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High temperature coatings for gas turbine blades: A review
Surface
and Coatings
Technology,
HIGH TEMPERATURE A REVIEW R. SIVAKUMAR*
37 (1989)
COATINGS
139
- 160
139
FOR GAS TURBINE
and B. L. MORDIKE
Znstitut fiir Werkstoffkunde und Werkstoffiechnik, Technische Agricolastrasse 2, 3392 Clausthal-Zellerfeld (F.R.G.) (Received
BLADES:
Uniuersitdt
Clausthal,
April 9, 1988)
Summary High temperature coatings are formed to protect many engineering components from environmental degradation. The operating temperature and the nature of the corrodant dictate the choice of the coating. The coatings can be formed by different methods, but invariably the properties are dependent on the coating process. The life of the coating depends upon the mode of degradation in service. It is also influenced by the nature of the substrate. The need for coatings is well illustrated by their wide use in gas turbines. This review covers the development of the oxidation- and hotcorrosion-resistant coatings for turbine blades. The rationale behind such protection, starting from the nickel aluminide to future thermal barrier oxide coatings, is outlined. It is followed by a discussion of current trends in post-coating modifications to improve the properties of the coatings. The principles underlying the development of these coatings should serve as a useful guide in the choice of coatings for other high temperature applications.
1. Introduction High temperature oxidation and other reactions in corrosive atmospheres are of considerable importance to aerospace, nuclear and similar energy-related industries. The widely used high temperature engineering alloys are based on the transition metals iron, nickel and cobalt, whose oxidation resistance can be improved by adding sufficient amounts of chromium, aluminium or silicon to form continuous external oxide scales. These oxides, namely Cr203, A120, and SiO,, offer the best protection owing to their low growth rates. Among these oxides, Al,O, is the most suitable as Cr,O, scale is not useful above 1000 “C, where volatile CrO, *On
leave
from
Defence
Metallurgical
Research
Laboratory,
Hyderabad
500258,
India. 0257-8972/89/$3.50
0 Elsevier Sequoia/Printed
in The Netherlands
140
forms. In fact, significant evaporation has been detected at even lower temperatures of 850 - 900 “C [l]. This has been attributed to the high gas velocities in the gas turbine. At low oxygen pressures, SiOz scale is not stable and decomposes to gaseous products such as SiO. Further details on the oxidation properties of these alloys can be found in the review article by Wood and Stott [a]. Wagner’s theory can predict the minimum concentration required for the formation of continuous external scale [3]. However, such predictions are not yet feasible for the oxidation of complex alloys or even for the case of a pure metal exposed to a mixture of reacting gases. Table 1 lists the variety of possible reactions that can occur depending on the alloy substrate, the temperature range, the activity of the reacting species and the thermodynamics and kinetics of the various reactions.
TABLE
1
Metal-gas
reactions
Reacting
components
Substrate
or coating
Metal
Reaction Atmosphere Single
reacting
gas (02)
Alloy
Single gas
Metal
or alloy
products
reaction (02)
Mixture of reacting gases (02,
so,,
CO2)
Single-phase oxide (NiO, Cr,Os) 02 in solution and multiphase oxides (TiO, Ti203, TiOa) Continuous oxides of different elements Mixed oxides (spinels) Internal oxides Simultaneous or sequential formation of oxides, sulphides, carbides, etc.
Another class of high temperature alloys is based on the refractory metals molybdenum, tungsten, niobium and tantalum. These alloys are also Molybdenum and tungsten oxidize very reactive at high temperatures. Niobium and tantalum have a high rapidly by forming volatile oxides. solubility for oxygen and they also do not form a protective oxide scale. Thus the choice of high temperature alloys with inherent oxidation resistance is limited. The selection of these alloys for a given application is determined by the mechanical properties which are often not compatible with good oxidation resistance. This problem has been successfully overcome by a composite approach by forming protective coatings on the surface.
141
2. High temperature
coatings
High temperature coatings are designed to increase the life of the underlying alloy during service. Apart from this, the coatings also help in eliminating or reducing the critical and scarce raw materials used in the preparation of the alloys. On formation of a coating, the original surface problem is eliminated, but there is the new coating surface, and in addition, a coating-substrate interface which can create additional problems. The coating composition is selected to form the protective scale depending upon the environment it has to combat. Table 2 is a summary of the various corrosion processes encountered in practice along with the type of oxide that can offer resistance to a particular environment. The coatings should also be well bonded to the substrate which is decided by the coating process and the post-coating treatment. The problem of adhesion is accentuated by strains caused by differential heating or thermal cycling. A review of the fundamentals and applications of various high temperature oxidationresistant coatings has been prepared by the U.S. National Materials Advisory Board, which is a good source of general references [ 41. There are a number of methods of forming the coatings. They can be classified as processes that involve transfer of atoms and those formed by the transfer of particles, as shown in Fig. 1. The former covers chemical vapour deposition and physical vapour deposition processes such as evaporation, sputtering and ion plating. Slurry spraying, flame and plasma spraying deal with the deposition of micron-sized powders at room temperature or in molten form. All these methods are at present in industrial use and the choice depends on the desired coating compound, the engineering aspects concerning the type of parts to be coated and, ultimately, the economics. TABLE
2
High temperature
corrosion
Type
Reactant
Oxidation
02
Carburization
co,
Preferred
oxide
Al203 on Ni, Fe or Co alloys SiOl on refractory alloy
co2
so2
Hot corrosion
Na2S04, v205
cr203,
SiO2
muffle
Heating elements Aircraft and space vehicle components in
Coal gassifiers
Cr203
NaCl
Furnace
Components petroleum industry
A1203
CH4
Sulphidation
Applications
A1203
Gas turbine blades and vanes Diesel engine components
142
Fig. 1, Classification
3. Protection
of coatings.
for turbine blades
Among the various high temperature combustion systems, gas turbines, especially aircraft turbines, use the maximum number of coated components. This is primarily due to the demand on the components for high performance and reliability. In addition, the pay-off is high, even for a marginal improvement in efficiency or specific fuel consumption. In the gas turbine, the turbine blades and vanes occupy a place of prime importance. These are small intricate components with the highest specific cost. The turbine blades and vanes are made of nickel-based alloys. The high temperature strength is produced by Ni,Al precipitates coherent with the nickel-solid-solution matrix. Figure 2 shows the various factors that limit the life of the turbine blades [5]. Between 800 and 900 “C, the blades undergo an accelerated form of attack known as hot corrosion. This is primarily caused by molten salts such as Na,SO, and NaCl. In spite of extensive investigation over the last 20 years, our theoretical understanding of this complex process is not very satisfactory [6]. Hot corrosion is a severe problem for gas turbines operating in the marine atmosphere. The nickel-based alloys that are specified for turbine blades contain not more than 10 wt.% Cr in order to have enough volume fraction of Ni,Al precipitates for strength. This chromium content is not sufficient to form a protective CrzO, scale. Beyond 900 “C, the oxidation process is rate controlling and the bare alloys undergo attack at an unacceptable rate in the absence of protective Cr,O, or Al,O, scale. These alloys also have refractory elements up to 15 wt.%, which impair the oxidation resistance. Thus the need for coatings is beyond doubt, especially to meet the life requirement of 5000 - 10 000 h. It is said that modern aircraft would not make a single transatlantic flight, but for the coatings. The demand for improved thrust and efficiency of gas turbines has led to the development of a number of alloys and special processing methods such as directional solidification and the production of single-crystal blades.
143
Metal temperature,
Fig. 2. Factors fatigue; - .. :,
‘C
influencing turbine component life [5]: thermal fatigue; - - -, oxidation; -,
creep; sulphidation.
-,
-
-,
mechanical
There has also been a corresponding demand on the performance of the coatings, which has led to changes in the type of coatings and their mode of formation. Taking the turbine blade as a typical example, the philosophy behind the coating development for its protection, with emphasis on the choice and the characteristics of the coating process, is outlined below.
4. Nickel aluminide diffusion
coatings
The earliest coatings used from the late 1950s are the diffusion aluminide coatings based on the intermetallic compound NiAl. These coatings are still in active use, especially in the older engines. The aluminium content of these coatings is sufficient to form AlzO, scale at high temperature. There are a number of methods of forming aluminide coatings such as slurry spraying followed by high temperature diffusion, electrolysis and pack cementation [7]. Of all these methods, pack cementation is the one widely used, as it is inexpensive and ideally suited for coating small components. Pack cementation is a modified vapour deposition process, in which the blades are embedded in a powder mixture containing aluminium, inert A120, filler to prevent the sintering of the pack and 1 - 2 wt.% of ammonium halide activators. The whole assembly is then heated in the temperature range 800 - 1100 “C in an H2 or argon atmosphere. In this temperature range, aluminium halides form which diffuse through the porous pack and react at the surface of the component to deposit aluminium either by decomposition of aluminium halides or by H2 reduction. The deposited aluminium diffuses into the substrate to form the NiAl coating. Figure 3 shows schematically the different coating steps. Levine and Caves
144
--__
\
Substrate
I
Gas Phase niffuslon Of Al HalIdes
Al
-_-IL
Concentrotlan ProfIle
--O1stanie Srhemot~c
D,ogrom
of Coating
Fig. 3. Schematic
Steps
diagram
During
Pack
Alumnwng
of coating
steps during
pack
aluminizing.
