SINGLE CRYSTAL NICKEL-BASED SUPERALLOY

Abstract: This invention relates to a single crystal casting to be used under high stress, high temperature conditions up to about 2030 F characterized by an increased resistance to creep under such conditions. The casting is made from a nickel-based superalloy consisting essentially of the following elements in percent by weight: from 6.2 to 6.8 percent rhenium, from 1.8 to 2.5 percent chromium, from 1.5 to 2.5 percent cobalt, from 8.0 to 9.0 percent tantalum, from 3.5 to 6.0 percent tungsten, from 5.5 to 6.1 percent aluminum, from 0.1 to 0.5 percent titanium, from 0.01 to 0.1 percent columbium, from 0.25 to 0.60 percent molybdenum, from 0 to 0.05 percent hafnium, from 0 to 0.04 percent carbon, from 0 to 0.01 percent boron, from 0 to 0.01 percent yttrium, from 0 to 0.01 percent cerium, from 0 to 0.01 percent lanthanum, from 0 to 0.04 percent manganese, from 0 to 0.05 percent silicon, from 0 to 0.01 percent zirconium, from 0 to 0.001 percent sulfur, from 0 to 0.10 percent vanadium, and the balance nickel + incidental impurities. The superalloy has a phasial stability number NV3B less than about 1.65.

Description: SINGLE CRYSTAL NICKEL-BASED SUPERALLOY BACKGROUND OF THE INVENTION 1. Field of the Invention This invention relates to single crystal nickel-based superalloys and, more particularly, single crystal nickel-based superalloys and articles made therefrom for use in advanced gas turbine engines under high stress, high temperature conditions. 2. Description of the Prior Art Advances over recent years in the metal temperature and stress capability of single crystal articles have been the result of the continuing development of single crystal superalloys, as well as improvements in casting processes and engine application technology. These single crystal superalloy articles include rotating and stationary turbine blades and vanes found in the hot sections of gas turbine engines. However, gas turbine engine design goals have remained the same during the past decades. These goals include the desire to increase engine operating temperature, rotational speed, thrust-to-weight ratio, fuel efficiency, and engine component durability and reliability. The basic technology of alloys for the casting of single crystal components is described in U.S. Patent Nos. 3,494,709; 4,116,723 and 4,209,348. Development work resulted in first generation nickel-based superalloys, which were materially improved over those described in the aforementioned patents. However, these first generation nickel-based superalloys contained no rhenium. Examples of such first generation nickel-based superalloys, commercially known as CMSX-2 alloy and CMSX-3 alloy produced by Cannon-Muskegon Corporation, assignee of the present application, are described in U.S. Patent No. 4,582,548. Further development work resulted in second generation nickel-based superalloys havin improved creep strength/creep rate. These second generation nickel-based superalloys have a moderate rhenium content of about 3 weight percent. An example of such a second generation nickel- based superalloy is described in U.S. Patent No. 4,643,782. This patent discloses a superalloy, commercially known as CMSX-4 alloy, having a specific nickel-based composition including a rhenium content in the range of 2.8-3.2 weight percent. The present invention provides the next generation of nickel-based superalloys having higher total refractory element ( +Re+Mo+Ta) content and improved mechanical properties. Single crystal articles are generally produced having the low-modulus (001) crystallographic orientation parallel to the component dendritic growth pattern or blade stacking axis. Face- centered cubic (FCC) superalloy single crystals grown in the (001) direction provide extremely good thermal fatigue resistance relative to conventionally cast articles. Since these single crystal articles have no grain boundaries, alloy design without grain boundary strengtheners, such as carbon, boron and zirconium, is possible. As these elements are alloy melting point depressants, their reduction from an alloy design provides a greater potential for high temperature mechanical strength achievement since more complete gamma prime solution and microstructural homogenization can be achieved relative to directionally solidified (DS) columnar grain and conventionally cast materials. Their reduction also makes possible a higher incipient melting temperature. These process benefits are not necessarily realized unless a multi-faceted alloy design approach is undertaken. Alloys must be designed to avoid tendency for casting defect formation such as freckles, slivers, spurious grains and recrystallization. Additionally, the alloys must provide an adequate heat treatment window (numeric difference between an alloy' s gamma prime solvus and incipient melting point) to allow for nearly complete gamma prime solutioning. At the same time, the alloy compositional balance should be designed to provide an adequate blend of engineering properties necessary for operation in gas turbine engines. Selected properties generally considered important by gas turbine engine designers include: elevated temperature creep- rupture strength, thermo-mechanical fatigue resistance, impact resistance plus hot corrosion and oxidation resistance. An alloy designer can attempt to improve one or two of these design properties by adjusting the compositional balance of known superalloys. However, it is extremely difficult to improve more than one or two of the design properties without significantly or even severely compromising the remaining properties. The unique superalloy of the present invention provides an excellent blend of the properties necessary for use in producing single crystal articles for operation in gas turbine engine hot sections. SUMMARY OF THE INVENTION This invention relates to a nickel-based superalloy comprising the following elements in percent by weight : from about 5.0 to about 7.0 percent rhenium, from about 1.8 to about 4.0 percent chromium, from about 1.5 to about 9.0 percent cobalt, from about 7.0 to about 10.0 percent tantalum, from about 3.5 to about 7.5 percent tungsten, from about 5.0 to about 7.0 percent aluminum, from about 0.1 to about 1.2 percent titanium, from about 0 to about 0.5 percent columbium, from about 0.25 to about 2.0 percent molybdenum, from about 0 to about 0.15 percent hafnium, and the balance nickel plus incidental impurities, the superalloy having a phasial stability number N V3B less than about 2.10. Advantageously, this superalloy composition may be further comprised of (percentages are in weight percent) from about 0 to about 0.04 percent carbon, from about 0 to about 0.01 percent boron, from about 0 to about 0.01 percent yttrium, from about 0 to about 0.01 percent cerium and from about 0 to about 0.01 percent lanthanum. Although incidental impurities should be kept to the least amount possible, the superalloy can also be comprised of from about 0 to about 0.04 percent manganese, from about 0 to about 0.05 percent silicon, from about 0 to about 0.01 percent zirconium, from about 0 to about 0.001 percent sulfur, and from about 0 to about 0.10 percent vanadium. In all cases, the base element is nickel. Furthermore, this superalloy can advantageously have a phasial stability number N V3B less than about 1.85, and a chromium content of from about 1.8 to about 3.0 percent, a rhenium content of from about 5.5 to about 6.5 percent, and a cobalt content of from about 2.0 to about 5.0 percent. This invention provides a superalloy having an increased resistance to creep under high stress, high temperature conditions, particularly up to about 2030°F. In one preferred embodiment, this invention relates to a single crystal casting to be used under high stress, high temperature conditions up to about 2030°F characterized by an increased resistance to creep under such conditions . In this embodiment, the casting is made from a nickel-based superalloy consisting essentially of the following elements in percent