||This article includes a list of references, but its sources remain unclear because it has insufficient inline citations. (July 2009)|
A superalloy, or high-performance alloy, is an alloy that exhibits several key characteristics: excellent mechanical strength, resistance to thermal creep deformation, good surface stability and resistance to corrosion or oxidation. The crystal structure is typically face-centered cubic austenitic. Examples of such alloys are Hastelloy, Inconel, Waspaloy, Rene alloys, Haynes alloys, Incoloy, MP98T, TMS alloys, and CMSX single crystal alloys.
Superalloy development has relied heavily on both chemical and process innovations. Superalloys develop high temperature strength through solid solution strengthening. An important strengthening mechanism is precipitation strengthening which forms secondary phase precipitates such as gamma prime and carbides. Oxidation or corrosion resistance is provided by elements such as aluminium and chromium.
The primary application for such alloys is in turbine engines, both aerospace and marine.
- 1 Chemical development
- 2 Process development
- 3 Metallurgy of superalloys
- 4 Coating of superalloys
- 5 Research and development of new superalloys
- 6 See also
- 7 References
- 8 External links
Because these alloys are intended to be used for high temperature applications, in addition to these materials being able to withstand loading at temperatures near their melting point, their creep and oxidation resistance are of primary importance. Ni based superalloys have emerged as the material of choice for these applications. The properties of these Ni based superalloys can be tailored to a certain extent through the addition of many other elements, both common and exotic, including not only metals, but also metalloids and nonmetals; chromium, iron, cobalt, molybdenum, tungsten, tantalum, aluminium, titanium, zirconium, niobium, rhenium, yttrium, vanadium, carbon, boron or hafnium are some examples of the alloying additions used. Each of these additions has been chosen to serve a particular purpose in optimizing the properties for high temperature application.
Creep resistance is dependent on slowing the speed of dislocation motion within a crystal structure. In modern Ni based superalloys the γ’-Ni3(Al,Ti) phase present acts as a barrier to dislocation motion. For this reason, this γ’ intermetallic phase, when present in high volume fractions, drastically increases the strength of these alloys due to its ordered nature and high coherency with the γ matrix. The chemical additions of aluminum and titanium promote the creation of the γ’ phase. The γ’ phase size can be precisely controlled by careful precipitation strengthening heat treatments. Many superalloys are produced using a two-phase heat treatment that creates a dispersion of cuboidal γ’ particles known as the primary phase, with a fine dispersion between these known as secondary γ’. In order to improve the oxidation resistance of these alloys, Al, Cr, B, and Y are added. The Al and Cr form oxide layers that passivate the surface and protect the superalloy from further oxidation while B and Y are used to improve the adhesion of this oxide scale to the substrate. Cr, Fe, Co, Mo and Re all preferentially partition to the γ matrix while Al, Ti, Nb, Ta, and V preferentially partition to the γ’ precipitates and solid solution strengthen the matrix and precipitates respectively. In addition to solid solution strengthening, if grain boundaries are present, certain elements are chosen for grain boundary strengthening. B and Zr tend to segregate to the grain boundaries which reduces the grain boundary energy and results in better grain boundary cohesion and ductility. Another form of grain boundary strengthening is achieved through the addition of C and a carbide former, such as Cr, Mo, W, Nb, Ta, Ti, or Hf, which drives precipitation of carbides at grain boundaries and thereby reduces grain boundary sliding.
While Ni based superalloys are excellent high temperature materials and have proven very useful, Co based superalloys potentially possess superior hot corrosion, oxidation, and wear resistance as compared to Ni based superalloys. For this reason, efforts have also been put into developing Co based superalloys over the past several years. Despite that, traditional Co based superalloys have not found widespread usage because they have a lower strength at high temperature than Ni based superalloys. The main reason for this is that they appear to lack the γ’ precipitation strengthening that is so important in the high temperature strength of Ni based superalloys. However, there has been a recent discovery of a stable γ’-Co3(Al,W) intermetallic compound with the L12 structure. The two-phase microstructure consists of cuboidal γ’ precipitates embedded in a continuous γ matrix and is therefore morphologically identical to the microstructure observed in Ni based superalloys. Like in the Ni based system, there is a high degree of coherency between the two phases which is one of the main factors resulting in the superior strength at high temperatures. This provides a pathway for the development of a new class of load-bearing Co based superalloys for application in severe environments. In addition to the fact that many of the properties of these new Co based superalloys could be better than those of the more traditional Ni based ones, Co also has a higher melting temperature than Ni. Therefore, if the high temperature strength could be improved, the development of novel Co based superalloys could allow for an increase in jet engine operation temperature resulting in an increased efficiency.