[8] have worked out a method of calculating the rate of supply of aluminium in the gas phase. Detailed thermodynamic calculations of the pack reactions have been made for different activators and a new theory of coatings formation which can predict the surface composition and coating thickness for the case of unalloyed nickel has been worked out by Sivakumar and Seigle [9]. The above understanding of the process can serve as a guide in selecting process parameters for coating of complex alloys. The aluminide coatings can be broadly divided into the low activity and the high activity coatings [lo]. In the low activity pack process, an alloy pack is used so that NiAl is formed directly on the blade at a temperature of 1050 - 1100 “C. Usually in this process, a nickel-rich NiAl compound is formed. In the high activity process, which is more often used, the N&Al, phase is initially formed using a pure aluminium pack. This is later diffusion heat treated at 1050 - 1100 ‘C, in the absence of the aluminium source, to form NiAl by reacting with the substrate. Figure 4 shows the typical microstructure of these two types of coatings [ll]. In addition, the figure shows the microstructures of chromoaluminized coatings, in which the NiAl with an increased chromium content is obtained to improve the hot corrosion properties of NiAl coatings. It is accomplished by a two-step process in which the nickel-based alloys are initially chromized and then aluminized. The structural features of these coatings can be explained in terms of the basic diffusion mechanism in the NiAl compound. The differences arise because nickel is the predominant diffusing species in nickel-rich NiAl and aluminium in the aluminium-rich NiAl and the Ni,Al, phase [12]. The structure and composition of the coating have been found to be dependent on the substrate alloy, making it necessary to optimize the process parameters for a given substrate. This implies that the coatings are generally tailor made for a given alloy. This is well illustrated in the case of aluminide coatings on nickel-based alloys containing no aluminium, such as Hastealloy X. In these cases, because of the outward diffusion of nickel from the alloy, either discrete or continuous voids form at the coating-substrate interface leading to thermally induced spalling of the coating [ 131.
145
(a)
Fig. 4. Microstructures of NiAl coatings [ 111: (a) high activity pack; (b) low pack. Left-hand micrographs, aluminide; right-hand micrographs, chromoaluminide.
activity
The coating is not in thermodynamic equilibrium with the alloy, and hence on long exposure aluminium from the coating diffuses into the substrate. The coating life is affected by this diffusional degradation. The other life-limiting factor is the formation of Al,O, scale which spalls during thermal cycling. This is followed by a transient period during which all possible oxides, namely NiO, Cr,03 and A1203, form on the surface. A layer of Cr,O, then forms under these oxides, below which a continuous layer of AlTO develops. This process is repeated until the outward flux of aluminium is insufficient to form a continuous scale. It is to be noted that these two degradation mechanisms can operate simultaneously on different parts of the same blade. Similarly, in the presence of Na,S04, the protective A1203 scale is thought to react by a fluxing reaction. The accelerated corrosion may also come about by the mechanical damage of the oxide scale caused by thermal cycling or cracking in the presence of NaCl vapour [ 61. An exact description of the mechanism is not available. It has been found that the NiAl compound by itself is not resistant to hot corrosion. In a coated alloy, however, chromium from the substrate diffuses to the coating and provides the necessary protection. Figure 5 shows a typical degradation undergone by an aluminium-rich NiAl coating in an accelerated laboratory coating crucible test at 950 “C in 85wt.%Na$O,-15wt.%NaCl molten salt mixture [ll]. After 20 h, only the outer part of the coating has been converted to NisAl,
146
Fig. 5. Scanning
electron
micrographs
of hot-corroded
NiAl [ll]
: (A) 20 h, (B) 40 h.
whereas after 40 h the transformation is complete throughout the coating. The outward diffusion of aluminium and the density difference between the two phases have resulted in the formation of pores in these coatings. Apart from the laboratory tests, the cyclic oxidation and hot corrosion properties are evaluated in a combustion rig which simulates the gas turbine atmosphere. This dynamic rig testing, over and above the laboratory tests, is essential as the rate of oxidation and the extent of spalling have been found to depend on the amplitude and frequency of fluctuations in the heat flow from the gas stream to the blades [14]. Unfortunately, the rig testing procedures have not been standardized. Each engine manufacturer has his own rig, whose characteristics are unknown and uses test conditions which have evolved empirically over a period of time. The diffusion aluminide coatings, being thin (75 - 100 pm) do not significantly affect the mechanical properties of the underlying alloy. This has been shown to be true as regards the stress-rupture properties of a number of coated alloys [15]. One of the major limitations of these coatings, however, is their brittleness below about 600 ‘C, which cannot be easily overcome in this intermetallic compound. The cyclic strains induced by the temperature gradients in the aerofoil can lead to thermal fatigue cracks in the coating. These cracks can propagate into the alloy and reduce the fatigue life. When the metal temperature exceeds 950 OC, there is a considerable reduction in the life of the aluminide coatings. Incorporation of various alloying elements such as chromium and silicon and of inert oxides such as Al,O, into the basic NiAl compound were tried. The improvement was only marginal. One of the successful breakthroughs came with the advent of noble metal plating (platinum, rhenium etc.) followed by aluminizing [16]. The aluminium reacts with platinum to form platinum aluminide, in addition to the formation of NiAl. This process improves the life of the coatings by a factor of 3. Only limited investigations have been carried out on the reasons behind such improvement.
147
5. Overlay coatings Before the platinum aluminide process gained wide commercial acceptance, the concept of forming overlay coatings based on MCrAlY (M = Fe, Co or Ni) compounds was advocated. In these coatings, aluminium is required to form protective Al,O, scale. Chromium promotes the formation of Al,O, scale, as well as improving the hot-corrosion resistance. The concept that small additions of rare earth elements (maximum of 1 wt.%) improve the adherence of oxide scales emanated from oxidation studies of dispersionstrengthened materials. This idea was adopted and the overlay coating compositions modified by the addition of yttrium. The reasons for such a beneficial effect are not well established [ 21. These coatings have at least four elements and their deposition by the pack method requires each element to be deposited separately. It is difficult to alter the partial pressures of the halides of the various elements which are dictated by the thermodynamics of the pack to deposit them in a single cycle. Thus the electron beam evaporation process was further developed to form these coatings 1171. In this method, an alloy ingot is melted by an electron beam and the evaporated elements deposit onto the blades. The line-of-sight nature of this process, unlike the pack process, demands complex movements for the blades to obtain uniform coatings. As the vapour pressures of the elements differ, suitable compositional adjustment must be incorporated in the starting ingot. The deposition rates are high (10 pm min-‘) and this is still an expensive process because of the high investment cost. Also, the control of coating composition is difficult when it contains elements with high vapour pressures. Consequently, the electron beam process is not appropriate for trying various compositions. Owing to its simplicity in operation and economy, the low pressure plasma spraying process has started to replace the electron beam evaporation method. There have been attempts to form the overlay coatings by sputtering [ 181, ion plating and a combination of slurry spraying and pack aluminizing which is called the controlled-composition reaction-sintered coating process [ 191. These methods have not taken off as viable production routes even though there are no major technical difficulties. A recent account of these and other alternative processes has been presented by Restall and Wood [20]. The plasma spraying process is normally carried out in air at atmospheric pressure. In this process, the alloy powder is melted in a high power d.c. plasma flame formed in a flowing Ar-Hz mixture [21]. The plasma torch acts not only as a heat source, but also accelerates the powders to velocities of up to 300 m s-l. In the short residence time of a few milliseconds, the powder particles are melted and they strike and flatten on the object to be coated. By repeated movement of the gun, the coating is built up of layer on layer of these flattened droplets. There are a number of process variables that need to be controlled to obtain coatings of high density. The development of this process has so far proceeded in an
148
empirical manner. Recently attempts have been made to calculate the timetemperature history of particles in the plasma from a fundamental consideration of heat and momentum transfer [22]. The very nature of this process yields coatings with a porosity in the range 2% - 10%. Spraying in air also causes oxidation of the powders because of turbulent mixing of the surrounding air with the plasma gas. Initially, attempts were made to eliminate these problems by spraying in an argon chamber. This was only partially successful and the next logical step was the development of spraying in a low pressure chamber (50 - 100 Torr). The low pressure chamber offers a number of advantages especially for spraying metallic powders [23]. The plasma becomes elongated in the low pressure chamber compared with that at atmospheric pressure. This implies a gradual drop in the plasma flame temperature. The main advantage of this process lies in the fact that the turbine blade can now be heated to high temperatures (800 - 900 “C) without the problem of oxidation. As the plasma gas from the gun expands into a low pressure chamber, the plasma gas velocity also increases compared with that for spraying in air. The molten droplets hit the hot blade with a high velocity and lose the particle shape, thus forming a recrystallized coating. The dramatic difference in the coating structure can be seen in Fig. 6 [24]. In this process a numerically controlled five-axis mainipulator system is required for moving the blade to obtain uniform coatings on the entire surface. Alternately, a robot has been used. A wide variety of compositions have been tried in the MCrAlY series based on iron, nickel, cobalt and their combinations. Some of the coatings are single phase, but most of them are two phase, containing NiAl or CoAl precipitates in an Ni-Cr or Co-Cr matrix. All these coatings rely on the
(a) Fig. 6. Structure of low pressure plasma
plasma
sprayed;
(b)
149
formation of a protective AlzO, scale. The powders for spraying are made by the argon atomization process and hence it is easy to obtain powders of different compositions. These powders can be sprayed without any difficulty using the low pressure plasma spraying process. Such flexibility helps to optimize the coating composition for the best protection. The coating ductility is also affected by the composition. Figure 7 shows how, by decreasing the chromium and aluminium contents of a CoCrAlY coating, the ductile-brittle transition temperature can be decreased from 540 “C (1000 “F) to 95 “C (200 “F). This is due to a smaller volume fraction of the brittle CoAl phase dispersed in a more ductile cobalt matrix for the Co-15.5wt.%Cr-10.2wt.%Al-0.3wt.%Y coating. The superiority of the overlay coatings to the nickel aluminide is also shown in the Fig. 10. Even though the overlay coatings by their very nature are independent of the substrate, a small diffusion zone of about 10% of the coating thickness is formed by post-coating heat treatment to improve the bonding. During service, the substrate elements and the coating elements interdiffuse making it necessary to evaluate the long-term diffusional stability of these coatings. The oxidation and hot-corrosion properties of these coatings have also been found to be dependent on the substrate composition. Pettit and Goward [ 251 have reported that the NiCoCrAlY coating (18wt.%Cr, 23wt.%Co, 12wt.%Al, 0.5wt.%Y, remainder Ni) shows better oxidation resistance on a hafnium-containing alloy. This has been attributed to greater adherence of the Al,O, scale which has been pegged by HfO,. Apparently the hafnium from the alloy has diffused through the coating to form these pegs. Attempts are being made to improve the oxidation and hot-corrosion resistance of the overlay coatings further by adding hafnium, tantalum and silicon to the base NiCoCrAlY composition. However, caution is required in translating the properties of a particular coating on a specific alloy to other types of alloys. The properties are also known to differ depending upon the specific environment in which the gas turbines operate.