by weight: from 6.2 to 6.8 percent rhenium, from 1.8 to 2.5 percent chromium, from 1.5 to 2.5 percent cobalt, from 8.0 to 9.0 percent tantalum, from 3.5 to 6.0 percent tungsten, from 5.5 to 6.1 percent aluminum, from 0.1 to 0.5 percent titanium, from 0.01 to 0.1 percent columbium, from 0.25 to 0.60 percent molybdenum, from 0 to 0.05 percent hafnium, from 0 to 0.04 percent carbon, from 0 to 0.01 percent boron, from 0 to 0.01 percent yttrium, from 0 to 0.01 percent cerium, from 0 to 0.01 percent lanthanum, from 0 to 0.04 percent manganese, from 0 to 0.05 percent silicon, from 0 to 0.01 percent zirconium, from 0 to 0.001 percent sulfur, from 0 to 0.10 percent vanadium, and the balance nickel + incidental impurities, wherein the superalloy has a phasial stability number N V3B less than about 1.65. Single crystal articles can be suitably made from the superalloy of this invention. The article can be a component for a turbine engine and, more particularly, the component can be a gas turbine blade or gas turbine vane. The superalloy compositions of this invention have a critically balanced alloy chemistry which results in a unique blend of desirable properties. These properties include: excellent single crystal component castability, particularly for moderately sized blade and vane components; adequate cast component solutionability; excellent resistance to single crystal cast component recrystallization; ultra-high creep-rupture strength to about 2030°F; extremely good smooth and notched low cycle fatigue strength; extremely good high cycle fatigue strength; high impact strength; very good bare hot corrosion resistance; very good bare oxidation resistance; adequate coatability; and adequate microstructural stability, such as resistance to the undesirable, brittle phases called topologically close-packed (TCP) phases. Accordingly, it is an object of the present invention to provide superalloy compositions and single crystal articles made therefrom having a unique blend of desirable properties. It is a further object of the present invention to provide superalloys and single crystal articles made therefrom for use in advanced gas turbine engines under high stress, high temperature conditions, such as up to about 2030°F. These and other objects and advantages of the present invention will be apparent to those skilled in the art upon reference to the following description of the preferred embodiments. BRIEF DESCRIPTION OF THE DRAWINGS FIG. 1 is a chart of hot corrosion test results performed to 117 hours on one embodiment of the alloy of this invention and on two prior art alloys; FIG. 2 is a chart of hot corrosion test results performed to 144 hours on another embodiment of the alloy of this invention and on a prior art alloy; FIG. 3 is a graphical comparison of bare alloy oxidation data from tests performed at 2012°F on two embodiments of this invention and on three other alloys; FIG. 4 is a graphical comparison of bare alloy oxidation data from tests performed at 1886°F on two embodiments of this invention and on three other alloys; FIG. 5 is a graphical comparison of bare alloy corrosion data from tests performed at 1742°F on two embodiments of this invention and on two other alloys; and FIG. 6 is a graphical comparison of bare alloy corrosion data from tests performed at 1742°F on four embodiments of this invention and on two other alloys. DESCRIPTION OF THE PREFERRED EMBODIMENTS The nickel-based superalloy of the present inventio comprises the following elements in percent by weight : Rhenium about 5.0-7.0 Chromium about 1.8-4.0 Cobalt about 1.5-9.0 Tantalum about 7.0-10.0 Tungsten about 3.5-7.5 Aluminum about 5.0-7.0 Titanium about 0.1-1.2 Columbium about 0-0.5 Molybdenum about 0.25-2.0 Hafnium about 0-0.15 Nickel + Incidental balance Impurities This superalloy composition also has a phasial stability number N V3B less than about 2.10. Further, this invention has a critically balanced alloy chemistry which results in a unique blend of desirable properties. These properties include increased creep- rupture strength relative to prior art single crystal superalloys, single crystal component castability, cast component solutionability, single crystal component resistance to recrystallization, fatigue strength, impact strength, bare hot corrosion resistance, bare oxidation resistance, component coatability, and microstructural stability, including resistance to TCP phase formation under high stress, high temperature conditions. Unlike prior nickel-based superalloys known in the art, the superalloys of the present invention have a low chromium, low cobalt and high rhenium content. The chromium is about 1.8-4.0% by weight. Advantageously, the chromium content is from 1.8% to 3.0% by weight. This chromium content is significantly lower than that typically found in prior art single crystal nickel-based superalloys. In the present superalloy, chromium provides hot corrosion resistance, although it may also assist with the alloy's oxidation capability. Tantalum and rhenium also assist toward hot corrosion property attainment, and aluminum is present at sufficient levels to provide adequate oxidation resistance, so that relatively low addition of chromium is tolerable in this alloy. Besides lowering the alloy's gamma prime solvus, chromium contributes to the formation of Cr, Re, W-rich TCP phase and must be balanced accordingly in these compositions. The cobalt content is about 1.5-9.0% by weight. Advantageously, the cobalt content is from 2.0% to 5.0% by weight. This cobalt content is lower than that typically found in prior art single crystal nickel-based superalloys. In the present superalloy, cobalt assists in providing an appropriate heat treatment window since it has the effect of lowering the alloy's gamma prime solvus while generally not affecting its incipient melting point. Rhenium-containing alloys are generally designed with much higher cobalt content than the present invention for the purpose of imparting increased solid solubility and phasial stability. However, the superalloys of the present invention unexpectedly show that much lower cobalt contents are possible and desirable toward providing optimized phasial stability, including control of TCP phase formation. The rhenium content is about 5.0-7.0% by weight and, advantageously, rhenium is present in an amount of from 5.5% to 6.5% by weight. The amount of rhenium in the superalloy of the present invention is significantly greater than the rhenium content of prior art single crystal nickel-based superalloys. Furthermore, the superalloys of this invention are generally designed with an increased level of refractory element content, e.g., W+Re+Mo+Ta. The tungsten content is about 3.5-7.5% by weight and, advantageously, the amount of tungsten is from 3.5% to 6.5% by weight. Tungsten is added since it is an effective solid solution strengthener and it contributes to strengthening the gamma prime. Additionally, tungsten is effective in raising the alloy's incipient melting temperature. The amount of tungsten added to these superalloys is balanced with the amount of rhenium added since they both contribute to the formation of "freckle" defects during the single crystal investment casting process. They also both strongly effect the propensity for TCP phase formation. Similar to tungsten, rhenium is effective in raising the alloy's incipient melting point. However, rhenium is a more effective strengthener than tungsten, molybdenum and tantalum in terms of elevated temperature creep-rupture and, therefore, rhenium is added appropriately. Additionally, rhenium has a positive influence on this alloy's hot corrosion resistance. Moreover, rhenium partitions primarily to the gamma matrix, and it is effective in slowing gamma prime particle growth during high temperature, high stress conditions. Besides requiring the balancing of rhenium with tungsten for castability reasons, W+Re must also be set at a level consistent with minimizing TCP phase formation. In general, the TCP phases which occur in such material are