The historical developments in superalloy processing have brought about considerable increases in superalloy operating temperatures. Superalloys were originally iron based and cold wrought prior to the 1940s. In the 1940s investment casting of cobalt base alloys significantly raised operating temperatures. The development of vacuum melting in the 1950s allowed for very fine control of the chemical composition of superalloys and reduction in contamination and in turn led to a revolution in processing techniques such as directional solidification of alloys and single crystal superalloys.
There are many forms of superalloy present within the gas turbine engine. In order to have fracture resistance, the disks of the high pressure turbine are polycrystalline, which are usually cast and then forged into shape. The cast discs have a large columnar grain structure and contain chemical segregation. The polycrystalline discs can also be made by powder metallurgy, where fine powders are hot-isostatically pressed, extruded, and forged into shape. On the other hand, turbine blades are usually monocrystalline or single crystal. Single crystal blades are free of γ/γ’ grain boundaries, which allow for increase in creep resistance. Turbine blades can also be polycrystalline, which are made via investment casting. The polycrystalline blades can contain either columnar grains or equiaxed grains. Columnar grain structured blades are created using directional solidification techniques and have grains parallel to the major stress axes while equiaxed grain structured blades are prone to creep deformation.
Single-crystal superalloys (SX or SC superalloys) are formed as a single crystal using a modified version of the directional solidification technique, so there are no grain boundaries in the material. The mechanical properties of most other alloys depend on the presence of grain boundaries, but at high temperatures, they would participate in creep and must be replaced by other mechanisms. In many such alloys, islands of an ordered intermetallic phase sit in a matrix of disordered phase, all with the same crystalline lattice. This approximates the dislocation-pinning behavior of grain boundaries, without introducing any amorphous solid into the structure.
Metallurgy of superalloys
The microstructure of most precipitation strengthened nickel-base superalloys consists of the gamma matrix, and of intermetallic γ' precipitates. The γ-phase is a solid solution with a face-centered crystal (fcc) lattice and randomly distributed different species of atoms. By contrast, the γ'-phase has an ordered crystalline lattice of type LI2. Modern alloys typically contain about 70% by volume fraction of cube-like γ' precipitates whose edge length is about 0.5 μm.
In pure Ni3Al phase atoms of aluminium are placed at the vertices of the cubic cell and form the sublattice A. Atoms of nickel are located at centers of the faces and form the sublattice B. The phase is not strictly stoichiometric. There may exist an excess of vacancies in one of the sublattices, which leads to deviations from stoichiometry. Sublattices A and B of the γ'-phase can solute a considerable proportion of other elements. The alloying elements are dissolved in the γ-phase as well. The γ'-phase hardens the alloy through an unusual mechanism called the yield strength anomaly. Dislocations dissociate in the γ'-phase, leading to the formation of an anti-phase boundary. It turns out that at elevated temperature, the free energy associated with the anti-phase boundary (APB) is considerably reduced if it lies on a particular plane, which by coincidence is not a permitted slip plane. One set of partial dislocations bounding the APB cross-slips so that the APB lies on the low-energy plane, and, since this low-energy plane is not a permitted slip plane, the dissociated dislocation is now effectively locked. By this mechanism, the yield strength of γ'-phase Ni3Al actually increases with temperature up to about 1000 °C, giving superalloys their currently unrivalled high-temperature strength.
Initial material selection for blade applications in Gas Turbine engines included alloys like the Nimonic series alloys in the 1940s. The early Nimonic series incorporated γ' Ni3(Al,Ti) precipitates in a γ matrix, as well as various metal-carbon carbides (e.g. Cr23C6) at the grain boundaries for additional grain boundary strength. Turbine blade components were forged until vacuum induction casting technologies were introduced in the 1950s. This process significantly improved cleanliness, reduced defects, and increased the strength and temperature capability of the material.
Modern superalloys were developed in the 1980s with the advent of single crystal, or monocrystal, solidification techniques (see Bridgman technique) for superalloys that enable grain boundaries to be entirely eliminated from a casting. Because the material contained no grain boundaries, carbides were unnecessary as grain boundary strengthers and were thus eliminated. Additionally, the volume fraction of the γ' precipitates increased to about 50-70%. The first generation superalloys incorporated increased Aluminium, Titanium, Tantalum, and Niobium content in order to increase the γ' volume fraction in these alloys. Examples of first generation superalloys include: PWA1480, René N4 and SRR99.
The second and third generation superalloys introduced about 3 and 6 weight per cent Rhenium, for increased temperature capability. Examples of second generation superalloys include PWA1484, CMSX-4 and René N5. Third generation alloys include CMSX-10, and René N6. Fourth, Fifth, and even Sixth generation superalloys have been developed which incorporate Ruthenium additions, making them more expensive still than the prior generation's Re-containing alloys.