0
0
200
400
Fig. 7. The ductility
600
BOO
1000
Tempetalure “F
of NiAl and CoCrAly
1 ‘loo
1600
coatings
[lo].
1200
150
6. Thermal
barrier coatings
The next generation of advanced coatings were based on ceramic materials which acted as a thermal barrier between the hot combustion gas and the turbine metal surface [26]. This approach has already been successfully used in combustor and after-burner applications. The reliability and durability of these coatings need to be improved for successful applications on turbine blades. The incentive to develop these coatings is high as they can decrease the metal temperature by up to 200 “C and thus can improve the creep and oxidation life. Alternatively, the cooling air consumption can be reduced by 6% for the same metal temperature. This will have a significant impact on the efficiency of the engine. Plasma spraying is the most attractive method of forming these ceramic coatings which are 200 - 300 pm thick. They are applied on a metallic bond coat for better adherence. The main limitation of the ceramic coatings is their low strain tolerance. The mismatch in thermal expansion between the coating and the underlying substrate is a major factor leading to the spalling of the coating. These coatings are porous and thus the oxidation of the underlying metallic bond coat also leads to spalling [ 271. Y,O,-stabilized ZrO, compounds have shown considerable promise as coating materials [ 281. A number of approaches have been tried to improve the thermal shock resistance of zirconia coatings. These include segmented structure, microcracked structure, graded coatings and composition control to obtain a cubic plus tetragonal two-phase structure [25]. In the presence of molten salt, it is reported that Y,O, reacts leading to destabilization of the cubic phase. Even when there is no reaction, as in the case of MgO. ZrO,, the molten salt can penetrate the coating and oxidize the underlying bond coat. Figure 8 shows 0
Fig. 8. Hot-corroded micrograph with salt;
Yz03-ZrOa (b) electron
coating and X-ray
on CoCrAlY images.
bond
coat
[27]:
(A)
optical
151
such an attack which leads to spalling of the MgO* ZrO, coating [27]. Already coated with nozzle vane platforms thermal barriers are in use in some aero engines [29]. Ceramic-coated vanes have been successfully engine tested and it may not be long before ceramic-coated blades propel aircraft. Another process that is competing to form these coatings is electron beam evaporation [30]. This process is attractive as it forms a strain-tolerant columnar structure of ZrO, which is denser than the plasma-sprayed coatings. The increased density lowers the insulative capacity, but increases the erosion resistance of these coatings. Recently the use of ion or plasma assistance in the electron beam physical vapour deposition of stabilized zirconia coatings has been shown to be highly beneficial in structure control and adhesion enhancement [ 31, 321.
7. Future
developments
The properties of the coatings depend critically upon the process, as the coating structure varies with the method of formation. The coating processes have their own inherent limitations. One of the viable methods to overcome some of these limitations is by post-coating treatments. The well known among these are high temperature diffusion treatment to improve bonding and shot-peening to induce compressive stresses and to obtain a smooth surface [33]. Of late, the laser has been recognized as a versatile tool to remelt the coatings. Laser remelting can help to remove the pores that are present in plasma-sprayed coatings or help to modify the structure of a fully dense coating. In view of its great potential, a summary of its use in the context of turbine blade coatings is presented below. 7.1. Metallic coatings One of the major limitations of diffusion aluminide coatings is its limited ductility. Hence, during temperature cycling, the thermal stress can induce cracking of the coated layer. These cracks have been found to extend into the substrate and to damage the component [29]. The mechanical effects of the coatings are complex. An overview of these problems in general has been presented by Evans et al. [ 341 and in particular reference to turbine blades by Schneider and Griinling [35]. The role of residual stresses and the effect of various mechanisms of stress relaxation have been outlined. The role of interface defects in initating buckling and eventual spalling of the coating has also been analysed [35]. The effect of coating composition and the type of coating on ductility are shown in Fig. 7. This in turn affects the thermomechanical fatigue life as shown in Fig. 9 1351. As shown in Fig. 7, in the case of MCrAlY coatings, the ductility can be altered by changing the aluminium content. Such a dramatic change is not possible with NiAl coatings. Laser remelting of NiAl coatings offers a distinct possibility of increasing the ductility. Such an increase has been
152
Fig. 9. Thermomechanical
fatigue
properties
of NiAl and CoCrAlY
coatings
[33].
demonstrated in NiAl-based alloys through refining the grain size by thermomechanical processing [ 361. For grain sizes below 20 pm, elongations greater than 10% were obtained at room temperature. There is some controversy whether such an increase can be obtained by rapid solidification, which can reduce the ordering in NiAl. Increased ductility by reduced ordering has been well established for Ni,Al [3’7]. In view of the above findings, it would be worth while to laser melt the NiAl coatings to increase their ductility. As-formed NiAl coatings are unstable as there is a gradient in the concentration of aluminium from the surface (highest) to the interior. Laser melting would remove this concentration gradient and thereby increase the microstructural stability. However, care must be taken to ensure an adequate amount of aluminium in the NiAl coatings after laser melting in order not to impair its resistance to hot corrosion [ 291. Laser melting is beneficial for air-plasma-sprayed MCrAlY types of overlay coatings. It offers an alternative route to form a dense and oxidefree coating which can replace the more expensive methods such as electron beam evaporation and low pressure plasma spraying. This approach was investigated by Dalliare and Cielo [ 38, 391 and Capp and Rigsbee [40] by melting Hastelloy, nickel and Stellite coatings by laser. Bhat et al. [41] have successfully sealed pores by melting a thin top layer (20 - 25 E.trn thick) of plasma-sprayed NiCrAlY coatings using a CO, laser operating in the continuous wave mode [ 411. Detailed investigations into the potential and limitations of laser melting of MCrAlY coatings have been carried out by Sivakumar and Mordike [42]. It was found that there is no problem of absorption of laser radiation and a desired thickness can be melted. Figure 10 shows the typical cross-sectional morphologies of the coatings that were melted partially and fully. However, there is a depletion of aluminium and yttrium in the laser-melted coatings owing to the formation of oxides of
153
(a) Fig. 10. Cross-sectional melted; (B) fully melted.
(b) micrographs
of
laser-melted
NiCoCrAlY
coatings:
(A) partially
the above elements. These oxides can be removed as they are driven to the surface by convection currents during laser melting. Figure 11 shows the typical surface appearance of laser-melted NiCoCrCrAlY coatings with the laser operating in the continuous wave and the pulsed modes. The dark regions are oxides of aluminium and yttrium. This loss is reflected by poor cyclic oxidation resistance. Hence the initial plasmasprayed coating compositions need to be adjusted to take account of the above losses and thereby to obtain optimum properties. It is also possible to carry out the same process with an electron beam [43]. The laser is more versatile as it requires no vacuum for its operation. Another approach is to avoid the plasma spraying step and to melt directly MCrAlY powders, attached to the substrate by a binder, using a laser [44]. The control over the melt thickness, uniformity and the extent of dilution with substrate are difficult and more work needs to be carried out to make this process viable. Another possibility is laser spraying or cladding, in which the powders are fed directly into the laser-melted bath by a delivery system [45 - 471. When the powders are delivered coaxially into the laser beam, they are accelerated to a few metres per second. Possible mechanisms of transfer of momentum are discussed by Schaefer et al. [45]. These velocities by themselves are not sufficient to form dense coatings, as these are very low compared with the velocities of the droplets reached during plasma spraying (a few hundred metres per second). Hence the laser must melt the particles that are falling onto the substrate along with a part of the substrate to form a fully dense coating. It is also possible to spray the
(a) Fig. Il. Surface morphology mode; (B) pulsed mode.