rich in chromium, tungsten, and rhenium content, with rhenium being present in the greatest proportion. Thus, careful Re/W ratio control is necessary in this alloy to control the propensity for TCP phase formation. The molybdenum content is about 0.25-2.0% by weight. Advantageously, molybdenum is present in an amount of from 0.25% to 1.5% by weight. Molybdenum is a good solid solution strengthener, but it is not as effective as tungsten, rhenium and tantalum. However, since the alloy's density is always a design consideration, and the molybdenum atom is lighter than the other solid solution strengtheners, the addition of molybdenum is a means of assisting control of the overall alloy density in the compositions of this invention. The tantalum content is about 7.0-10.0% by weight and, advantageously, the tantalum content is from 8.0% to 10.0% by weight. Tantalum is a significant contributor to this alloy's strength through means of solid solution strengthening and enhancement of gamma prime particle strength (tantalum also partitions to the gamma prime phase) . In this alloy, tantalum is able to be utilized at relatively high concentration since it does not contribute to TCP phase formation. Additionally, tantalum is an attractive single crystal alloy additive in this composition since it assists in preventing "freckle" defect formation during the single crystal casting process. Tantalum is also beneficial in this composition since it tends to raise this alloy's gamma prime solvus, and it is effective toward promoting good alloy oxidation and hot corrosion resistance, along with aluminide coating durability. The aluminum content is about 5.0-7.0% by weight. Furthermore, the amount of aluminum present in this composition is advantageously from 5.3% to 6.5% by weight. Aluminum and titanium are the primary elements comprising the gamma prime phase. These elements are added in this composition in a proportion and ratio consistent with achieving adequate alloy castability, solution heat treatability, phasial stability and high mechanical strength. Aluminum is also added to this alloy in proportions sufficient to provide oxidation resistance. The titanium content is about 0.1-1.2% by weight. Advantageously, titanium is present in this composition in an amount from 0.2% to 0.8% by weight. Titanium is generally beneficial to the alloy's hot corrosion resistance, but it can have a negative effect to oxidation resistance, alloy castability and alloy response to solution heat treatment. Accordingly, the titanium content must be maintained within the stated range of this composition. The columbium content is about 0-0.5% by weight and, advantageously, the columbium content is from 0 to 0.3% by weight. Columbium is a gamma prime forming element and it is an effective strengthener in the nickel-based superalloys of this invention. Generally, however, columbium is a detriment to alloy oxidation and hot corrosion properties, so its addition to the composition of this invention is minimized. Moreover, columbium is added to this invention's composition for the purpose of gettering carbon, which can be chemi-sorbed into component surfaces during non-optimized vacuum solution heat treatment procedures. Any carbon pick-up will tend to form columbium carbide instead of titanium or tantalum carbide, thereby preserving the greatest proportion of titanium and/or tantalum for gamma prime and/or solid solution strengthening in this alloy. The hafnium content is about 0-0.15% by weight and, advantageously, hafnium is present in an amount from 0.02 to 0.05% by weight . Hafnium is added in a small proportion to the present composition in order to assist with coating adherence. Hafnium generally partitions to the gamma prime phase. The balance of this invention's superalloy composition is comprised of nickel and small amounts of incidental impurities . Generally, these incidental impurities are entrained from the industrial process of production, and they should be kept to the least amount possible in the composition so that they do not affect the advantageous aspects of the superalloy. For example, these incidental impurities may include up to about 0.04% by weight manganese, up to about 0.05% by weight silicon, up to about 0.01% by weight zirconium, up to about 0.001% by weight sulfur, and up to about 0.10% by weight vanadium. Amounts of these impurities which exceed the stated amounts could have an adverse effect upon the resulting alloy's properties. Additionally, the superalloy may optionally contain about 0-0.04% by weight carbon, about 0-0.01% by weight boron, about 0- 0.01% by weight yttrium, about 0-0.01% by weight cerium and about 0-0.01% by weight lanthanum. Not only does the superalloy of this invention have a composition within the above specified ranges, but it also has a phasial stability number N V3B less than about 2.10. Advantageously, the phasial stability number N V3B is less than 1.85 and, preferably, the phasial stability number N V3B is less than 1.65. As can be appreciated by those skilled in the art, N V3B is defined by the PW N-35 method of nickel-based alloy electron vacancy TCP phas control factor calculation. This calculation is as follows: EQUATION 1 Conversion for weight percent to atomic percent : Atomic percent of element i = Pi = Wi/A1 X100 ∑i ( i/Ai) where : Wi = weight percent of element i Ai = atomic weight of element i EQUATION 2 Calculation for the amount of each element . present in the continuous matrix phase: Element Atomic amount Rii remaining Cr R Cr =0.97P Cr -0.375P B -1.75P c Ni . 525P B -3 (P„ι + 0 . 03P„+P tl - 0 .5P-+0 .5P V +P T ,+P cb +P H[ ) Ti, Al, B, Ri=0 C, Ta, Cb, Hf V Rv=0.5P v W * R ( „ =P-,,-Q.l67P r . P w +P W * Note : weight percentage Re is added to weight percentage W for the calculation above. EQUATION 3 Calculation of N v3B using atomic factors from Equations 1 and 2 above : where: i = each individual element in turn. N ± i = the atomic factor of each element in matrix. (N v ) i = the electron vacancy No. of each respective element. This calculation is exemplified in detail in a technical paper entitled "PHACOMP Revisited", by H. J. Murphy, C. T. Sims and A. M. Beltran, published in Volume 1 of International Symposium on Structural Stability in Superalloys (1968) , the disclosure which is incorporated by reference herein. As can be appreciated by those skilled in the art, the phasial stability number for th superalloys of this invention is critical and must be less than th stated maximum to provide a stable microstructure and capabilit for the desired properties under high temperature, high stres conditions. The phasial stability number can be determine empirically, once the practitioner skilled in the art is i possession of the present subject matter. The superalloy of this invention can be used to suitabl make single crystal articles, such as components for turbin engines. Preferably, this superalloy is utilized to make a singl crystal casting to be used under high stress, high temperatur conditions characterized by an increased resistance to creep unde such conditions, particularly high temperature conditions up t about 2030°F. Furthermore, it is believed that this invention ha an increased resistance to creep under high stress, hig temperature conditions at about 2125°F and above compared t similar prior art materials, such as the CMSX-4 superalloy. Whil this superalloy can be used for any purpose requiring high strengt castings incorporating a single crystal, its particular use is i the casting of single crystal blades and vanes for gas turbin engines. This alloy possesses an unusual resistance to componen recrystallization during solution heat treatment, which i considered an important alloy characteristic that is necessary whe producing advanced technology, multi-piece, cast bonded singl crystal airfoils. Additionally, this superalloy provides the allo castability characteristics believed necessary to produc conventional-process-cast, moderately-sized turbine airfoils wit intricate cooling passages. While this