The current trend is to avoid very expensive and very heavy elements. A possible remedy to this is Eglin steel, a budget material with compromised temperature range and chemical resistance. It does not contain rhenium or ruthenium and its nickel content is limited. To reduce fabrication costs, it was chemically designed to melt in a ladle (though with improved properties in a vacuum crucible). Also, conventional welding and casting is possible before heat-treatment. The original purpose was to produce high-performance, inexpensive bomb casings, but the material has proven widely applicable to structural applications, including armor.
In addition, it is often beneficial for grain boundaries that the nickel-base superalloy contains carbides (or boron or zirconium) for improvements in creep strength. Where the carbides (e.g. MC where M is a metal and C is a carbon atom) are precipitated at the grain boundaries, they act to pin the grain boundaries and improve the resistance to sliding and climbing and migration that would occur during creep diffusion. However if they precipitate as a continuous grain boundary film, the fracture toughness of the alloy may be reduced, together with the ductility and rupture strength.
Coating of superalloys
Superalloy products that are subjected to high working temperatures and corrosive atmosphere (such as high pressure turbine region of jet engines) are coated with various kinds of coating. Several kinds of coating process are applied: pack cementation process, gas phase coating (both are a type of chemical vapor deposition (CVD)), thermal spraying, and physical vapor deposition. In most cases, after the coating process near-surface regions of parts are enriched with aluminium, the matrix of the coating being nickel aluminide.
Pack cementation process
The pack cementation process is carried out at lower temperatures, about 750 °C. The parts are loaded into boxes that contain a mixture of powders: active coating material, containing aluminum, activator (chloride or fluoride), and thermal ballast, like aluminum oxide. At high temperatures the gaseous aluminum chloride is transferred to the surface of the part and diffuses inside (mostly inward diffusion). After the end of the process the so-called "green coating" is produced, which is too thin and brittle for direct use. A subsequent diffusion heat treatment (several hours at temperatures about 1080 °C) leads to further inward diffusion and formation of the desired coating.
Thermal spraying is a process of applying coatings by heating a feedstock of precursor material and spraying it on a surface. Different specific techniques are used depending on desired particle size, coat thickness, spray speed, desired area, etc. The coatings applied by thermal spraying of any kind, however, rely on adhesion to the surface. As a result, the surface of the superalloy must be cleaned and prepared, usually polished, before application of the thermal coating.
Of the various thermal spray methods, one of the more ideal and commonly used techniques for coating superalloys is plasma spraying. This is due to the versatility of usable coatings, and the high-temperature performance of plasma-sprayed coatings. Plasma spraying can accommodate a very wide range of materials, much more so than other techniques. As long as the difference between melting and decomposition temperatures is greater than 300 Kelvin, a material can be melted and applied as a coating via plasma spraying.
Gas phase coating
This process is carried out at higher temperatures, about 1080 °C. The coating material is usually loaded onto special trays without physical contact with the parts to be coated. The coating mixture contains active coating material and activator, but usually does not contain thermal ballast. As in the pack cementation process, the gaseous aluminium chloride (or fluoride) is transferred to the surface of the part. However, in this case the diffusion is outwards. This kind of coating also requires diffusion heat treatment.
The bond coat adheres the thermal barrier coating to the superalloy substrate. Additionally, the bond coat provides oxidation protection and functions as a diffusion barrier against the motion of substrate atoms towards the environment.
There are five major types of bond coats, the aluminides, the platinum-aluminides, MCrAlY, cobalt-cermets, and nickel-chromium.
For the aluminide bond coatings, the final composition and structure of the coating depends on the composition of the substrate. Aluminides also lack ductility below 750 °C, and exhibit a limited by thermomechanical fatigue strength.
The Pt-aluminides are very similar to the aluminide bond coats except for a layer of Pt (5-10 μm) deposited on the blade. The Pt is believed to aid in oxide adhesion and contributes to hot corrosion. The cost of Pt plating is justified by the increased blade life span.
The MCrAlY is the latest generation of bond coat and does not strongly interact with the substrate. Normally applied by plasma spraying, MCrAlY coatings are secondary aluminum oxide formers. This means that the coatings form an outer layer of chromium oxide (chromia), and a secondary aluminum oxide (alumina) layer underneath. These oxide formations occur at high temperatures in the range of those that superalloys usually encounter. The chromia provides oxidation and hot-corrosion resistance. The alumina controls oxidation mechanisms by limiting oxide growth by self-passivating. The yttrium enhances the oxide adherence to the substrate, and limits the growth of grain boundaries (which can lead to flaking of the coating). Investigation indicates that addition of rhenium and tantalum increases oxidation resistance.
