Volume 01 - Properties and Selection Irons, Steels, and High-Performance Alloys Part 16 potx

M-22 (a) 515 (75) 385 (56) 395 (57) 285 (41) 200 (29) 130 (19) MAR-M 421 (a) 450 (65) 305 (44) 310 (46) 215 (31) 125 (18) 83 (12) MAR-M 432 (a) 435 (63) 330 (48) 295 (40) 215 (31) 140 (20) 97 (14) MC-102 (a) 195 (28) 145 (21) 145 (21) 105 (15) . . . . . . Nimocast 90 (a) 160 (23) 110 (17) 125 (18) 83 (12) . . . . . . Nimocast 242 (a) 110 (16) 83 (12) 90 (13) 59 (8.6) 45 (6.5) . . . Udimet 500 (a) 330 (48) 240 (35) 230 (33) 165 (24) 90 (13) . . . Udimet 710 (a) 420 (61) 325 (47) 305 (44) 215 (31) 150 (22) 76 (11) CMSX-2 (b) . . . . . . . . . 345 (50) . . . 170 (25) GMR-235 (b) . . . . . . . . . 180 (26) . . . 75 (11) IN-939 (b) . . . . . . . . . 195 (28) . . . 60 (9) MM 002 (b) . . . . . . . . . 305 (44) . . . 125 (18) IN-713 Hf (MM 004) (b) . . . . . . . . . 205 (30) . . . 90 (13) René 125 Hf (MM 005) (b) . . . . . . . . . 305 (44) . . . 115 (17) SEL-15 (b) . . . . . . . . . 295 (43) . . . 75 (11) UDM 56 (b) . . . . . . . . . 270 (39) . . . 125 (18) (a) Ref 3. (b) Ref 1 Table 6 Stress-rupture strengths for selected polycrystalline cobalt-base superalloys Rupture stress Alloy At 815 °C (1500 °F) At 870 °C (1600 °F) At 980 °C (1800 °F) At 1095 °C (2000 °F) 100 h MPa (ksi) 1000 h MPa (ksi) 100 h MPa (ksi) 1000 h MPa (ksi) 100 h MPa (ksi) 1000 h MPa (ksi) 100 h MPa (ksi) 1000 h MPa (ksi) HS-21 150 (22) 95 (14) 115 (17) 90 (13) 60 (9) 50 (7) . . . . . . X-40 (HS-31) 180 (26) 140 (20) 130 (19) 105 (15) 75 (11) 55 (8) . . . . . . MAR-M 509 270 (39) 225 (33) 200 (29) 140 (20) 115 (17) 90 (13) 55 (8) 41 (6) FSX-414 150 (22) 115 (17) 110 (16) 85 (12) 55 (8) 35 (5) 21 (3) . . . WI-52 . . . 195 (28) 175 (25) 150 (25) 90 (13) 70 (10) . . . . . . Fig. 1 Progress in the high-temperature capabilities of superalloys since the 1940s. Source: Ref 2 The development of new polycrystalline alloys continued through the 1970s, however, at a more moderate rate. Attention was concentrated instead on process development, with specific interest directed toward grain orientation and directional- solidification (DS) turbine blade and vane casting technology (Fig. 2). Fig. 2 Advances in turbine blade materials and processes since 1960. Source: Ref 4 Applied to turbine blades and vanes, the DS casting process results in the alignment of all component grain boundaries such that they are parallel to the blade/vane stacking fault axis, essentially eliminating transverse grain boundaries (Fig. 3). Because turbine blades/vanes encounter major operating stress in the direction which is near normal to the stacking fault axis, transverse grain boundaries provide relatively easy fracture paths. The elimination of these paths provides increased strain elasticity by virtue of the lower <001> elastic modulus, thereby creating opportunities for further exploitation of the nickel-base alloy potential. Fig. 3 The evolution of the processing of nickel- base superalloy turbine blades. (a) From left, equiaxed, directionally solidified, and single- crystal blades. (b) An exposed view of the internal cooling passages of an aircraft turbine blade. Source: Ref 5 The logical progression to grain-boundary reduction is the total elimination thereof. Thus, single-crystal turbine blade/vane casting technology soon developed, providing further opportunity for nickel-base alloy design innovation. The late 1970s and the 1980s have, therefore, been a productive development period for nickel-base alloys designed specifically for directionally solidified columnar-grain and single-crystal cast components. These new process technologies, which are more fully discussed in the article "Directionally Solidified and Single-Crystal Superalloys" in this Volume, have contributed to dramatic improvements in gas turbine engine operating efficiency. References 1. R.W. Fawley, Superalloy Progress, in The Superalloys, C.T. Sims and W.C. Hagel, Ed., John Wiley & Sons, 1972, p 12 2. R.F. Decker, Superalloys Does Life Renew at 50?, in Proceedings of the Fourth International Symposium on Superalloys, American Society for Metals, 1980, p 2 3. Appendix B: Superalloy Data, in Superalloys II, C.T. Sims, N.S. Stoloff, and W.C. Hagel, Ed., John Wiley & Sons, 1987, p 575-597 4. M. Gell and D.N. Duhl, The Development of Single-Crystal Superalloy Turbine Blades, in Advanced High- Temperature Alloys: Processing and Properties, American Society for Metals, 1986, p 41-49 