superalloy' s primary use is in aircraf turbine engines, there are stationary engine applications requiring the specialized high performance characteristics of this alloy. This is particularly the case in turbine engines which require performing characteristics with very restricted clearances, thereby materially limiting the amount of permissible creep. Engines designed to develop high performance characteristics are normally operated at higher component temperatures and, therefore, the problem of creep is increased. Generally, creep in excess of 1% is considered unacceptable in these cases . The creep characteristics of known state of the art alloys have limited operating temperatures and, thus, maximum performance capability. The superalloy of this invention has an increased resistance to creep under high stress, high temperature conditions, particularly up to 2030°F. The single crystal components made from this invention's compositions can be produced by any of the single crystal casting techniques known in the art. For example, single crystal directional solidification processes can be utilized, such as the seed crystal process and the choke process. The single crystal castings made from the superalloy of the present invention are advantageously subjected to a high temperature aging heat treatment in order to optimize the creep- rupture properties of these alloys. This invention's single crystal castings can be aged at a temperature of from about 1950°F to about 2125°F for about 1 to about 20 hours. Advantageously, this invention's single crystal castings can be aged at a temperature of from about 2050°F to about 2125°F for about 1 to about 20 hours. However, as can be appreciated by those skilled in the art, the optimum aging temperature and time for aging depends on the precise composition of the superalloy. This invention provides superalloy compositions having a unique blend of desirable properties. These properties include: excellent single crystal component castability, particularly for moderately sized blade and vane components; excellent cast component solutionability; excellent resistance to single crystal cast component recrystallization; ultra-high creep-rupture strength to about 2030°F; extremely good low cycle fatigue strength; extremely good high cycle fatigue strength; high impact strength; very good bare hot corrosion resistance; very good bare oxidation resistance; adequate component coatability; and microstructural stability, such as resistance to formation of the undesirable TCP phases. As noted above, this superalloy has a precise composition with only small permissible variations in any one element if the unique blend of properties is to be maintained. In order to more clearly illustrate this invention and provide a comparison with representative superalloys outside the claimed scope of the invention, the examples set forth below are presented. The following examples are included as being illustrations of the invention and its relation to other superalloys and articles, and should not be construed as limiting the scope thereof. EXAMPLES A large number of superalloy test materials were prepared to investigate the compositional variations and ranges for the superalloys of the present invention. Some of the alloy compositions tested and reported below fall outside the claimed scope of the present invention, but are included for comparative purposes to assist in the understanding of the invention. Representative alloy aim chemistries of those materials tested are reported in Table 1 below. I&BLE L Third generation single crystal alloy development to investigat the compositional variations for the superalloys of the present invention began with the definition and evaluation of a series experimental compositions. Increased creep-rupture strength wa the primary objective of the initial development effort, with elemental balancing to provide a combination of useful engineering characteristics following the definition of a base concept for increased strength. The initial materials explored the utility of higher levels of refractory element and gamma prime forming elements than are present in similar prior art compositions. As shown in Table 1 the alloy chromium content was reduced to improve alloy stability. Cobalt content, initially thought to be required fo increased solid solubility, could be signi icantly reduced. Refractory element content ( +Re+Mo+Ta) was varied, while the summation of the primary gamma prime partitioning elementb (Al+Ti+Ta+Cb) was also varied. The alloy's Re content was initially explored at conventional levels, but it was found tha the Re level had to be increased. Standard N v3B calculations were performed during the initial all design stage to assist respective alloy phasial stability predictions, with that number varying from one alloy compositi to another. -19- S ome of the alloys were produced using production-type procedures. These alloys were vacuum induction melted in the C annon-Muskegon Corporation V-1 furnace, yielding approximately 200 -300 lbs. of bar product per alloy (see Table 2 below). Q uantities of each compositional iteration, as reported in Table 2 , were made into test bars and test blades by vacuum investment casting. S olution heat treatment procedures were developed in the laboratory in 3" and 6" diameter tube furnaces. Gamma prime aging treatments were also performed in the laboratory. .TAg^E. z - H R -20- All other specimens reported in Table 1 above were pro d uced by blending base alloy bar stock with the virgin elemental additi necessary to achieve the desired composition. The blen d ing wa done during test bar and blade manufacture. The base alloy ba stock plus virgin additions were placed into the casting furna melt crucible, melted and the bath homogenized prior to pourin into an appropriate shell mold. It is believed that good correlation between alloy aim chemistry and test bar/bla d e chemistry was routinely achieved (see Table 3 below) . TABLE 3, ALLOT TEST BAR CHEMISTRIES • Mloy J-r_ Co Ho Cb TI Al Ta_ Re Hf HI Nv CHSX-10A -10B -IOC -100 -10E -10F .OH .027 -10G -1C-3B -10Gb (10K) -10H -101 -10l» -10J .017.01 -10L -12A -12B -12C -12C» (120) -12E -10R1 -12RI * PUA M-35 He-hod For the CHSX-10D specimen (see Table 1) , high quality virgin elemental additions were vacuum melted and the refined material was poured into 2 diameter bars. In turn, a quantity of the resulting bar was. used to produce single crystal test bar/blade specimens by investment casting. It was apparent that considerable variation in the investment casting process integrity may have occurred during specimen manufacture since varying levels of test bar freckle formation, secondary dendrite arm spacing and property attainment were apparent. Derivative alloy response to solution treatment (reporte * d in Table 4 below) varied, and was a function of both alloy composition and test specimen quality. Heat treatments developed for the alloy iterations are reported in Table 4 below. Full gamma prime solutioning was desired for each material, however, this objective was not universally achieved. Primary gamma prime aging was performed to effect a more desirable gamma prime particle size and distribution. Secondary gamma prime aging was performed to effect precipitatio of conventional matrix gamma prime precipitates along with ultra fine gamma prime precipitates located within the matrix channels between the primary gamma prime particles for these specimens. TAP^E 4 Mt ..Tr9«t'nW1t..P.C .U Determined by visual estimation Specimens βlr cooled from ell •oino treatments Fully heat treated test bars were creep-rupture tested. The specimens were machined and low-stress ground to ASTM standard proportional specimen dimension. The specimens were creep- rupture tested at various conditions of temperature and stress according to standard ASTM procedure. A significant factor of the CMSX-IOA alloy design was the shift to higher Re content. At the same time, W, Cr, Ta and other gamma prime etrengthenβrs were balanced to provide the desired