NICKEL BASE SUPERALLOY

Abstract: A novel alloy is provided having approximately the following ingredient formula: The alloy has a solvous temperature of below 1200°C. Fatigue crack propagation rate is low for metal samples cooled at rates of 50°C/min to 200°C/min. The alloy has uniquely high strength at temperatures at and above 1200°F.

Description: The subject matter of this application relates generally to that of EP-A-0260510, EP-A-0260511, EP-A-0260512 and EP-A-0260513, all of which were published on 23 March 1988. The subject application also relates generally to the subject matter of US-A-4685977, published on 11 August 1987. It is well known that nickel based superalloys are extensively employed in high performance environments. Such alloys have been used extensively in jet engines and in gas turbines where they must retain high strength and other desirable physical properties at elevated temperatures of a 1000°F or more. Many of the nickel-based superalloys depend for part of their strength and other properties at high temperature on γ' precipitates. Some detailed characteristics of the phase chemistry of γ' are given in "Phase Chemistries in Precipitation-Strengthening Superalloy" by E.L. Hall, Y.M. Kouh and K.M. Chang, "Proceedings of 41st Annual Meeting of Electron Microscopy Society of America, August 1983 (p. 248)]. The following U.S. patents disclose various nickel-base alloy compositions: U.S. 2,570,193; U.S. 2,621,122; U.S. 3,046,108; U.S. 3,061,426; U.S. 3,151,981; U.S. 3,166,412; U.S. 3,322,534; U.S. 3,343,950; U.S. 3,575,734; U.S. 3,576,681, U.S. 4,207,098 and U.S. 4,336,312. The aforementioned patents are representative of the many alloying situations reported to date in which many of the same elements are combined to achieve distinctly different functional relationships between the elements such that phases providing the alloy system with different physical and mechanical characteristics are formed. Nevertheless, despite the large amount of data available concerning the nickel-base alloys, it is still not possible for workers in the art to predict with any degree of accuracy the physical and mechanical properties that will be displayed by certain concentrations of known elements used in combination to form such alloys even though such combination may fall within broad, generalized teachings in the art, particularly when the alloys are processed using heat treatments different from those previously employed. Most of the alloys that are processed either by cast and wrought or by powder metallurgy technologies can only retain their superior strength at temperatures up to about 649°C (1200°F). A dramatic drop in strength is observed when the temperature of these alloys is raised above 649°C (1200°F). A problem which has been recognized to a greater and greater degree with many nickel based superalloys is that they are subject to formation of cracks or incipient cracks, either in fabrication or in use, and that the cracks can actually initiate or propagate or grow while under stress as during use of the alloys in such structures as gas turbines and jet engines. The propagation or enlargement of cracks can lead to part fracture or other failure. The consequence of the failure of the moving mechanical part due to crack formation and propagation is well understood. In jet engines it can be particularly hazardous and can be catastrophic. However, what has been poorly understood until recent studies were conducted was that the formation and the propagation of cracks in structures formed of superalloys is not a monolithic phenomena in which all cracks are formed and propagated by the same mechanism and at the same rate and according to the same criteria. By contrast the complexity of the crack generation and propagation and of the crack phenomena generally and the interdependence of such propagation with the manner in which stress is applied is a subject on which important new information has been gathered in recent years. The period during which stress is applied to a member to develop or propagate a crack, the intensity of the stress applied, the rate of application and of removal of stress to and from the member and the schedule of this application was not well understood in the industry until a study was conducted under contract to the National Aeronautics and Space Administration. This study is reported to a technical report identified as NASA CR-165123 issued from the National Aeronautics and Space Administration in August 1980, identified as "Evaluation of the Cyclic Behavior of Aircraft Turbine Disk Alloys" Part II, Final Report, by B.A. Cowles, J.R. Warren and F.K. Hauke, and prepared for the National Aeronautics and Space Administration, NASA Lewis Research Center, Contract NAS3-21379. A principal unique finding of the NASA sponsored study was that the rate of propagation based on fatigue phenomena or in other words the rate of fatigue crack propagation (FCP) was not uniform for all stresses applied nor to all manners of applications of stress. More importantly, the finding was that fatigue crack propagation actually varied with the frequency of the application of stress to the member where the stress was applied in a manner to enlarge the crack. More surprising still, was the finding from the NASA sponsored study that the application of stress of lower frequencies rather than at the higher frequencies previously employed in studies, actually increased the rate of crack propagation. In other words the NASA study revealed that there was a time dependence in fatigue crack propagation. Further the time dependence of fatigue crack propagation was found to depend not on frequency alone but on the time during which the member was held under stress for a so-called