5. L.E. Dardi, R.P. Dalal, and C. Yaker, Metallurgical Advancements in Investment Casting Technology, in Advanced High-Temperature Alloys: Processing and Properties, American Society for Metals, 1986, p 25-39 Polycrystalline Cast Superalloys Gary L. Erickson, Cannon-Muskegon Corporation Superalloy Design Nickel-base superalloys have microstructures consisting of an austenitic face-centered cubic (fcc) matrix (γ) dispersed intermetallic fcc γ' Ni 3 (Al,Ti) precipitates coherent with the matrix (0 to 0.5% lattice mismatch), and carbides, borides, and other phases distributed throughout the matrix and along the grain boundaries. These complex alloys generally contain more than ten different alloying constituents. Various combinations of carbon, boron, zirconium, hafnium, cobalt, chromium, aluminum, titanium, vanadium, molybdenum, tungsten, niobium, tantalum, and rhenium result in the commercial alloys used in today's gas turbine engines. Some alloying elements have single-function importance, whereas others provide multiple functions. For example, chromium is primarily added to nickel-base alloys for sulfidation resistance (Cr 2 O 3 protective-scale formation), whereas aluminum not only is a strong γ' former but also helps provide oxidation resistance when present in sufficient quantity by forming a protective Al 2 O 3 scale. Many of the other alloying elements also have multiple roles. Titanium, while primarily partitioning to the γ', also participates in the formation of primary (MC) carbides, the hexagonal close-packed (hcp) eta (η) phase, and undesirable nitride and carbosulfide formation. Molybdenum, tungsten, tantalum, rhenium, cobalt, and chromium additions promote solid-solution strengthening, but it is known that tantalum, tungsten, and rhenium may also partition to the γ' to varying degrees and that tantalum and rhenium may also be beneficial to environmental resistance properties. Vanadium is a γ' partitioner, but it also promotes the formation of M 3 B 2 -type borides. Niobium forms the intermetallic phases delta (δ) (orthorhombic Ni 3 Nb) and γ'' (body-centered tetragonal Ni 3 Nb), but it is also involved in the formation of Laves (Fe,Ni 2 Nb) phase, carbides, borides, and/or nitrides. Hafnium is a strong carbide former that is added to polycrystalline alloys to improve grain-boundary ductility. However, at the same time, it increases the volume fraction of γ/γ' eutectic and increases oxidation resistance. Carbon, boron, and zirconium are used at varying levels for grain- boundary strengthening. All of these constituents interact in various ways to provide high tensile, creep, and fatigue strengths, plus oxidation and sulfidation resistance. Proper control of the cast microstructure and subsequent solutioning and aging treatments generally result in satisfactory component performance. Under the extreme temperature/stress conditions in which superalloy components operate, however, microstructural features change, often with attendant property changes. The microstructural instabilities that may occur include: • Intermetallic phase precipitation (σ, μ, Laves) • Phasial decomposition (carbides, borides, nitrides) • Phase coalescence and coarsening (γ') • Phasial solutioning and reprecipitation (γ') • Order-disorder transition • Material oxidation • Stress-corrosion cracking The formation of topologically close-packed phases (σ, μ, and so on) generally decreases rupture properties. Their occurrence is controlled through chemistry adjustment and is fairly predictable through use of the commonly accepted methods of calculating the so-called electron vacancy number, N v , of the given alloys. Different calculation methods exist; however, all provide a useful key to the prediction of σ formation when proper reference points are known. Although it occurs during both solidification and heat treatment, carbide precipitation is generally promoted during component heat treatment to effect an optimum grain-boundary carbide morphology and population. Discrete, blocky M 23 C 6 particles distributed in a discontinuous fashion are preferred. High-temperature, stressed exposure tends to cause carbide degeneration, often resulting in grain-boundary overload and compromised rupture strength. MC-type carbides generally occur during alloy solidification. They are titanium-rich (MC-1) or tantalum-rich (MC-2) and may partially degenerate with high-temperature exposure to form hafnium-rich (MC-3) carbides and/or M 23 C 6 , M 7 C 3 , and M 6 C carbides (secondary carbides); the specific type depends upon alloy chemistry and exposure temperature. The chromium-rich M 23 C 6 generally forms at the grain boundaries in polycrystalline materials; when present as discrete, discontinuous particles, it provides the grain-boundary strength and resistance to fracture needed to prolong service life. On the other hand, carbide degeneration also releases titanium and tantalum to the solid-solution matrix, resulting in further matrix saturation. Oversaturation can result in the formation of undesirable secondary phases such as (tungsten- and/or molybdenum-rich), α-W, α-Cr, and or M 6 C carbides, making chemistry balancing and controlled thermal treatment necessary for ultimate success. Superalloys are, indeed, complex. However, careful alloy design and processing will provide the desired results. Simply stated, superalloy property attainment is principally a function of the amount and morphology of the γ', grain size and shape, and carbide distribution. Early superalloys contained less than 25 vol% γ'. However, commercial vacuum induction refining and casting provided the opportunity for greater γ' volume fraction, to the extent that today's commercial superalloys generally contain approximately 60 vol% γ'. This increased level of γ' results in greater alloy creep strength (Fig. 4), but it can be fully exploited only in single-crystal components, where full γ' solutioning is generally possible. For polycrystalline superalloy components, high-temperature strength is affected by the condition of the grain boundaries and, in particular, the grain-boundary carbide morphology and distribution. Optimized properties can be achieved if solutioning and aging treatments are developed to attain discrete, globular carbide formation along the grain boundaries in conjunction with the optimized γ' volume fraction/morphology and component grain structure. Representative stress-rupture curves for selected nickel-base superalloys are shown in Fig. 5. Table 5 also provides stress-rupture data. Fig. 4 The relationship between γ' volume percent and stress-rupture strength for nickel- base superalloys. Source: Ref 6 Fig. 5 Stress-rupture curves for selected superalloys. (a) and (b) Nickel-base superalloys. 1000 h. (c) Cobalt- base superalloys. 1000 h. Source: Ref 3. (d) Larson-Miller stress-rupture curves for selected nickel- base superalloys. Source: Ref 7. (e) Larson-Miller stress-rupture curves for selected cobalt-base superalloys. Source: Ref 8 Cobalt-base alloys (see Table 2) are designed around a cobalt-chromium matrix with chromium contents ranging from 18 to 35 wt%. The high chromium content contributes to oxidation and sulfidation resistance, but also participates in carbide formation (Cr 7 C 3 and M 23 C 6 ) and solid-solution strengthening. Carbon content generally ranges from 0.25 to 1.0%, with nitrogen occasionally substituting for carbon. Cobalt-base alloys are often designed with significant levels of both nickel and tungsten. The addition of nickel helps to stabilize the desired fcc matrix, while tungsten provides solid-solution strengthening and promotes carbide formation. Molybdenum also contributes to solid-solution strengthening but is less effective and potentially more deleterious than a tungsten addition. Other alloying elements contributing to the solid solution and/or carbide formation are tantalum, niobium, zirconium, vanadium, and titanium. These additions provide strength by means of solid-solution and second-phase strengthening. No intermetallic precipitated phase has been discovered to equal the benefit imparted by γ' nickel-base superalloys. Solid-solution strengthening results principally from the chromium, tantalum, niobium, and tungsten additions, while second-phase strengthening is obtained primarily from the carbides and carbonitrides formed with chromium. The multiple-composition complex carbides may be present as MC, M 6 C, M 7 C 3 , M 23 C 6 , and M 2 C 3 . As with nickel-base superalloys, carbides must be precipitated at grain boundaries to control gross grain-boundary sliding and migration. Optimum mechanical properties are obtained through the careful balancing of the carbides at grain boundaries and within the matrix. When carbides are present in sufficient quantity, the skeletal carbide network that results can contribute to component strength much