alloy characteristics and properties. The alloys higher Re leve resulted in significantly improved creep-rupture strength throughout the entire test regime, as indicated by the results reported in Table 5 below for the CMSX-IOA specimens * TABLE 5 • As-to.utler»d Condition Microstructural review of the failed rupture specimens of this alloy revealed that TCP phase precipitation occurred during the respective creep-rupture tests, particularly those at 1900*F and above. It became apparent that the N v3B phasial stability number calculation would be an effective tool in predicting alloy stability and, effectively, high temperature creep strength for the invention. Wherein the CMSX-IOA specimen's N v3B number was 2.08, CMSX-IOB was designed to the 2.02 level. This was accomplished by the further reduction of alloy Cr content and similar reduction to Co and W+Re level. W was reduced more than the Re in this specimen since Re is more effective in the solid solution. Additionally, wherein some loss in W contribution to the gamma prime could be anticipated, it was sufficiently replaced by the modest increase to Ta content in this composition. These changes resulted in the CMSX-IOB alloy specimen exhibiting even more improved creep strength at 1800*F. Table 6 reported below illustrates that three specimens achieved an average life of 961 hours, with 1.0% creep occurring at an average of 724 hours. However, it was observed that TCP phase was present at higher temperature. 1ΔELE_£ txx-iΛ atεtv-wrtwε 21WF/15.0 k«1 69.7 54.2 48.1 69.4 47.674 25.3 36 * Aβ-solutlon-βd Condition Only about 97-98% gamma prime solutioning was achieved in the CMSX-IOA and -10B materials (see Table 4) which was insufficient for the purpose of optimizing alloy mechanical properties and microstructural homogeneity. Attainment of a greater level of gamma prime solutioning, therefore, became an equal priority in tandem with improving microstructural stability at temperatures above 1900T. To confirm the suspected composition of the TCP phase forming in the alloys, scanning electron microscope (SEM) wavelength dispersive x-ray (WDX) microchemistry analyses of CMSX-IOB test bar contained needles was undertaken and compared to the alloys gamma and gamma prime compositions. The results, reported in Table 7 below, confirm that the needles were enriched in Cr, W and Re. -IΔBI----_Z O.5X- 10B Hlero-Ch«nl »trv An vβeβ Tranβvtrβt Section, tottoa Iβr Location, tolutlo d to 2465°F A «d 1975 β F/19. C The calculated N v3B numbers were 1.90 for CMSX-10C and 1.95 for CMSX-10D. Re was maintained at around 5% while W was further reduced to improve stability in these specimens. Alloy Ta was increased since it did not participate in TCP formation and the Ta/W ratio was effectively improved, which assisted with alloy castability. Chromium was reduced in the -IOC specimens but increased to 4.0% in the -lOD specimens to provide an opportunit to determine the suitability of the Cr levels from a hot corrosion standpoint. Co was reduced in both materials, significantly in the -lOD specimen, while Al+Ti level was also reduced to assist in achieving more complete gamma prime solutioning. Creep-rupture results for the two specimens are reported below in Tables 8 and 9, respectively. Even though t -10D alloy specimens were observed to exhibit full gamma prime βolutioning (as opposed to 99.-99.5% for CMSX-IOC) the alloys greater Cr content, which necessitated a lower λl+Ti level, effected lower properties than attained with CMSX-IOC. Howeve both materials exhibited improved alloy stability and higher temperature properties, so that attempts to balance the alloys low and high temperature creep response were favorable. -30- TΛBLE 9 The acceptability of the alloys' low Cr content was confirmed through extremely aggressive short-term burner rig hot corrosio tests performed at 1650'F, 1% sulfur, 10 ppm sea salt condition FIGS. 1 and 2 illustrate the results for tests performed to 117 and 144 hours for the CMSX-10C and CMSX-10D specimens, respectively. In both cases, the materials performed similar t MAR M 247-type materials, thereby confirming the suitability of the low Cr alloy design concept. With the -above-noted results, another series of alloys, CMSX-IO -10F, -10G, -10H, -101,and -12A were designed, produced and evaluated. The alloys explored Re level ranging 4.8-6.3%, 2.2 3.0% Cr level, 4.7-7.6% Co level and the remainder balanced to maintain castability, improve solutionability and improve phasia stability. The N y3l number ranged between 1.81-1.89. One of the series, CMSX-IOF, contained .02% C and .02% B. These additions were observed to improve casting yield and may have assisted in providing more consistent control of single crystal cast article orientation. However, the melting point depressants, C and B, restricted the specimen's response to solution heat treatment. The CMSX-IOF creep-rupture properties are reported in Table 10 below. TABLE IP OT5X-107 QIEEP-Rlj jRE The CMSX-IOE, G, H and I, plus CMSX-12A creep-rupture specime results are reported below in Tables ll, 12, 13, 14, and 15, respectively. The results show a general improvement to allo creep-rupture strength above 1900*F while maintaining extreme good strength at lower temperatures. TABLE 11 CNSX-10E OtEEP-IUPTlSE TAB E 12 P«?*;1oo WtlT-BfTWE TABLE 13 X- -R TAB E 14 otgr-io i CREEP artxmε TABLE 15 Q-5X-12A CREEP-RUPTURE Varying the primary gamma prime aging treatment was explored wit most of the development activity concentrated on achieving optimized gamma prime size and distribution through longer soak times at 1975*F (see Table 4) since higher temperature aging treatments accelerated TCP phase formation during the aging cycle. Ten to twenty-one hour soak times at 1975'F were successful sinc they provided average gamma prime particles of about 0.5 urn dimension. However, it appeared that shorter primary gamma prim aging time at higher temperature may be more practical, once mor stable microεtructures were defined. Microchemical SEM WDX needle particle analyses was performed on failed CMSX-10G creep-rupture specimen. The specimen, tested at 1976T/28.1 ksi condition, exhibited needles in its microstructure. The results of the analysis are reported in Table 16 below and indicate, again, that the needles formed in this class of material are particularly rich in Re, but are als enrichened with Cr and W. -37- TABLE 16 CMSX-lQg 1976 β F/28. 1 ksi 104 . 6 HRS . ELEM A standardized test for resistance to recrystallization was performed on a CMSX-IOG test bar. The test method and the results are reported in Table 17 below. The test results indicate that the CMSX-IOG specimen exhibited similar resistanc to cast process/solution treatment/bonding process recrystallization level in comparison to CMSX-4 alloy. TABLE 17 TABLE 17 Method: A controlled level of eompreββlve βtreββ 1β itaperted on the entire surface of •s-caβt test ber. The bar Is then solution heet treated. Following solution treatment, the bar Iβ eectioned and the tra sverβe section Iβ observed ■wtallographieally. Depth of recrystallization e-eaβurementβ are taken. Evaluation Standards: Resistance To RX Anticipated In Hade Alloy Depth of RX CastInge The CMSX-lOGa -101a, -12B,-12C, -10J, -lORi and -12Ri compositions were defined and evaluated. No creep-rupture properties were generated for the CMSX-IOJ specimen, although test bars were produced and a solution heat treatment developed. Again, the inclusion of C and B in the -10J composition appeared to have positive effect to single crystal test specimen yield. Additionally, the lower levels of C and B than evaluated in CMSX IOF specimen, particularly lower B, made the material more amenable to solution heat treatment. Ninety-eight to ninety-nin percent gamma prime solutioning was achieved, as opposed to the approximate 95% level typical of the CMSX-10F composition. The CMSX-lOGa and -101a alloys were designed with N v3B numbers of about 1.70. These alloy specimens contain about 2.5% Cr, 3.3- 4.0% Co, 5.6-6.0% Re, greater Ta/W ratio, reduced Cb, and reduce