Cobalt-cermet based coatings consisting of materials such as tungsten carbide/cobalt can be used due to excellent resistance to abrasion, corrosion, erosion, and heat. These cermet coatings perform well in situations where temperature and oxidation damage are significant concerns, such as boilers. One of the unique advantages of cobalt cermet coatings is a minimal loss of coating mass over time, due to the strength of carbides within the mixture. Overall, cermet coatings are useful in situations where mechanical demands are equal to chemical demands for superalloys.
Nickel-chromium coatings are used most frequently in boilers fed by fossil fuels, electric furnaces, and waste incineration furnaces, where the danger of oxidizing agents and corrosive compounds in the vapor must be dealt with. The specific method of spray-coating depends on the composition of the coatings. Nickel-chromium coatings that also contain iron or aluminum perform much better (in terms of corrosion resistance) when they are sprayed and laser glazed, while pure nickel-chromium coatings perform better when thermally sprayed exclusively.
Research and development of new superalloys
The availability of superalloys during past decades has led to a steady increase in the turbine entry temperatures and the trend is expected to continue. Sandia National Laboratories is studying a new method for making superalloys, known as radiolysis. It introduces an entirely new area of research into creating alloys and superalloys through nanoparticle synthesis. This process holds promise as a universal method of nanoparticle formation. By developing an understanding of the basic material science behind these nanoparticle formations, there is speculation that it might be possible to expand research into other aspects of superalloys.
There may be considerable disadvantages in making alloys by this method. About half of the use of superalloys is in applications where the service temperature is close to the melting temperature of the alloy. It is common therefore to use single crystals. The above method produces polycrystalline alloys, which suffer from an unacceptable level of creep.
Future paradigm in alloy development focus on reduction of weight, improving oxidation and corrosion resistance while maintaining the strength of the alloy. Furthermore, with the increasing demand for turbine blade for power generation, another focus of alloy design is to reduce the cost of super alloys.
- Reed, Roger C. The Superalloys: Fundamentals and Applications. Cambridge, UK: Cambridge UP, 2006.
- Klein, L., Y. Shen, M. S. Killian, and S. Virtanen. "Effect of B and Cr on the High Temperature Oxidation Behavior of Novel γ/γ′Strengthened Co-base Superalloys." Corrosion Science 53 (2011): 2713-720.
- Shinagawa, K., Toshihiro Omori, Katsunari Oikawa, Ryosuke Kainuma, and Kiyohito Ishida. "Ductility Enhancement by Boron Addition in Co–Al–W High-temperature Alloys." Scripta Materialia 61.6 (2009): 612-15.
- Sato, J. "Cobalt-Base High-Temperature Alloys." Science 312.5770 (2006): 90-91.
- Suzuki, A., Garret C. DeNolf, and Tresa M. Pollock. "Flow Stress Anomalies in γ/γ′ Two-phase Co–Al–W-base Alloys." Scripta Materialia 56.5 (2007): 385-88.
- C. Sims, N. Stoloff, W. Hagel, Superalloys II: High Temperature Materials for Aerospace and Industrial Power, 1987, John Wiley & Sons
- R.C. Reed, The Superalloys. Fundamentals and Applications
- D. Bombač, M. Fazarinc, G. Kugler, S. Spajić, Microstructure development of Nimonic 80A superalloys during hot deformation, Materials and Geoenvironment, 55:3 (2008) 319-328.
- G. R. Heath, P. Heimgartner, G. Irons, R. Miller, S. Gustafsson, Mater. Sci. Forum 1997, 251–54, 809
- O. Knotek, Handbook of Hard Coatings: Deposition Technologies, Properties and Applications, Ed. R. F. Bunshah, Noyes Pub. Park Ridge, New Jersey, U. S. A./William Andrew Publishing, LLC, Norwich, New York, U.S.A. 2001.
- P. Niranatlumpong, C. B. Ponton, H. E. Evans, Oxid. Met. 2000, 53, 241
- P. Fauchais, A. Vardelle, M. Vardelle, Modelling of Plasma Spraying of Ceramic Films and Coatings, Ed. Vinenzini, Pub. Elsevier State Publishers B.V 1991.
- B. M. Warnes, Surf. Coat. Technol. 2003, 163–164, 106.
- H. M. Tawancy, N. M. Abbas, A. Bennett, Surf. Coat. Technol. 1994, 68–69, 10.
- C. Chuanxian, H. Bingtang, L. Huiling, Thin Solid Films 1984, 118, 485.
- Y. Kawahara, Mater. High Temp. 1997, 14, 261.
- Y. Longa-Nava, M. Takemoto, Corros. 1992, 48, 599.
- Levitin, Valim (2006). High Temperature Strain of Metals and Alloys: Physical Fundamentals. WILEY-VCH. ISBN 978-3-527-31338-9.