Verlag.
   
8. Goswami, B., Sahay, S.K., and Ray, A.K. (2004). High Temp. Mater. Processes 23: 211.
   
9. Hampikian, J.M. and Dahotre, N.B. (eds.) (1999). Elevated Temperature Coatings. Warrendale, PA: TMS.
   
10. He, J. and Schoenung, J.M. (2002). Mater. Sci. Eng., A 336: 274.
    
11. Herman, H. (1988). Plasma‐Sprayed Coatings. New York: Scientific American.
    
12. Jackson, M.R., Bewlay, B.P., Rowe, R.G. et al. (1996). JOM 48: 39.
    
13. Kawahara, Y. (2005). Corros. Eng. 54: 213.
    
14. Lang, E. (ed.) (1983). Coatings for High Temperature Applications. London: Applied Science Publishers.
    
15. Lang, E. (ed.) (1989). The Role of Active Elements in the Oxidation Behaviour of High Temperature Metals and Alloys. Amsterdam: Elsevier.
    
16. Lau, M.L., Strock, E., Fabel, A. et al. (1998). Nanostruct. Mater. 10: 723.
    
17. Malyshev, V.V. (2004). Prot. Met. 40: 525.
    
18. Mévrel, R. (1989). Mater. Sci. Eng., A 120: 13.
    
19. Miller, R.A. (1987). Surf. Coat. Technol. 30: 1.
    
20. Newkirk, J.B. (ed.) (1980). High Temperature Materials. Tel‐Aviv: Freund Publ. House.
    
21. Padture, N.P., Gell, M., and Jordan, E.H. (2002). Science 296: 280.
    
22. Pawlowksi, L. (1995). The Science and Engineering of Thermal Spray Coatings. New York: Wiley.
    
23. Peters, M., Leyens, C., Schulz, U., and Kaysser, W.A. (2001). Adv. Eng. Mater. 3: 193.
    
24. Picas, J.A., Forn, A., Igartua, A., and Mendoza, G. (2003). Surf. Coat. Technol. 174: 1095.
    
25. Rohr, V., Schütze, M., Fortuna, E. et al. (2005). Mater. Corros. 52: 12.
    
26. Sathiyanarayanan, S., Rajagopal, G., Palaniswamy, N., and Raghavan, M. (2005). Corros. Rev. 23: 355.
    
27. Sequeira, C.A.C. (2013). Corrosion and Materials in the Oil and Gas Industries (ed. R. Javaherdashti, C. Newaoha and H. Jan), 243. Boca Raton, FL: CRC Press.
    
28. Shah, S.I. and Glocker, D. (eds.) (1996). Handbook of Thin Film Process Technology. Bristol: IOP Publishing Ltd.
    
29. Stern, K.H. (1996). Metallurgical and Ceramic Protective Coatings. London: Chapman and Hall.
    
30. Stover, D., Pracht, G., Lehman, H. et al. (2004). J. Therm. Spray Technol. 13: 76.
    
31. Tellkamp, V.L., Lau, M.L., Fabel, A., and Lavernia, J. (1997). Nanostruct. Mater. 9: 489.
    
32. Verma, A.R.B. and van Ooij, W.J. (1997). Surf. Coat. Technol. 89: 132.
    
33. Warner, B.M. and Punola, D.C. (1997). Surf. Coat. Technol. 94–95: 1.