Al+Ti content. Such reduction to Cb+Al+Ti level improved the solutioning characteristics of the materials (see Table 4) , plus assisted achievement of increased alloy stability. Both specimens exhibited nearly full gamma prime solutioning. The lowered N γ3l number continued to show e fectiveness in providing better creep-rupture capability at temperature greater than 1900*F, while maintaining extremely good creep-strength at lower temperature. CMSX-lOGa test results from specimens produced with improved casting process controls exhibited 700 hours or more life with about 475 hours required to creep to 1.0% for 1800'F/36.0 ksi condition. For higher temperature exposure, the specimen provided the improved average life of about 500 hours at 2050*F/15.0 ksi condition and average 1.0% creep deformation that occurred at about 250 hours, as indicated by the results reported in Table 18 below. TABLE 18 C -a-IQC CTKF-WfTWe TIME II MUMS ■UPTUtE TIME X s π-uL at-----* KADIMG TO REAOI τ«τ m-cπ-α- "CUBS gt-OMO. ---- t. hcnrm X defw— t.on 1.9* i. 1800*r 36.0 kit 300.7 19.9 25.2 499.7 19.541 316.5 360 384.2 29.1 25.4 583.9 26.395 370.0 401 505.1 22.6 29.8 508.7 18.212 307.4 347 730.9 42.0 42.8 730.7 40.216 477.6 516 460.6 428.5 1850/36.0 184.5 41.0 33.9 183.2 37.154 82.3 94 291.5 27.3 29.9 290.2 19.323 191.6 207 279.5 33.9 32.3 278.1 29.054 155.3 180 323.9 30.9 36.6 322.9 29.218 194.1 217 326.5* 8.9 12.6 295.2* 33.3 33.5 174.1 162.3 300.1* 22.8 22.4 1976/28.1 88.6 34.9 33.9 39.7 48 100.1 28.2 29.2 53.9 61 107.9 28.8 31.4 51.1 62 87.1 27.0 33.8 39.2 48 82.8 23.3 27.7 20.6 38 88.2 31.2 35.2 33.6 44 83.7 34.0 34.3 36.3 45 114.1 24.3 26.3 62.1 73 122. .3 18.3 21.3 76.5 86 117. .7 23.2 25.6 78.0 85 (INTERRUPTED TESTS) 39.9 42.3 99.3* 60.1 38.8 127.9 41.5 34.3 31.2 62 96.8 22.9 27.9 45.9 54 118.9 31.3 27.1 49.5 61 111.1 25.0 22.8 46.4 58 96.6* 24.1 22.9 120.5* 25.8 29.4 113.0* 27.6 20.5 1976/18.85 (INTERRUPTED TESTS) 261.5 1.015 207.2 1.017 592.1 25.8 22.4 590.4 23.596 570.7 27.2 26.9 570.7 26.289 535.5 19.3 23.9 535.2 17.513 2050/15.0 536.8 28.3 27.5 535.6 20.662 497.0 23.7 23.9 496.2 17.600 514.8 23.4 24.4 513.1 12.500 454.1 16.6 35.2 453.7 15.476 420.3* 33.7 33.2 (INTERRUPTED TESTS) 560.1* 22.9 2012/14.5 536.6* 7.3 8.1 2100/12.0 354.1 14.8 36.5 353.8 12.646 343.4* 27.2 491.0* 16.7 1700/50.0 ♦ •* Machined Free Hade Specimen 1% creep β trength is a significant property. Limiting creep ■trains to 1.0% and 2.0% is extremely important to gas turbine component design, since a component's usefulness is generally measured by its resistance to creep to an approximate 1-2% leve not its ultimate rupture strength. Many prior art alloys may exhibit attractive rupture strength at the >190ό'F level, however, they lack the level of ussful strength, i.e., creep strength to 2.0%, that this invention provides in tandem with i far superior strength in test conditions below 1900*F. The CMSX-lOIa specimens also provided significantly increased creep strength at the higher temperature extremes, but it did n appear to develop strength as good as the CMSX-lOGa specimens i lower temperature tests, as indicated by the results in Table 1 below. TABLE 19 P^-19 ll PtEy. pyTVtt Si ilarly, CMSX-12B, with Nv3B at 1.80 level and additional chemistry balance as presented in Table 1, provided attractive creep strength at test condition greater than 1900*F, but did n perform -quite as well as CMSX-lOGa in lower temperature tests, indicated by the results reported in Table 20 below. Alloy composition has the greatest effect on ultimate creep strength. However, some of the variation experienced between alloy derivatives, and particularly for tests exhibiting inconsistent results for a given alloy, can be caused by variation in casting process condition.. Casting process therm gradient variation affects the cast specimen dendrite arm spac and ultimately, its response to solution heat treatment and primary gamma prime aging treatment. It must, therefore, be recognized that much of the creep-rupture results reported herein may have been generated under non-optimized conditions and may be capable of improvement. Improved casting process control may provide casting microstructures more amenable to solution treatment and study to determine the appropriate primary gamma prime aging treatment to provide the optimum gamma prime particle size, which may result in further mechanical property enhancement. The CMSX-12C composition was designed to provide a calculated N v3B number of 1.70. The alloy Cr content was designed at 2.8% and Co set at 3.5% aim for this alloy. An attractive Ta/W ratio was maintained while Re content was moderate at 5.8%. The alloy's Al+Ti content was reduced, in comparison to the CMSX-12A and CMSX-12B specimens, to provide improved alloy response to solution procedure. Similar to the CMSX-lOGa specimen, the CMSX-12C specimen exhibited an improved balance of creep strength for test condition ranging 1800-2100*F, as reported in Table 21 below. TABLE 21 P*«-ia BEff-WTy E 395.7 With improved casting process controls, this specimen has shown the following 1.0% longitudinal creep strengths, as reported in Table 22 below. TABLE 22 Test Condition Time to 1.0* Strain H-C5 1800T/36.0 ksi 455 2100 T/12.0 ksi 309.3 Both alloys provide similarly greater rupture strength than CMS 4 alloy at condition to 1976*F. Respective improvements to met temperature capability are reported below in Table 23. TABLE 23 Approx. Strength Advanta Temperature Relative to CMSX-4 1800T 40'F 1850'F 45*F 1976*F 43'F Based on 1.0% creep strength, the respective approximate advantages are: 1800'F +46*F 1850*F +60T 1976'F +55'F Note that the comparison is not density corrected. For test temperature above 1976*F, the test results indicate tha the CMSX-lOGa and CMSX-12C specimens provided slightly lower strength than CMSX-4 alloy. The reduction in strength advantage for these alloys is believed to be the result of TCP phase formation. To address this issue, the alloys CMSX-lOGb, CHSX-10L, CMSX-12Ca, and CMSX-12E, are designed with N v3B number as low as 1.50 (see Table 1) to provide greater phasial stability, and effect much improved high temperature creep- strength while maintaining most of the creep advantage demonstrated for the 1800-1976*F test regime. The CMSX-lORi and CMSX-12R1 compositions were designed at the •v 1.91 and 1.92 N vJB levels, respectively. These specimens were subjected to the most extensive testing of properties. They wer designed with 2.65% and 3.4% respective Cr levels, with other features remaining similar to the aforementioned alloy design considerations. The properties generated for these two material confirm the overall invention design concept with the other material iterations able to provide similar physical properties and relatively better blends of mechanical properties. The CMSX-lORi and CMSX-12Ri specimens' respective creep-rupture capabilities are reported below in Tables 24 and 25. TABI-E 2-1 PCX-IO-R.) CREEP-RUPTURE The method and results of W and Re microstructural segregation investigation undertaken on fully solutioned and partially εolutioned CMSX-12R1 test specimens are reported in Table 26 below. The investigation indicated that it is desirable to minimize the amount of microstructure-contained residual eutectic and that for fully solutioned specimens, the solution treatments developed for the invention are successful in minimizing elemental segregation, which is important in attaining optimized mechanical properties and microstructural stability. TAB E . 