(b) of laser-melted
NiCoCrAlY
coatings:
(A) continuous
wave
powder at an angle close to the substrate and to use the laser to melt the powder and alloy at the same time as the substrate [ 471. Because of process control problems, these technologies are at various stages of development. At present, they can be used for less critical applications. Laser melting of fully dense MCrAlY coatings formed by electron beam evaporation or low pressure plasma spraying has also been found to be beneficial in increasing the cyclic oxidation resistance [ 48, 491. Bornstein and Smeggil [48] found a thin tenaceous oxide scale followed by a singlephase zone in the case of laser-melted NiCrAlY coatings after testing in a burner rig at 985 “C for 100 h. The corresponding non-laser-melted specimen showed some spalling of the oxide, and the matrix below the oxide consisted of a two-phase structure of NiAl and Ni-Cr solid solution. It was also observed that there was an accumulation of yttrium beneath the scale, whereas no such yttrium was detected in the laser-melted specimens. In their later study, it was reported that the as-formed coating contained discrete yttrium-enriched particles and yttrium could not be detected on the laser-melted surface [49]. Also the laser-melted surface had a fine micron-sized distribution of NiAl phase as opposed to a coarse distribution of the same in the as-formed coating. Laser melting made the coated surface more homogeneous and thus promoted the formation a uniform, thin and adherent oxide. Further investigations are needed to understand the reason behind such an improvement. Also interesting would be the hot corrosion properties of such laser-melted layers.
155
7.2. Ceramic coatings In the development of thermal barrier ceramic coatings, laser melting is likely to play a key role. The laser is uniquely suited to remelt a part of the ceramic coatings with a minimal input of heat. Such a remelted layer would be fully dense and so would prevent the corrosive gases and salts from attacking the coating and the underlying substrate. A few investigations towards achieving the same have been reported [ 50 - 541. The major problem which is yet to be solved is the cracking of the laser-melted layer owing to non-uniform cooling [55]. Figure 12 shows the typical surface morphology of laser-melted surfaces of the most popular 7Y,O,*ZrO, coatings and the corresponding cross-section showing the transverse cracks
(b) Fig. 12. Laser-melted
7Y203*ZrOz
coatings:
(a) surface
morphology;
(b) cross-section
156
across the melted layer. Detailed investigations on the various aspects of laser melting of different ceramic coatings, namely Zr02, Al,O, and Ti02, have recently been carried out [56]. In this study it was shown how the extent of cracking can be reduced by proper choice of laser operation
(b)
(a)
(cl Fig. 13. Surface temperature;(c)
morphology laser melted,
of A1203 800 “C.
coatings:
(a) as sprayed;
(b) laser melted,
room
157
parameters. Also by a simple analysis of stresses that arise during cooling it was shown that it is possible to melt a thin surface layer of Al,O, and TiO,? without cracking by preheating the coating to 800 “C!. Figure 13 shows the typical surface morphology of an Al,O, surface that was laser melted at room temperature and the surface without cracks which was melted with preheating to 800 “C. Figure 14 shows the benefits in terms of avoidance of spalling of the coatings in a laboratory hot-corrosion test with 95%Na,SO,5%NaCl salt at 900 “C for 100 h [56]. The above solution is not practical for ZrO, systems, as the coatings have to be heated to temperatures above 1000 “C to avoid cracking. It may be possible, however, to reduce this temperature by altering the Y,O,stabilized ZrOz compositions to increase the fracture toughness. One possibility is to add up to 20wt.% Al,Os to Y,O,-stabilized ZrO, to increase the toughness from 6 to 7.6 MN m-” [57]. The other option is to choose a low Y203 content ZrO, composition, in which the tetragonal-to-monoclinic phase transformation takes place during cooling. The expansion associated with the above phase change may counteract the thermal contraction. It is also feasible to try Zr02* SiO* coatings, in which the presence of SiO, helps to relieve the stresses and thereby prevents cracking [ 541.
10 mm Fig. 14. Hot-corroded A1203: laser melted at 800 “C.
0, as sprayed;
1, laser
melted
at room
temperature;
2, 3,
8. Conclusion It has been shown how the demand for better performance has led to the development of new coatings and coating methods. Table 3 gives a summary of the coating developments for turbine blades. This development encompasses an interaction between the physical metallurgy of the coating and the processing. A fundamental understanding of the mechanism of formation of different coatings and their mode of degradation is essential for the design of suitable coatings. In the case of high temperature coatings, the properties are considerably affected by the structure and chemistry of the substrate. It is hence prudent to design a coated alloy as a single entity
158 TABLE Coatings
3 for turbine
blades
Coating type
Coating method
Coating phase
Mode of protection
Diffusion
Pack aluminizing
NiAl
A1203
Overlay
Electron beam evaporation
MCrAlY (M = Ni, Co, Fe)
More spallresistant
Overlay
Low pressure plasma spraying
MCrAlY (M = Ni, Co, Fe)
More spallresistant
Thermal
Air plasma spraying Electron beam processes
Partially stabilized ZrOz
Insulative
barrier
Limitations
Hot corrosion Brittleness Diffusional stability Hot corrosion Thermal fatigue
Al203
Hot corrosion Thermal fatigue
A1203
Thermal spalling, oxidation and hot corrosion of bond coat
as opposed to the current practice of trying to marry the existing ones to the various alloys, with minor modifications here and there. The future trend will be to go for more post-coating modifications to suit specific applications, rather than major changes in coating processes.
References 1 A. S. Radcliff, Mater Sci. Technol., 3 (1987) 554. 2 G. C. Wood and F. H. Stott, Mater. Sci. Technol., 3 (1987) 519. 3 N. Birks and G. H. Meier, Introduction to High Temperature Edward
Arnold,
London,
1983,
Oxidation
of Metal,
p. 91.
4 High Temperature Oxidation Resistant Coatings, National Academy of Sciences, Washington, DC, 1970. 5 D. A. Spera and S. J. Grisaffe, NASA Tech. Memo. TMX-2664, (available from National Technical Information Service, Springfield, VA). 6 9. Stringer, Mater Sci. Technol., 3 (1987) 482. 7 S. J. Grisaffe, in C. T. Sims and W. C. Hagel (eds.), The Superalloys, Wiley, New York, 1972, p. 341. 8 S. R. Levine and R. M. Caves, J. EZectrochem. Sot., 121 (1974) 1051. 9 R. Sivakumar and L. L. Seigle, Metall. Trans. A, 7 (1976) 1073. 10 G. W. Goward, in Z. A. Foroulis and F. S. Pettit (eds.), Properties of High Temperature Alloys, Electrochemical Society, Princeton, NJ, 1976, p. 806. 11 R. Sivakumar, Oxid. Met., 17 (l-2) (1982) 27. 12 R. Sivakumar and L. L. Seigle, Trans. Indian Inst. Met., 31 (5) (1978) 336. 13 J. Chin, W. R. Johnson and K. Chen, Thin Solid Films, 95 (1982) 85. 14 R. R. Dils, in Fatigue at Eleuated Temperatures, ASTM Spec. Tech. Publ., 520 (1973) 102.