like the strengthening that is achieved in a composite (Ref 9). Most cobalt-base alloy aerospace castings are not heat treated apart from a relatively low-temperature stress-relief anneal. Carbide distribution is, therefore, determined during solidification, thereby high-lighting the need for stringent control of the alloy pouring temperature and cooling rate during and after solidification. Exceptions to this may be found in medical applications, where cobalt-base alloy castings for orthopedic implants are sometimes solution treated. Representative stress-rupture curves for selected cobalt-base alloys are shown in Fig. 5(c) and 5(e). Table 6 also provides stress-rupture data. References cited in this section 3. Appendix B: Superalloy Data, in Superalloys II, C.T. Sims, N.S. Stoloff, and W.C. Hagel, Ed., John Wiley & Sons, 1987, p 575-597 6. R.F. Decker, "Strengthening Mechanisms in Nickel- Base Superalloys," Paper presented at the Steel Strengthening Mechanisms Symposium, Zurich, 1969 7. W. Betteridge, Nickel and Its Alloys, Ellis Horwood, 1984 8. W. Betteridge, Cobalt and Its Alloys, Ellis Horwood, 1982 9. M.J. Donachie, Introduction to Superalloys, in Superalloys Source Book, American Society for Metals, 1984, p 9 Polycrystalline Cast Superalloys Gary L. Erickson, Cannon-Muskegon Corporation Vacuum Induction Melting of Superalloys Commercial vacuum induction melting was developed in the early 1950s, having been stimulated by the need to produce superalloys containing reactive elements within an evacuated atmosphere. The process is relatively flexible, featuring the independent control of time, temperature, pressure, and mass transport through melt stirring. As such, VIM offers more control over alloy composition and homogeneity than all other vacuum melting processes. The primary purification reaction occurring in the process is the removal of melt-contained oxygen by means of a reaction with carbon to form carbon monoxide (CO). The reaction occurs most readily at or near the melt surface with the reaction kinetics being affected by crucible geometry and melt stirring. The removal of oxygen from the melt as CO is favored by decreased melt chamber pressure, elevated bath temperature, and increased carbon activity (Ref 10). The melting crucible material is not inert and is actually another source of oxygen and other impurities, depending on refractory type and condition. Therefore, both melt refining temperature and refining duration are carefully scrutinized. Proper melt stirring is integral to the deoxidation process and must be optimized through proper furnace power frequency and application procedure to prevent refractory lining erosion, a potential problem particularly during the controlled but more vigorous CO boiling portion of the process. Vacuum induction melting deoxidation, that is, the generation of CO gas, proceeds as CO bubbles are nucleated heterogeneously along the walls and, sometimes, bottom of the melt/lining-refractory interface. This occurs preferentially at small crevices existing in the lining, with the bubbles growing during movement toward the molten metal/vacuum interface (Ref 11, 12, 13). Actual bubble formation is dependent on the number of gas molecules present; the pressure in the liquid at the level of the bubble; the temperature of the gas; and, for very small bubbles, the interfacial tension between the gas and the liquid metal. Figure 6 shows bubble formation during the VIM process. Fig. 6 Vacuum induction refining process Following formation, bubble growth and mass transport within the liquid toward the liquid/vacuum interface is dependent on: • The quantity of the dissolved gas • The decreased pressure exerted on the bubble as it rises in the melt • The bath temperature • The time it takes for the bubble to rise through the melt to the surface, which, in turn, is a function of melt stirring • The pressure above the melt • The interfacial tension between the bubble and the liquid metal The relatively vigorous, but controlled, portion of the boiling process results in the greatest CO removal. Concurrently, a slight nitrogen loss is realized because of scavenging associated with the CO bubbles, and a slight sulfur reduction may occur during the CO supersaturation stage via sulfur dioxide (SO 2 ) evolution. Minor tramp elements such as lead, silver, bismuth, selenium, and tellurium, which are deleterious to alloy elevated-temperature