26 Alley! OβX-12 Ri Teat tpeelawn: 3/8" Diameter Solid Bar Specimen Condition: Fully Solutioned Solutioned with 2.OX Residual Eutectic Analyses Method: Mieroprobe Analyses ♦ Random array of 350 points across a section at right angles to the growth direction ♦ Seven line scans, 51* apart, 50 point analyses per line The standard deviation of the U and Re measurements are the Measure of homogeneity Table 27 below reports results of burner rig hot corrosion test undertaken with the CMSX-12R1 specimen. The measurements were taken at the bar location which experienced the maximum attack, i.e., 1652*F location, with the results showing the DS MAR M 002 alloy experienced approximately 20X more metal loss than the CMSX-12Ri specimen. Visual obseirvation showed a similar result for the CMSX-lORi alloy. Both CMSX-lORi alloy and CMSX-12Ri alloy showed similar resistance to attack as CMSX-4 alloy based on visual specimen review at 60, 90 and 120 hours. TAB E 27 NOT CORROSION fgTHOO ♦ 8umtr Rig 1742*F (9S0*C) 2 ppβ βalt, standard fuel ♦ Measurements taken at point of saaxlaua attack which was at 1652*F (900*C) ♦ Measurements reported were taken at the average ninim-n diameter of useful a-etal Table 20 below reports the results of cyclic oxidation tests undertaken at 2012*F with March 1 gas velocity. The CMSX-12Ri specimen was similarly resistant to oxidation attack at 2012*F, however, it was not as good as CMSX-4 at approximately 1886*F exposure. TAB E 28 Cyclic Oxidation Test . 15 Minute Cyeleβ to 2012*F (1100'C), Cooled To Aβ lβ t Retween Cycles Meeh 1 Caβ Velocity . 89 Hours Total With 77 Hours at 2012*F CMgX-12 Bl RESULT: at 1100'C Approx.0.1 en loss per aide for every 300 cycles Q£&zk Approx.0.1 an loss per aide for every 380 cycles at 1030*C CMSX-12 HI Approx. .105 an lose per aide after 355 cycles BMzk Approx. .03 am loss per aide after 355 cycles CMSX-12R1 elevated temperature tensile data is reported in Tabl 29 below, while the results of impact tests are reported in Ta 30 below. The CMSX-12Ri elevated temperature impact strength minimum is similar to CMSX-4 and its maximum occurring at 3-742' is better. TABLE 29 TENSILE DATA CKSX-12 Ri Alloy Teat R 1922 3.8* 69.5 74.3 104.1 19 36 1922 3.8* 72.4 77.6 106.0 19 36 TABLE 3Q IMPA v CMSX-12 Ri 0.35 Inch Diameter Plain Cylindrical SpecI Further Impact Property Compariβon . CMSX-2 -- Mln. Impact Strength 16.5 Joules . SRR 99 -- Mln. I-rpaet Strength 20 Joules The results of CMSX-12R1 low cycle fatigue tests undertaken at 1382*F and 1742 * F test conditions, with R - O, are reported in Table 31 below. The data indicates that CMSX-12Ri performance similar to CMSX-4 at 1382 * F condition, while the alloy exhibit approximately 2.5 times the typical CMSX-4 life at 1742'F condition. TABLE 31 LOU CYCLE FATIGUE CMSX-12 Ri Alloy . Performance Is βlβllar to CMSX-4 at 1382*F (750*C) . Con-pared to CMSX-4. at 1742*f (950*0 and In the 20000 cycle region, CMSX-12RI exhibits 2.5 times life or 15X on βtrength. Notched low cycle fatigue test results show the CMSX-12Ri is 2 1/2 times better than CMSX-4 out to about 30000 cycles, while 50000 cycles and above, the alloy performance is similar to C 4. The results of these tests performed at 1382'F, -^ ■» i.O a R«0 test condition, are reported in Table 32 below. TABLE 32 ROTCHED LOW CYCLE FATICUE CMSX-12R1 Alloy 1382*F (750*C), K t 2.0, R PEAK STRESS CYCLES ksi (MPa) 113.13 (780) 4879 107.33 (740) 9784 95.72 (660) 28470 84.12 (580) 49810 81.22 (560) 78.32 (540) > 115,000 75.42 (520) > 115,000 . Results are 2-1/2 times better than CMSX-4 out to about'3000 cycles. . Results are similar to CMSX-4 at 50000 cycles, and above. High cycle fatigue test results for the CMSX-lORi alloy are reported in Table 33 below. For 1742 * F, 100 Hz, R - O test condition, the alloy exhibited about 2-1/2 times the typical CMSX-4 lives. HfgH CTCVE FπTlWE CMSX-lORi Alloy 1742*F (950*C), 100 Hx., R ■ 0 PEAK STRESS CYCLES ksi (MPa) (N f ) 81.22 (560) 15.2 x IO 6 92.82 (640) 3.59 x 10* 104.43 (720) 0.6 X IO 6 * Lives are 2-1/2 times better than CMSX-4 The CMSX-lORi and CMSX-12R1 test data indicates that adequate corrosion and oxidation resistance can be achieved with extrem low alloy chromium content. Additionally, extremely good ther mechanical fatigue tensile and impact strengths are apparent w the superalloys of this invention. The results of alloy specimen density measurements are reporte in Table 34 below. TABLE 34 pit-GLE CRYSTAL ALLOT DENSITY DATA The alloys of this invention are amenable to HIP processing. Specimens HIP treated as reported in Table 35 below, showed nearly complete pore closure and absence from incipient melting. IP condition 1. Heat Specimens in the HIP vessel to 2455 β F at mininuiα Argon pressure (approximately 1500 psi) and hold for 4 hours while maintaining 2455°F/1500 psi condition. 2. While maintaining the 2 55*F operating temperature, increase the Argon pressure over 1 hour to 20 ksi. Soak specimens at 2455°F/20 ksi condition for 4 hours. Further evaluation of the CMSX-lORi and CMSX-12Ri alloys was undertaken. The data reported below further shows the unique capabilities exhibited by the alloys of this invention. More specifically, very attractive hot corrosion and dynamic oxidation resistance are exhibited by these alloys in tandem with exceptionally high creep-rupture, thermo-mechanical fatigue, tensile and impact strengths, despite the extremely low level of chromium content employed, thereby providing a unique blend of desirable properties. Additional tensile data is reported in Table 36 below for the CMSX-lORi and CMSX-12Ri alloys. These alloys were evaluated at temperatures ranging from Room Temperature (RT) to about 2100°F. This data complements the data reported in Table 29 above. TABLE 36 Tensile Data CMSX-10RΪ Alloy TABLE 36. CONTINUED CMSX-12Ri Alloy Further results of impact tests, in addition to the results reported in Table 30 above, are reported in Table 37 below. Bot the CMSX-lORi and CMSX-12Ri compositions were evaluated at temperatures ranging from 1382-1922°F. TABLE 37 Impact Data Plain low cycle fatigue data for the CMSX-lORi and CMSX-12R specimens is reported in Table 38 below. This data complements the data reported in Table 31 above. The tests were performed a 1382°F and 1742°F, with R=0 and 0.25 Hz test conditions. TABLE 38 Plain Low Cycle Fatigue Data CMSX- 12Ri Al loy; 1742°F, R=0, 0.25 Hz CMSX-12RJ Al loy; 1382°F, R=0, 0.25 Hz The results of notched low cycle fatigue tests (K t = 2.2) undertaken with CMSX-lORi and CMSX-12Ri specimens is reported in Table 39 below. The tests were performed at 1382°F and 1742°F, with R=0 and 0.25 Hz test conditions. TABLE 39 K*. = 2.2 NOTCHED LOW CYCLE FATIGUE DATA CMSX- 10Ri Al loy; 1742°F, R=0, 0.25 Hz CMSX-12Ri Al loy; 1382°F, R=0, 0.25 Hz CMSX- 12RJ Al loy; 1742°F, R=0, 0.25 Hz High cycle fatigue data, which complements the data report in Table 33 above, is reported in Table 40 below. The data is reported for tests undertaken with CMSX-lORi and CMSX-12Ri specimens at the respective test conditions of: a)1742°F, R=0, 100 Hz; and b) 1022°F, R=-l, 100Hz. TABLE 40 HIGH CYCLE FATIGUE DATA CMSX-lORi Alloy; 1742°F, R---0, 100 Hz CMSX-12Ri Al loy; TESTED AT 1022°F, R=- 1 , 100 Hz Bare alloy oxidation data are presented in Figures 3 and 4. Both Figures compare the test results for CMSX-lORi and CMSX-12R alloys to results from identical tests performed on the DSM002, CMSX-4 and CMSX-4+Y (120 ppm) alloys. The data presented in these Figures complements the data reported in Table 28 above. Figure 3 illustrates the results of tests performed out to about 375 cycles at about 2012°F, while Figure 4 illustrates the results obtained at about 1886°F. Figure 5 illustrates a similar alloy capability comparison, except for bare alloy hot corrosion resistance. The burner-rig tests were performed to about 120 hours duration at about 1742°F with ingestion of 2 ppm salt. Additional alloy compositions were defined in accordance with this invention for the purpose of further optimizing alloy solution heat treatment characteristics, alloy stability, and creep-rupture strength, while maintaining the extremely good oxidation, hot corrosion and thermo-mechanical fatigue strengths already achieved with the invention. Table 41 below reports the aim chemistries for additional materials tested, which provide an even better blend of useful engineering property characteristics than the above-described alloys. These optimized compositions typically contain lower chromium and cobalt contents than the materials described in Table 1 above. Furthermore, these optimized alloys typically contain higher