159 15 16
17
18
19
20
21 22
23 24 25 26
27 28 29
30 31
32 33 34 35 36 37 38 39
40 41 42
43 44
45
46
47 48
P. Lane, Jr., and N. M. Geyer, J. Met., 18 (2) (1966) 186. G. Lehnert and H. W. Meinhardt, Electrodep. Surf. Treat., I (1972) 189. G. W. Goward, Proc. 1974 Gas Turbine Materials in the Marine Environment Conf., MCZC 75-27, (1975) 277 (available from National Technical Information Service, Springfield, VA). J. Fairbanks, R. Busch, J. Patten and E. McClanahan, Proc. 1974 Gas Turbine Materials in the Marine Environment Conf., MCIC 75-27 (1975) 429 (available from National Technical Information Service, Springfield, VA). A. R. Stetson and V. S. Moore, Proc. 1974 Gas Turbine Materials in the Environment Information Conf., MCZC 75-27 (1975) 457 ( available from National Technical Service, Springfield, VA). J. E. Restall and M. I. Wood, Mater. Sci. Technol., 2 (1986) 225. D. A. Gerdeman and N. L. Hecht, Arc Plasma Technology in Materials Science, Springer, New York, 1972. J. H. Zaat, Annu. Rev. Mater. Sci., 13 (1983) 9. H. D. Steffens, in E. Lang (ed.), Coatings for High Temperature Applications, Applied Science, London, 1983, p. 121. J. P. Merutka, NASA Tech. Memo. TM-82740, 1981 (available from National Technical Information Servie, Springfield, VA). F. S. Pettit and G. W. Goward, in J. K. Tien and J. F. Elliott (eds.), Metallurgical Treatises The Metallurgical Society of AIME, Warrendale, PA, 1981, p. 603. S. R. Levine, R. A. Miller and P. E. Hodge, SAMPE Q., 12 (1) (1980) 20. R. Sivakumar and M. P. Srivastava, Oxid. Met., 20 (3-4) (1983) 67. G. Johner and K. K. Schweitzer, Thin Solid Films, 119 (1984) 301. G. W. Meetham, Mater. Sci. Technol., 2 (1986) 290. T. E. Strangman, Thin Solid Films, 127 (1985) 93. K. S. Fancey and A. Matthews, J. Vat. Sci. Technol. A, 4 (1986) 2652. A. S. James, K. S. Fancey and A. Matthews, Surf. Coat. Technol., 32 (1987) 377. D. F. Bettridge and R. G. Ubank, Mater. Sci. Technol., 2 (1986) 232. A. G. Evans, G. B. Crumley and R. E. Demaray, Oxid. Met., 20 (5-6) (1983) 193. K. Schneider and H. W. Griinling, Thin Solid Films, 107 (1983) 395. E. M. Schulson, in B. Giessen, D. E. Polk and A. I. Taub (eds.), High temperature ordered intermetallic alloys, Mater. Res. Sot. Symp. Proc., 39 (1985) 193. A. J. Taub and M. R. Jackson, Mater. Res. Sot. Symp. Proc., 58 (1986) 389. S. Dallaire and P. Cielo, Thin Solid Films, 108 (1983) 19. S. Dallaire and P. Cielo, Metall. Trans. B, 13 (1982) 479. M. L. Capp and J. M. Rigsbee, Mater. Sci. Eng., 62 (1984) 49. H. Bhat, H. Herman and R. J. Royle, in S. C. Singhal (ed.), High Temperature Protective Coatings, The Metallurgical Society of AIME, Warrendale, PA, 1983, p. 37. R. Sivakumar and B. L. Mordike, Laser meiting of plasma sprayed NiCoCrAlY coatings, Surf. Eng., 3 (4) (1987) 299. C. Burman and T. Ericsson, in S. C. Singhal (ed.), High Temperature Protective Coatings, The Metallurgical Society of AIME, Warrendale, PA, 1983, p. 51. C. A. Liu, M. J. Humphries and D. W. Mason, Thin Solid Films, 107 (1983) 251. C. A. Liu, M. J. Humphries and R. C. Krutenat, Thin Solid Films, 107 (1983) 269. R. J. Schaefer, T. R. Tucker and J. D. Ayres, in E. A. Metzbower (ed.), Applications of Lasers in Materials Processing, American Society for Metals, Metals Park, OH, 1979, p. 135. T. R. Tucker, A. H. Clauer, I. G. Wright and J. T. Stropki, Thin Solid Films, 118 (1984) 73. W. M. Steen, in D. Belforte and M. Levitt (eds.), The Industrial Laser-Annual Handbook, Penn Well Books, OK, 1986, p. 158. N. S. Bornstein and J. G. Smeggil, in C. R. Clayton and C. M. Preece (eds.), Corrosion of Metals Processed by Direct Energy Beams, Metallurgical Society of AIME, Warrendale, PA, 1982, p. 147.
160 49 50 51 52 53 54 55 56 57
J. G. Smeggil, A. W. Funkenbusch and N. S. Bornstein, Thin Solid Films, 119 (1984) 327. F. S. Galasso and R. Veltri, Ceram. Bull., 62 (2) (1983) 253. I. Zaplatiynsky, Thin Solid Films, 95 (3) (1982) 275. N. Iwamoto, N. Umesaki and S. Endo, in Aduances in Thermel Spraying, Proc. 11 th Int. Thermal Spraying Conf., Pergamon, Oxford, 1986, p. 563. A. Adamski and R. McPherson, in Advunces in Thermal Spraying, Proc. 11 th Int. Thermal Spraying Conf., Pergamon, Oxford, 1986, p. 555. M. Havrda, K. Volenik, J. Wagner and P. Mraz, in Advances in Thermal Spraying, Proc. 11 th Znt. Thermal Spraying Conf., Pergamon, Oxford, 1986, p. 569. R. Sivakumar and B. L. Mordike, in B. L. Mordike (ed.), Proc. Conf. on Laser Treutment ofMaterials, Deutsche Gesellschaft fur Metallkunde, Oberursel, 1987, p. 373. R. Sivakumar and B. L. Mordike, Laser Melting of Plasma Sprayed Ceramic Coatings, Surf. Eng., 4 (2) (1988) 127. Hou Yao Yong,
Gao Ji-Qiang
and Zhang
Hong-Tu,
J. Muter.
Sci. Left.,
6 (1987)
246.
and Coatings
Technology,
HIGH TEMPERATURE A REVIEW R. SIVAKUMAR*
37 (1989)
COATINGS
139
- 160
139
FOR GAS TURBINE
and B. L. MORDIKE
Znstitut fiir Werkstoffkunde und Werkstoffiechnik, Technische Agricolastrasse 2, 3392 Clausthal-Zellerfeld (F.R.G.) (Received
BLADES:
Uniuersitdt
Clausthal,
April 9, 1988)
Summary High temperature coatings are formed to protect many engineering components from environmental degradation. The operating temperature and the nature of the corrodant dictate the choice of the coating. The coatings can be formed by different methods, but invariably the properties are dependent on the coating process. The life of the coating depends upon the mode of degradation in service. It is also influenced by the nature of the substrate. The need for coatings is well illustrated by their wide use in gas turbines. This review covers the development of the oxidation- and hotcorrosion-resistant coatings for turbine blades. The rationale behind such protection, starting from the nickel aluminide to future thermal barrier oxide coatings, is outlined. It is followed by a discussion of current trends in post-coating modifications to improve the properties of the coatings. The principles underlying the development of these coatings should serve as a useful guide in the choice of coatings for other high temperature applications.
1. Introduction High temperature oxidation and other reactions in corrosive atmospheres are of considerable importance to aerospace, nuclear and similar energy-related industries. The widely used high temperature engineering alloys are based on the transition metals iron, nickel and cobalt, whose oxidation resistance can be improved by adding sufficient amounts of chromium, aluminium or silicon to form continuous external oxide scales. These oxides, namely Cr203, A120, and SiO,, offer the best protection owing to their low growth rates. Among these oxides, Al,O, is the most suitable as Cr,O, scale is not useful above 1000 “C, where volatile CrO, *On
leave
from
Defence
Metallurgical
Research
Laboratory,
Hyderabad
500258,
India. 0257-8972/89/$3.50
0 Elsevier Sequoia/Printed
in The Netherlands
140
forms. In fact, significant evaporation has been detected at even lower temperatures of 850 - 900 “C [l]. This has been attributed to the high gas velocities in the gas turbine. At low oxygen pressures, SiOz scale is not stable and decomposes to gaseous products such as SiO. Further details on the oxidation properties of these alloys can be found in the review article by Wood and Stott [a]. Wagner’s theory can predict the minimum concentration required for the formation of continuous external scale [3]. However, such predictions are not yet feasible for the oxidation of complex alloys or even for the case of a pure metal exposed to a mixture of reacting gases. Table 1 lists the variety of possible reactions that can occur depending on the alloy substrate, the temperature range, the activity of the reacting species and the thermodynamics and kinetics of the various reactions.
TABLE
1
Metal-gas
reactions
Reacting
components
Substrate
or coating
Metal
Reaction Atmosphere Single
reacting
gas (02)
Alloy
Single gas
Metal
or alloy
products
reaction (02)
Mixture of reacting gases (02,
so,,
CO2)
Single-phase oxide (NiO, Cr,Os) 02 in solution and multiphase oxides (TiO, Ti203, TiOa) Continuous oxides of different elements Mixed oxides (spinels) Internal oxides Simultaneous or sequential formation of oxides, sulphides, carbides, etc.
Another class of high temperature alloys is based on the refractory metals molybdenum, tungsten, niobium and tantalum. These alloys are also Molybdenum and tungsten oxidize very reactive at high temperatures. Niobium and tantalum have a high rapidly by forming volatile oxides. solubility for oxygen and they also do not form a protective oxide scale. Thus the choice of high temperature alloys with inherent oxidation resistance is limited. The selection of these alloys for a given application is determined by the mechanical properties which are often not compatible with good oxidation resistance. This problem has been successfully overcome by a composite approach by forming protective coatings on the surface.
141
2. High temperature
coatings
High temperature coatings are designed to increase the life of the underlying alloy during service. Apart from this, the coatings also help in eliminating or reducing the critical and scarce raw materials used in the preparation of the alloys. On formation of a coating, the original surface problem is eliminated, but there is the new coating surface, and in addition, a coating-substrate interface which can create additional problems. The coating composition is selected to form the protective scale depending upon the environment it has to combat. Table 2 is a summary of the various corrosion processes encountered in practice along with the type of oxide that can offer resistance to a particular environment. The coatings should also be well bonded to the substrate which is decided by the coating process and the post-coating treatment. The problem of adhesion is accentuated by strains caused by differential heating or thermal cycling. A review of the fundamentals and applications of various high temperature oxidationresistant coatings has been prepared by the U.S. National Materials Advisory Board, which is a good source of general references [ 41. There are a number of methods of forming the coatings. They can be classified as processes that involve transfer of atoms and those formed by the transfer of particles, as shown in Fig. 1. The former covers chemical vapour deposition and physical vapour deposition processes such as evaporation, sputtering and ion plating. Slurry spraying, flame and plasma spraying deal with the deposition of micron-sized powders at room temperature or in molten form. All these methods are at present in industrial use and the choice depends on the desired coating compound, the engineering aspects concerning the type of parts to be coated and, ultimately, the economics. TABLE
2
High temperature
corrosion
Type
Reactant
Oxidation
02
Carburization
co,
Preferred
oxide
Al203 on Ni, Fe or Co alloys SiOl on refractory alloy
co2
so2
Hot corrosion
Na2S04, v205
cr203,
SiO2
muffle
Heating elements Aircraft and space vehicle components in
Coal gassifiers
Cr203
NaCl
Furnace
Components petroleum industry
A1203
CH4
Sulphidation
Applications
A1203
Gas turbine blades and vanes Diesel engine components
142
Fig. 1, Classification
3. Protection
of coatings.