rupture strength and ductility (Ref 14, 15), are partially evaporated during this period as well as throughout the entire refining process (Ref 16, 17). Some undesirable elements, however, such as arsenic and tin, must be controlled through raw material selection because they are not removed by vacuum refining. Figure 7 shows the effects of VIM time and temperature on tramp element concentration. Once the boiling subsides, surface desorption of additional CO occurs, and it is during this nonboiling period that nitrogen removal (desorption) is most effective (Ref 18). Fig. 7 Evaporation of elements from an 80Ni-20Cr alloy during VIM Refractory Materials. Superalloy melting is generally undertaken in a relatively unreactive, high-bond strength, high- purity MgO-Al 2 O 3 spinel refractory lining. Refractories may be monolithic or brick and mortar, with the former providing the greater potential for alloy quality. By minimizing extremes in thermal cycling; optimizing alloy sequencing; and refining temperature, time, and pressure, alloys practically void of any lining-related nonmetallics can be produced. The types of raw materials and the melt procedure vary depending on the quality of alloy being produced. Alloys destined for critical application, the components of which may be difficult to cast, are produced using the highest-quality raw materials commercially available, in conjunction with sophisticated melt processing. Lower-quality raw materials, such as GMR-235 or IN-713 C, are used for commercial application, such as nonrotating, noncritical, or nonaerospace use, for example, turbocharger wheels, because a higher-quality alloy product is not necessary. The more sophisticated alloy systems are generally produced to premium-quality or integral wheel quality levels, with special attention given not only to cleanliness but also to specific microstructural characteristics and/or chemistry requirements to assist castability, weldability, and component mechanical properties. The Melt Process. Base charge materials are layered in the relatively warm furnace, in a manner which recognizes and accommodates the elemental melting point of the material and bridging tendency. Only those materials with oxides that are relatively easily reduced for the encountered melt conditions are placed in the initial furnace charge along with a small, controlled carbon addition. Also, those elements that have a particularly strong affinity for nitrogen may be withheld from the base charge because they lower the activity of the dissolved nitrogen. Following furnace evacuation and particular heat-up cycles that ensure proper closure of any refractory lining cracks prior to metal liquation, optimum temperature and vacuum pressure, consistent with promoting a somewhat vigorous CO boil, is attained. Bath refining is undertaken at a temperature and duration long enough to reach the so-called system equilibrium conditions, the assurance of which is provided by the attainment of consistent furnace leak-up rates. At this point, those elements that were held from the base charge because of their relative reactivity toward oxygen, for example, aluminum, titanium, zirconium, and hafnium, are added with an associated solutioning and homogenization procedure. Dip sample alloy chemistry is checked in a relatively short time period with any necessary chemistry adjustment undertaken. A similar analytical check is undertaken prior to pouring. Pouring proceeds once the correct chemistry is ensured, the bath is properly solutioned/homogenized, and the proper pour temperature is attained. It is generally undertaken under high vacuum condition and proceeds from the furnace crucible into a relatively sophisticated multicompartment tundish, thereby ensuring that extensive time for flotation is achieved [...]... ppm 20+ 6 0-1 00 60 5-1 5 1 0-2 5 1 O, ppm 5+ 5-1 0 . γ' volume percent and stress-rupture strength for nickel- base superalloys. Source: Ref 6 Fig. 5 Stress-rupture curves for selected superalloys. (a) and (b) Nickel-base superalloys 15+ 1 0-4 0 1 0-3 0 5-1 5 10 <5 Zr, wt% . . . <0 .01 0. 001 . . . <10 ppm . . . Fe, wt% 0.10+ . . . . . . 0.0 5-0 .10 . . . 0.03 Cu, wt% 0.002+ 0.08 0 .0 1- 0.05 0.00 2-0 .005 <0. 001 <0. 001. MC-type carbides generally occur during alloy solidification. They are titanium-rich (MC-1) or tantalum-rich (MC-2) and may partially degenerate with high-temperature exposure to form hafnium-rich