rhenium, lower tungsten and lower phasial stability number N V3B . TABLE 41 Key: 1 - U + Re 2 - Al + Ti 3 - Al + Ti + Ta + Cb 4 - W + Re + Mo + Ta * Calculated using PWA N-35 Method Table 42 below reports the chemistries for additional alloy heats (200-320 lb. quantities) . As with the alloys reported in Table 1 above, the alloys having optimized compositions were produced using production-type procedures, and were made into test bars and test blades by vacuum investment casting. TABLE 42 V-1 VIM FURNACE HEAT CHEMISTRIES Table 43 below reports heat treatment detail for the optimized alloy compositions. They typically require solution heat treatment to 2490-2500°F peak process temperature, which generally results in complete γ' solutioning without the occurrence of incipient melting. A three-step aging treatment was utilized. TABLE 43 Heat Treatment Detail *Determined by visual estimation -Specimens air cooled from all aging treatments The optimized alloy compositions were evaluated. However, significant data was generated for the CMSX-10K and CMSX-12D compositions, while some test results were generated for the remaining compositions reported in Table 41 above. Table 44 below reports the results of stress- and creep- rupture tests undertaken with CMSX-10K specimens. Variously sized test bar and blade specimens, which were cast at industrial sources, were utilzed for these tests. The test results show extremely good stress and creep-rupture properties for these alloys . TABLE 44 CMSX-10K ALLOY STRESS- AND CREEP-RUPTURE DATA TABLE 44. CONTINUED CMSX-10K ALLOY STRESS- AND CREEP-RUPTURE DATA TABLE 44. CONTINUED CMSX- 10K ALLOY STRESS- AND CREEP-RUPTURE DATA TABLE 44. CONTINUED CMSX-10K ALLOY STRESS- AND CREEP-RUPTURE DATA * 2106°F/6 Hr/AC Primary Age ** Interrupted Creep-rupture Test. All results with 2075°F/6 Hr/AC Primary Age. CMSX-IOK alloy tensile data are reported in Table 45 below. The tests were performed at temperatures ranging from RT-2012°F. The test results show improvement relative to earlier alloy test results presented in Table 36 above. TABLE 45 Tensile Data CMSX-10K Alloy Dynamic elastic modulus data are reported in Table 46 below. The CMSX-IOK alloy data are reported for test conditions ranging from RT-2012°F. TABLE 46 Dynamic Elastic Modulus, E CMSX-10K Alloy Elevated temperature (1202-1922°F) impact data for CMSX-IOK and CMSX-12D specimens are reported in Table 47 below. Compariso to other material capabilities can be made by review of the respective Tables 30 and 37, above. TABLE 47 Impact Data Plain low cycle fatigue data for CMSX-10K and CMSX-12D specimens is reported in Table 48. The tests were performed at 1382°F, 1562°F and 1742°F, respectively, with R=0 and 0.25 Hz test conditions. Comparative results are found above in Tables 31 and 38, respectively. TABLE 48 Plain Low Cycle Fatigue Data CMSX-10K Alloy; R=0, 0.25 Hz * Test Stopped: Specimen pulled out of threads TABLE 48, CONTINUED CMSX-12D ALLOY; R=0, 0.25 Hz CMSX-IOK notched low cycle fatigue (K t = 2.2) data is reported in Table 49 below. The tests were performed at 1742°F, R=0 and 0.25 Hz test conditions. Comparative results are found above in Tables 32 and 39, respectively. TABLE 49 K.=2.2 Notched Low Cycle Fatigue Data Bare alloy oxidation data for CMSX-10K and CMSX-12D specimens are reported in Table 50 below. The tests were performed with test conditions of 2012°F, 0.25 ppm salt and cycled four times per hour. Test duration was 200 hours, or 800 cycles. The test results indicate that both alloys exhibit extremely good resistance to oxidation. TABLE 50 Bare Alloy Oxidation Data Figure 6 illustrates the results of initial measurements taken on corrosion pins having CMSX-10K and CMSX-12D compositions, which were run at 1742°F, 2 ppm salt conditions out to about 200 hours. CMSX-12D stress- and creep-rupture data are reported in Table 51 below, while this alloy's respective tensile properties for tests conducted at RT-2102°F conditions are reported in Table 52 below. TABLE 51 CMSX-12D ALLOY STRESS- AND CREEP-RUPTURE DATA TIME IN HOURS FINAL CREEP READING TO REACH TEST CONDITION RA t HRS % DEFORMATION 1.0% 2. 0% 1675°F/75.0 ksi (913°C/517 MPa) 1750°F/50.0 ksi (954°C/345 MPa) 1800°F/36.0 ksi (982°C/248 MPa) 1850°F/36.0 ksi (1010°C/248 MPa) 1922°F/32.63 ksi (1050°C/225 MPa) 1976°F/28.1 ksi (1080°C/194 MPa) 1976°F/18.85 ksi (1080°C/130 MPa) 300.5 TABLE 51. CONTINUED CMSX-12D ALLOY STRESS- AND CREEP-RUPTURE DATA TIME IN HOURS RUPTURE TIME FINAL CREEP READING TO REACH TEST CONDITION HRS % El % RA . HRS % DEFORMATION 1.0% 2.0% 1995°F/27.5 ksi 77.0 27.8 37.0 75.9 15.867 25.3 36.4 (1090°C/190 MPa) TABLE 51. CONTINUED CMSX-12D ALLOY STRESS- AND CREEP-RUPTURE DATA * 2106°F/6 Hr/AC Primary Age ** Interrupted Creep-rupture Test. All results with 2075°F/6 Hr/AC Primary Age. TABLE 52 Tensile Data Initial creep-rupture test data for CMSX-10M and CMSX-12F specimens is reported below in Tables 53 and 54, respectively. TABLE 53 CMSX-10M ALLOY STRESS- AND CREEP-RUPTURE DATA RUPTURE TIME FINAL CREEP READING TIME IN HRS TO REACH TEST CONDITION HRS % EL % RA t. HOURS % DEFORMATION 1.0% 2.0% 1675°F/75.0 ksi * 241.4 25.2 32.5 240.6 21.422 15.5 56.3 (913°C/517 MPa) 1750°F/50.0 ksi * 433.4 23.5 30.6 433.4 22.209 224.1 271.1 (954°C/345 MPa) 1800°F/36.0 ksi 812.5 (982°C/248 MPa) 832.5 902.3 1850°F/36.0 ksi 368.9 (1010°C/248 MPa) 338.7 306.6 279.2 300.4 TABLE 53 . CONTINUED CMSX-10M ALLOY STRESS- AND CREEP-RUPTURE DATA 2030°F/18.85 ksi * 278.9 13.9 30.3 277.8 12.367 80.6 160.8 (1110°C/130 MPa) 2200-F/10.0 Ksi * 186.1 18.2 44.9 185.3 17.445 15.4 69.4 (1204°C/69 MPa) * 2106-F/6 Hr/AC Primary Age. ** Interrupted Creep- rupture Test . TABLE 54 CMSX-12F ALLOY STRESS- AND CREEP-RUPTURE DATA RUPTURE TIME FINAL CREEP READING TIME IN HRS TO REACH TEST CONDITION HRS % EL % RA t. HOURS % DEFORMATION 1.0% 2.0% 1675°F/75.0 ksi * 199.8 26.6 34.2 197.0 21.502 8.1 34.0 (913°C/517 MPa) 2030°F/18.85 ksi * 355.4 26.8 27.1 355.1 26.607 139.7 188.1 (1110°C/130 MPa) 2150°F/12.0 ksi * ** 40.8 89.7 (1177°C/83 MPa) 2200-F/10.0 ksi (1204°C/69 MPa) * 2106°F/6 Hr/AC Primary Age ** Interrupted Creep-rupture Test. hile this invention has been described with respect to particular embodiments thereof, it is apparent that numerous other forms and modifications of this invention will be obvious to those skilled in the art. The appended claims and this invention generally should be construed to cover all such obvious forms and modifications which are within the true spirit and scope of the present information.

Claims: CLAIMS What is claimed is: 1. A single crystal casting to be used under high stress, high temperature conditions up to about 2030°F characterized by an increased resistance to creep under such conditions, said casting being made from a nickel-based superalloy consisting essentially of the following elements in percent by weight : Rhenium 6.2-6.8 Chromium 1.8-2.5 Cobalt 1.5-2.5 Tantalum 8.0-9.0 Tungsten 3.5-6.0 Aluminum 5.5-6.1 Titanium 0.1-0.5 Columbium 0.01-0.1 Molybdenum 0.25-0.60 Hafnium 0-0.05 Carbon 0-0.04 Boron 0-0.01 Yttrium 0-0.01 Cerium 0-0.01 Lanthanum 0-0.01 Manganese 0-0.04 Silicon 0-0.05 Zirconium 0-0.01 Sulfur 0-0.001 Vanadium 0-0.10 Nickel + Incidental balance Impurities said superalloy having a phasial stability number N V3B less than about 1.65. 2. The single crystal casting of Claim 1 wherein said casting has been aged at a temperature of from 2050°F to 2125°F for 1 to 20 hours. 3. The single crystal casting of Claim 1 wherein said casting is a component for a turbine engine. 4. The single crystal casting of Claim 1 wherein said casting is a gas turbine blade. 5. The single crystal casting of Claim 1 wherein said casting is a gas turbine vane. 6. The single crystal casting of Claim 1 wherein said casting is further characterized by increased oxidation and hot corrosion resistance under said conditions.

Inventors: ERICKSON GARY L, ERICKSON, GARY, L.
Applicants: CANNON MUSKEGON CORP, ERICKSON GARY L, CANNON-MUSKEGON CORPORATION
Application year: 1993
Main classification: C22C19
IPC codes: C22C19, C30B11, C30B29