for turbine blades
Among the various high temperature combustion systems, gas turbines, especially aircraft turbines, use the maximum number of coated components. This is primarily due to the demand on the components for high performance and reliability. In addition, the pay-off is high, even for a marginal improvement in efficiency or specific fuel consumption. In the gas turbine, the turbine blades and vanes occupy a place of prime importance. These are small intricate components with the highest specific cost. The turbine blades and vanes are made of nickel-based alloys. The high temperature strength is produced by Ni,Al precipitates coherent with the nickel-solid-solution matrix. Figure 2 shows the various factors that limit the life of the turbine blades [5]. Between 800 and 900 “C, the blades undergo an accelerated form of attack known as hot corrosion. This is primarily caused by molten salts such as Na,SO, and NaCl. In spite of extensive investigation over the last 20 years, our theoretical understanding of this complex process is not very satisfactory [6]. Hot corrosion is a severe problem for gas turbines operating in the marine atmosphere. The nickel-based alloys that are specified for turbine blades contain not more than 10 wt.% Cr in order to have enough volume fraction of Ni,Al precipitates for strength. This chromium content is not sufficient to form a protective CrzO, scale. Beyond 900 “C, the oxidation process is rate controlling and the bare alloys undergo attack at an unacceptable rate in the absence of protective Cr,O, or Al,O, scale. These alloys also have refractory elements up to 15 wt.%, which impair the oxidation resistance. Thus the need for coatings is beyond doubt, especially to meet the life requirement of 5000 - 10 000 h. It is said that modern aircraft would not make a single transatlantic flight, but for the coatings. The demand for improved thrust and efficiency of gas turbines has led to the development of a number of alloys and special processing methods such as directional solidification and the production of single-crystal blades.
143
Metal temperature,
Fig. 2. Factors fatigue; - .. :,
‘C
influencing turbine component life [5]: thermal fatigue; - - -, oxidation; -,
creep; sulphidation.
-,
-
-,
mechanical
There has also been a corresponding demand on the performance of the coatings, which has led to changes in the type of coatings and their mode of formation. Taking the turbine blade as a typical example, the philosophy behind the coating development for its protection, with emphasis on the choice and the characteristics of the coating process, is outlined below.
4. Nickel aluminide diffusion
coatings
The earliest coatings used from the late 1950s are the diffusion aluminide coatings based on the intermetallic compound NiAl. These coatings are still in active use, especially in the older engines. The aluminium content of these coatings is sufficient to form AlzO, scale at high temperature. There are a number of methods of forming aluminide coatings such as slurry spraying followed by high temperature diffusion, electrolysis and pack cementation [7]. Of all these methods, pack cementation is the one widely used, as it is inexpensive and ideally suited for coating small components. Pack cementation is a modified vapour deposition process, in which the blades are embedded in a powder mixture containing aluminium, inert A120, filler to prevent the sintering of the pack and 1 - 2 wt.% of ammonium halide activators. The whole assembly is then heated in the temperature range 800 - 1100 “C in an H2 or argon atmosphere. In this temperature range, aluminium halides form which diffuse through the porous pack and react at the surface of the component to deposit aluminium either by decomposition of aluminium halides or by H2 reduction. The deposited aluminium diffuses into the substrate to form the NiAl coating. Figure 3 shows schematically the different coating steps. Levine and Caves
144
--__
\
Substrate
I
Gas Phase niffuslon Of Al HalIdes
Al
-_-IL
Concentrotlan ProfIle
--O1stanie Srhemot~c
D,ogrom
of Coating
Fig. 3. Schematic
Steps
diagram
During
Pack
Alumnwng
of coating
steps during
pack
aluminizing.
[8] have worked out a method of calculating the rate of supply of aluminium in the gas phase. Detailed thermodynamic calculations of the pack reactions have been made for different activators and a new theory of coatings formation which can predict the surface composition and coating thickness for the case of unalloyed nickel has been worked out by Sivakumar and Seigle [9]. The above understanding of the process can serve as a guide in selecting process parameters for coating of complex alloys. The aluminide coatings can be broadly divided into the low activity and the high activity coatings [lo]. In the low activity pack process, an alloy pack is used so that NiAl is formed directly on the blade at a temperature of 1050 - 1100 “C. Usually in this process, a nickel-rich NiAl compound is formed. In the high activity process, which is more often used, the N&Al, phase is initially formed using a pure aluminium pack. This is later diffusion heat treated at 1050 - 1100 ‘C, in the absence of the aluminium source, to form NiAl by reacting with the substrate. Figure 4 shows the typical microstructure of these two types of coatings [ll]. In addition, the figure shows the microstructures of chromoaluminized coatings, in which the NiAl with an increased chromium content is obtained to improve the hot corrosion properties of NiAl coatings. It is accomplished by a two-step process in which the nickel-based alloys are initially chromized and then aluminized. The structural features of these coatings can be explained in terms of the basic diffusion mechanism in the NiAl compound. The differences arise because nickel is the predominant diffusing species in nickel-rich NiAl and aluminium in the aluminium-rich NiAl and the Ni,Al, phase [12]. The structure and composition of the coating have been found to be dependent on the substrate alloy, making it necessary to optimize the process parameters for a given substrate. This implies that the coatings are generally tailor made for a given alloy. This is well illustrated in the case of aluminide coatings on nickel-based alloys containing no aluminium, such as Hastealloy X. In these cases, because of the outward diffusion of nickel from the alloy, either discrete or continuous voids form at the coating-substrate interface leading to thermally induced spalling of the coating [ 131.
145
(a)
Fig. 4. Microstructures of NiAl coatings [ 111: (a) high activity pack; (b) low pack. Left-hand micrographs, aluminide; right-hand micrographs, chromoaluminide.
activity
The coating is not in thermodynamic equilibrium with the alloy, and hence on long exposure aluminium from the coating diffuses into the substrate. The coating life is affected by this diffusional degradation. The other life-limiting factor is the formation of Al,O, scale which spalls during thermal cycling. This is followed by a transient period during which all possible oxides, namely NiO, Cr,03 and A1203, form on the surface. A layer of Cr,O, then forms under these oxides, below which a continuous layer of AlTO develops. This process is repeated until the outward flux of aluminium is insufficient to form a continuous scale. It is to be noted that these two degradation mechanisms can operate simultaneously on different parts of the same blade. Similarly, in the presence of Na,S04, the protective A1203 scale is thought to react by a fluxing reaction. The accelerated corrosion may also come about by the mechanical damage of the oxide scale caused by thermal cycling or cracking in the presence of NaCl vapour [ 61. An exact description of the mechanism is not available. It has been found that the NiAl compound by itself is not resistant to hot corrosion. In a coated alloy, however, chromium from the substrate diffuses to the coating and provides the necessary protection. Figure 5 shows a typical degradation undergone by an aluminium-rich NiAl coating in an accelerated laboratory coating crucible test at 950 “C in 85wt.%Na$O,-15wt.%NaCl molten salt mixture [ll]. After 20 h, only the outer part of the coating has been converted to NisAl,
146
Fig. 5. Scanning
electron
micrographs
of hot-corroded
NiAl [ll]
: (A) 20 h, (B) 40 h.
whereas after 40 h the transformation is complete throughout the coating. The outward diffusion of aluminium and the density difference between the two phases have resulted in the formation of pores in these coatings. Apart from the laboratory tests, the cyclic oxidation and hot corrosion properties are evaluated in a combustion rig which simulates the gas turbine atmosphere. This dynamic rig testing, over and above the laboratory tests, is essential as the rate of oxidation and the extent of spalling have been found to depend on the amplitude and frequency of fluctuations in the heat flow from the gas stream to the blades [14]. Unfortunately, the rig testing procedures have not been standardized. Each engine manufacturer has his own rig, whose characteristics are unknown and uses test conditions which have evolved empirically over a period of time. The diffusion aluminide coatings, being thin (75 - 100 pm) do not significantly affect the mechanical properties of the underlying alloy. This has been shown to be true as regards the stress-rupture properties of a number of coated alloys [15]. One of the major limitations of these coatings, however, is their brittleness below about 600 ‘C, which cannot be easily overcome in this intermetallic compound. The cyclic strains induced by the temperature gradients in the aerofoil can lead to thermal fatigue cracks in the coating. These cracks can propagate into the alloy and reduce the fatigue life. When the metal temperature exceeds 950 OC, there is a considerable reduction in the life of the aluminide coatings. Incorporation of various alloying elements such as chromium and silicon and of inert oxides such as Al,O, into the basic NiAl compound were tried. The improvement was only marginal. One of the successful breakthroughs came with the advent of noble metal plating (platinum, rhenium etc.) followed by aluminizing [16]. The aluminium reacts with platinum to form platinum aluminide, in addition to the formation of NiAl. This process improves the life of the coatings by a factor of 3. Only limited investigations have been carried out on the reasons behind such improvement.
147
5. Overlay coatings Before the platinum aluminide process gained wide commercial acceptance, the concept of forming overlay coatings based on MCrAlY (M = Fe, Co or Ni) compounds was advocated. In these coatings, aluminium is required to form protective Al,O, scale. Chromium promotes the formation of Al,O, scale, as well as improving the hot-corrosion resistance. The concept that small additions of rare earth elements (maximum of 1 wt.%) improve the adherence of oxide scales emanated from oxidation studies of dispersionstrengthened materials. This idea was adopted and the overlay coating compositions modified by the addition of yttrium. The reasons for such a beneficial effect are not well established [ 21. These coatings have at least four elements and their deposition by the pack method requires each element to be deposited separately. It is difficult to alter the partial pressures of the halides of the various elements which are dictated by the thermodynamics of the pack to deposit them in a single cycle. Thus the electron beam evaporation process was further developed to form these coatings 1171. In this method, an alloy ingot is melted by an electron beam and the evaporated elements deposit onto the blades. The line-of-sight nature of this process, unlike the pack process, demands complex movements for the blades to obtain uniform coatings. As the vapour pressures of the elements differ, suitable compositional adjustment must be incorporated in the starting ingot. The deposition rates are high (10 pm min-‘) and this is still an expensive process because of the high investment cost. Also, the control of coating composition is difficult when it contains elements with high vapour pressures. Consequently, the electron beam process is not appropriate for trying various compositions. Owing to its simplicity in operation and economy, the low pressure plasma spraying process has started to replace the electron beam evaporation method. There have been attempts to form the overlay coatings by sputtering [ 181, ion plating and a combination of slurry spraying and pack aluminizing which is called the controlled-composition reaction-sintered coating process [ 191. These methods have not taken off as viable production routes even though there are no major technical difficulties. A recent account of these and other alternative processes has been presented by Restall and Wood [20]. The plasma spraying process is normally carried out in air at atmospheric pressure. In this process, the alloy powder is melted in a high power d.c. plasma flame formed in a flowing Ar-Hz mixture [21]. The plasma torch acts not only as a heat source, but also accelerates the powders to velocities of up to 300 m s-l. In the short residence time of a few milliseconds, the powder particles are melted and they strike and flatten on the object to be coated. By repeated movement of the gun, the coating is built up of layer on layer of these flattened droplets. There are a number of process variables that need to be controlled to obtain coatings of high density. The development of this process has so far proceeded in an
148
empirical manner. Recently attempts have been made to calculate the timetemperature history of particles in the plasma from a fundamental consideration of heat and momentum transfer [22]. The very nature of this process yields coatings with a porosity in the range 2% - 10%. Spraying in air also causes oxidation of the powders because of turbulent mixing of the surrounding air with the plasma gas. Initially, attempts were made to eliminate these problems by spraying in an argon chamber. This was only partially successful and the next logical step was the development of spraying in a low pressure chamber (50 - 100 Torr). The low pressure chamber offers a number of advantages especially for spraying metallic powders [23]. The plasma becomes elongated in the low pressure chamber compared with that at atmospheric pressure. This implies a gradual drop in the plasma flame temperature. The main advantage of this process lies in the fact that the turbine blade can now be heated to high temperatures (800 - 900 “C) without the problem of oxidation. As the plasma gas from the gun expands into a low pressure chamber, the plasma gas velocity also increases compared with that for spraying in air. The molten droplets hit the hot blade with a high velocity and lose the particle shape, thus forming a recrystallized coating. The dramatic difference in the coating structure can be seen in Fig. 6 [24]. In this process a numerically controlled five-axis mainipulator system is required for moving the blade to obtain uniform coatings on the entire surface. Alternately, a robot has been used. A wide variety of compositions have been tried in the MCrAlY series based on iron, nickel, cobalt and their combinations. Some of the coatings are single phase, but most of them are two phase, containing NiAl or CoAl precipitates in an Ni-Cr or Co-Cr matrix. All these coatings rely on the
(a) Fig. 6. Structure of low pressure plasma
plasma
sprayed;
(b)
149
formation of a protective AlzO, scale. The powders for spraying are made by the argon atomization process and hence it is easy to obtain powders of different compositions. These powders can be sprayed without any difficulty using the low pressure plasma spraying process. Such flexibility helps to optimize the coating composition for the best protection. The coating ductility is also affected by the composition. Figure 7 shows how, by decreasing the chromium and aluminium contents of a CoCrAlY coating, the ductile-brittle transition temperature can be decreased from 540 “C (1000 “F) to 95 “C (200 “F). This is due to a smaller volume fraction of the brittle CoAl phase dispersed in a more ductile cobalt matrix for the Co-15.5wt.%Cr-10.2wt.%Al-0.3wt.%Y coating. The superiority of the overlay coatings to the nickel aluminide is also shown in the Fig. 10. Even though the overlay coatings by their very nature are independent of the substrate, a small diffusion zone of about 10% of the coating thickness is formed by post-coating heat treatment to improve the bonding. During service, the substrate elements and the coating elements interdiffuse making it necessary to evaluate the long-term diffusional stability of these coatings. The oxidation and hot-corrosion properties of these coatings have also been found to be dependent on the substrate composition. Pettit and Goward [ 251 have reported that the NiCoCrAlY coating (18wt.%Cr, 23wt.%Co, 12wt.%Al, 0.5wt.%Y, remainder Ni) shows better oxidation resistance on a hafnium-containing alloy. This has been attributed to greater adherence of the Al,O, scale which has been pegged by HfO,. Apparently the hafnium from the alloy has diffused through the coating to form these pegs. Attempts are being made to improve the oxidation and hot-corrosion resistance of the overlay coatings further by adding hafnium, tantalum and silicon to the base NiCoCrAlY composition. However, caution is required in translating the properties of a particular coating on a specific alloy to other types of alloys. The properties are also known to differ depending upon the specific environment in which the gas turbines operate.
0
0
200
400
Fig. 7. The ductility
600
BOO
1000
Tempetalure “F
of NiAl and CoCrAly
1 ‘loo
1600
coatings
[lo].
1200
150
6. Thermal
barrier coatings
The next generation of advanced coatings were based on ceramic materials which acted as a thermal barrier between the hot combustion gas and the turbine metal surface [26]. This approach has already been successfully used in combustor and after-burner applications. The reliability and durability of these coatings need to be improved for successful applications on turbine blades. The incentive to develop these coatings is high as they can decrease the metal temperature by up to 200 “C and thus can improve the creep and oxidation life. Alternatively, the cooling air consumption can be reduced by 6% for the same metal temperature. This will have a significant impact on the efficiency of the engine. Plasma spraying is the most attractive method of forming these ceramic coatings which are 200 - 300 pm thick. They are applied on a metallic bond coat for better adherence. The main limitation of the ceramic coatings is their low strain tolerance. The mismatch in thermal expansion between the coating and the underlying substrate is a major factor leading to the spalling of the coating. These coatings are porous and thus the oxidation of the underlying metallic bond coat also leads to spalling [ 271. Y,O,-stabilized ZrO, compounds have shown considerable promise as coating materials [ 281. A number of approaches have been tried to improve the thermal shock resistance of zirconia coatings. These include segmented structure, microcracked structure, graded coatings and composition control to obtain a cubic plus tetragonal two-phase structure [25]. In the presence of molten salt, it is reported that Y,O, reacts leading to destabilization of the cubic phase. Even when there is no reaction, as in the case of MgO. ZrO,, the molten salt can penetrate the coating and oxidize the underlying bond coat. Figure 8 shows 0
Fig. 8. Hot-corroded micrograph with salt;
Yz03-ZrOa (b) electron
coating and X-ray
on CoCrAlY images.
bond
coat
[27]:
(A)
optical
151
such an attack which leads to spalling of the MgO* ZrO, coating [27]. Already coated with nozzle vane platforms thermal barriers are in use in some aero engines [29]. Ceramic-coated vanes have been successfully engine tested and it may not be long before ceramic-coated blades propel aircraft. Another process that is competing to form these coatings is electron beam evaporation [30]. This process is attractive as it forms a strain-tolerant columnar structure of ZrO, which is denser than the plasma-sprayed coatings. The increased density lowers the insulative capacity, but increases the erosion resistance of these coatings. Recently the use of ion or plasma assistance in the electron beam physical vapour deposition of stabilized zirconia coatings has been shown to be highly beneficial in structure control and adhesion enhancement [ 31, 321.
7. Future
developments
The properties of the coatings depend critically upon the process, as the coating structure varies with the method of formation. The coating processes have their own inherent limitations. One of the viable methods to overcome some of these limitations is by post-coating treatments. The well known among these are high temperature diffusion treatment to improve bonding and shot-peening to induce compressive stresses and to obtain a smooth surface [33]. Of late, the laser has been recognized as a versatile tool to remelt the coatings. Laser remelting can help to remove the pores that are present in plasma-sprayed coatings or help to modify the structure of a fully dense coating. In view of its great potential, a summary of its use in the context of turbine blade coatings is presented below. 7.1. Metallic coatings One of the major limitations of diffusion aluminide coatings is its limited ductility. Hence, during temperature cycling, the thermal stress can induce cracking of the coated layer. These cracks have been found to extend into the substrate and to damage the component [29]. The mechanical effects of the coatings are complex. An overview of these problems in general has been presented by Evans et al. [ 341 and in particular reference to turbine blades by Schneider and Griinling [35]. The role of residual stresses and the effect of various mechanisms of stress relaxation have been outlined. The role of interface defects in initating buckling and eventual spalling of the coating has also been analysed [35]. The effect of coating composition and the type of coating on ductility are shown in Fig. 7. This in turn affects the thermomechanical fatigue life as shown in Fig. 9 1351. As shown in Fig. 7, in the case of MCrAlY coatings, the ductility can be altered by changing the aluminium content. Such a dramatic change is not possible with NiAl coatings. Laser remelting of NiAl coatings offers a distinct possibility of increasing the ductility. Such an increase has been
152
Fig. 9. Thermomechanical
fatigue
properties
of NiAl and CoCrAlY
coatings
[33].
demonstrated in NiAl-based alloys through refining the grain size by thermomechanical processing [ 361. For grain sizes below 20 pm, elongations greater than 10% were obtained at room temperature. There is some controversy whether such an increase can be obtained by rapid solidification, which can reduce the ordering in NiAl. Increased ductility by reduced ordering has been well established for Ni,Al [3’7]. In view of the above findings, it would be worth while to laser melt the NiAl coatings to increase their ductility. As-formed NiAl coatings are unstable as there is a gradient in the concentration of aluminium from the surface (highest) to the interior. Laser melting would remove this concentration gradient and thereby increase the microstructural stability. However, care must be taken to ensure an adequate amount of aluminium in the NiAl coatings after laser melting in order not to impair its resistance to hot corrosion [ 291. Laser melting is beneficial for air-plasma-sprayed MCrAlY types of overlay coatings. It offers an alternative route to form a dense and oxidefree coating which can replace the more expensive methods such as electron beam evaporation and low pressure plasma spraying. This approach was investigated by Dalliare and Cielo [ 38, 391 and Capp and Rigsbee [40] by melting Hastelloy, nickel and Stellite coatings by laser. Bhat et al. [41] have successfully sealed pores by melting a thin top layer (20 - 25 E.trn thick) of plasma-sprayed NiCrAlY coatings using a CO, laser operating in the continuous wave mode [ 411. Detailed investigations into the potential and limitations of laser melting of MCrAlY coatings have been carried out by Sivakumar and Mordike [42]. It was found that there is no problem of absorption of laser radiation and a desired thickness can be melted. Figure 10 shows the typical cross-sectional morphologies of the coatings that were melted partially and fully. However, there is a depletion of aluminium and yttrium in the laser-melted coatings owing to the formation of oxides of
153
(a) Fig. 10. Cross-sectional melted; (B) fully melted.
(b) micrographs
of
laser-melted
NiCoCrAlY
coatings:
(A) partially
the above elements. These oxides can be removed as they are driven to the surface by convection currents during laser melting. Figure 11 shows the typical surface appearance of laser-melted NiCoCrCrAlY coatings with the laser operating in the continuous wave and the pulsed modes. The dark regions are oxides of aluminium and yttrium. This loss is reflected by poor cyclic oxidation resistance. Hence the initial plasmasprayed coating compositions need to be adjusted to take account of the above losses and thereby to obtain optimum properties. It is also possible to carry out the same process with an electron beam [43]. The laser is more versatile as it requires no vacuum for its operation. Another approach is to avoid the plasma spraying step and to melt directly MCrAlY powders, attached to the substrate by a binder, using a laser [44]. The control over the melt thickness, uniformity and the extent of dilution with substrate are difficult and more work needs to be carried out to make this process viable. Another possibility is laser spraying or cladding, in which the powders are fed directly into the laser-melted bath by a delivery system [45 - 471. When the powders are delivered coaxially into the laser beam, they are accelerated to a few metres per second. Possible mechanisms of transfer of momentum are discussed by Schaefer et al. [45]. These velocities by themselves are not sufficient to form dense coatings, as these are very low compared with the velocities of the droplets reached during plasma spraying (a few hundred metres per second). Hence the laser must melt the particles that are falling onto the substrate along with a part of the substrate to form a fully dense coating. It is also possible to spray the
(a) Fig. Il. Surface morphology mode; (B) pulsed mode.
(b) of laser-melted
NiCoCrAlY
coatings:
(A) continuous
wave
powder at an angle close to the substrate and to use the laser to melt the powder and alloy at the same time as the substrate [ 471. Because of process control problems, these technologies are at various stages of development. At present, they can be used for less critical applications. Laser melting of fully dense MCrAlY coatings formed by electron beam evaporation or low pressure plasma spraying has also been found to be beneficial in increasing the cyclic oxidation resistance [ 48, 491. Bornstein and Smeggil [48] found a thin tenaceous oxide scale followed by a singlephase zone in the case of laser-melted NiCrAlY coatings after testing in a burner rig at 985 “C for 100 h. The corresponding non-laser-melted specimen showed some spalling of the oxide, and the matrix below the oxide consisted of a two-phase structure of NiAl and Ni-Cr solid solution. It was also observed that there was an accumulation of yttrium beneath the scale, whereas no such yttrium was detected in the laser-melted specimens. In their later study, it was reported that the as-formed coating contained discrete yttrium-enriched particles and yttrium could not be detected on the laser-melted surface [49]. Also the laser-melted surface had a fine micron-sized distribution of NiAl phase as opposed to a coarse distribution of the same in the as-formed coating. Laser melting made the coated surface more homogeneous and thus promoted the formation a uniform, thin and adherent oxide. Further investigations are needed to understand the reason behind such an improvement. Also interesting would be the hot corrosion properties of such laser-melted layers.
155
7.2. Ceramic coatings In the development of thermal barrier ceramic coatings, laser melting is likely to play a key role. The laser is uniquely suited to remelt a part of the ceramic coatings with a minimal input of heat. Such a remelted layer would be fully dense and so would prevent the corrosive gases and salts from attacking the coating and the underlying substrate. A few investigations towards achieving the same have been reported [ 50 - 541. The major problem which is yet to be solved is the cracking of the laser-melted layer owing to non-uniform cooling [55]. Figure 12 shows the typical surface morphology of laser-melted surfaces of the most popular 7Y,O,*ZrO, coatings and the corresponding cross-section showing the transverse cracks
(b) Fig. 12. Laser-melted
7Y203*ZrOz
coatings:
(a) surface
morphology;
(b) cross-section
156
across the melted layer. Detailed investigations on the various aspects of laser melting of different ceramic coatings, namely Zr02, Al,O, and Ti02, have recently been carried out [56]. In this study it was shown how the extent of cracking can be reduced by proper choice of laser operation
(b)
(a)
(cl Fig. 13. Surface temperature;(c)
morphology laser melted,
of A1203 800 “C.
coatings:
(a) as sprayed;
(b) laser melted,
room
157
parameters. Also by a simple analysis of stresses that arise during cooling it was shown that it is possible to melt a thin surface layer of Al,O, and TiO,? without cracking by preheating the coating to 800 “C!. Figure 13 shows the typical surface morphology of an Al,O, surface that was laser melted at room temperature and the surface without cracks which was melted with preheating to 800 “C. Figure 14 shows the benefits in terms of avoidance of spalling of the coatings in a laboratory hot-corrosion test with 95%Na,SO,5%NaCl salt at 900 “C for 100 h [56]. The above solution is not practical for ZrO, systems, as the coatings have to be heated to temperatures above 1000 “C to avoid cracking. It may be possible, however, to reduce this temperature by altering the Y,O,stabilized ZrOz compositions to increase the fracture toughness. One possibility is to add up to 20wt.% Al,Os to Y,O,-stabilized ZrO, to increase the toughness from 6 to 7.6 MN m-” [57]. The other option is to choose a low Y203 content ZrO, composition, in which the tetragonal-to-monoclinic phase transformation takes place during cooling. The expansion associated with the above phase change may counteract the thermal contraction. It is also feasible to try Zr02* SiO* coatings, in which the presence of SiO, helps to relieve the stresses and thereby prevents cracking [ 541.
10 mm Fig. 14. Hot-corroded A1203: laser melted at 800 “C.
0, as sprayed;
1, laser
melted
at room
temperature;
2, 3,
8. Conclusion It has been shown how the demand for better performance has led to the development of new coatings and coating methods. Table 3 gives a summary of the coating developments for turbine blades. This development encompasses an interaction between the physical metallurgy of the coating and the processing. A fundamental understanding of the mechanism of formation of different coatings and their mode of degradation is essential for the design of suitable coatings. In the case of high temperature coatings, the properties are considerably affected by the structure and chemistry of the substrate. It is hence prudent to design a coated alloy as a single entity
158 TABLE Coatings
3 for turbine
blades
Coating type
Coating method
Coating phase
Mode of protection
Diffusion
Pack aluminizing
NiAl
A1203
Overlay
Electron beam evaporation
MCrAlY (M = Ni, Co, Fe)
More spallresistant
Overlay
Low pressure plasma spraying
MCrAlY (M = Ni, Co, Fe)
More spallresistant
Thermal
Air plasma spraying Electron beam processes
Partially stabilized ZrOz
Insulative
barrier
Limitations
Hot corrosion Brittleness Diffusional stability Hot corrosion Thermal fatigue
Al203
Hot corrosion Thermal fatigue
A1203
Thermal spalling, oxidation and hot corrosion of bond coat
as opposed to the current practice of trying to marry the existing ones to the various alloys, with minor modifications here and there. The future trend will be to go for more post-coating modifications to suit specific applications, rather than major changes in coating processes.
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