hold-time. Following the discovery of this unusual and unexpected phenomenon of increased fatigue crack propagation at lower stress frequencies there was some belief in the industry that this newly discovered phenomenon represented an ultimate limitation on the ability of the nickel based superalloys to be employed in the stress bearing parts of the turbines and aircraft engines and that all design effort had to be made to design around this problem. However, it has now been discovered that it is feasible to construct parts of nickel based superalloys for use at high stress in turbines and aircraft engines with greatly reduced crack propagation rates. Crack growth, i.e., the crack propagation rate, in high-strength alloy bodies is known to depend upon the applied stress (σ) as well as the crack length (a). These two factors are combined by fracture mechanics to form one single crack growth driving force; namely, stress intensity K, which is proportional to σ√ a . Under the fatigue condition, the stress intensity in a fatigue cycle represents the maximum variation of cyclic stress intensity (ΔK), i.e., the difference between K max and K min . At moderate temperatures, crack growth is determined primarily by the cyclic stress intensity (ΔK) until the static fracture toughness K IC is reached. Crack growth rate is expressed mathematically as da/dN ∝(ΔK) n . N represents the number of cycles and n is a constant which is between 2 and 4. At high temperatures, however, the cyclic frequency and the shape of the waveform become the important parameters determining the crack growth rate. For a given cyclic stress intensity, a slower cyclic frequency can result in a faster crack growth rate. This undesirable time-dependent behavior of fatigue crack propagation can occur in most existing high strength superalloys. The most undesirable time-dependent crack-growth behavior has been found to occur when a hold time is superimposed on a sine wave variation in stress. In such case a test sample may be subjected to stress in a sine wave pattern but when the sample is at maximum stress the stress is held constant for a hold time. When the hold time is completed the sine wave application of stress is resumed. According to this hold time pattern the stress is held for a designated hold time each time the stress reaches a maximum in following the normal sine curve. This hold time pattern of application of stress is a separate criteria for studying crack growth. This type of hold time pattern was used in the NASA study referred to above. Progress has been made in forming superalloy metal compositions containing high volume percents of strengthening precipitates and in processing of these metals into parts for advanced turbine engines and jet aircraft engines. This metal processing technology has been developed to introduce such superalloys into gas turbines and jet engines because of the higher temperature capabilities of the alloys themselves and because the engines built with such alloys also have higher temperature capabilities and resulting higher efficiencies and thrust per unit weight of engine. While some studies such as the NASA studies described above have been made of a number of these alloys, not all of the alloys have been examined comprehensively with respect to fatigue cracking and with respect to resistance to fatigue cracking. It has been determined that at low temperatures the fatigue crack propagation depends essentially entirely on the intensity at which stress is applied to components and parts of such structures in a cyclic fashion. As is partially explained in the background statement above, the crack growth rate at elevated temperatures cannot be determined simply as a function of the applied cyclic stress intensity ΔK. Rather the fatigue frequency can also affect the propagation rate. The NASA study demonstrated that the slower the cyclic frequency, the faster the crack grows per unit cycle of applied stress. It has also been observed that faster crack propagation occurs when a hold time is applied during the fatigue cycle. Time-dependence is a term which is applied to such cracking behavior at elevated temperatures where the fatigue frequency and hold time are significant parameters. It is known that some of the most demanding sets of properties for superalloys are those which are needed in connection with jet engine construction. Of the sets of properties which are needed, those which are needed for the moving parts of the engine are usually greater than those needed for static parts although the sets of needed properties are different for the different components of an engine. Because some sets of properties have not been attainable in cast alloy materials, resort is sometimes had to the preparation of parts by powder metallurgy techniques. However, one of the limitations which attends the use of powder metallurgy techniques in preparing moving parts for jet engines is that of the purity of the powder. If the powder contains impurities such as a speck of ceramic or oxide, the place where that speck occurs in the moving part becomes a latent weak spot where a crack may initiate or a latent crack. As alloy products for use in turbines and jet engines have developed, it has become apparent that different sets of properties are needed for parts which are employed in different parts of the engine or turbine. For jet engines, the material requirements of more advanced aircraft engines continue to become more strict as the performance requirements of the aircraft engines are increased. The different requirements are evidenced, for example, by the fact that many blade alloys display very good high temperature properties in the cast form. However, the direct conversion of cast blade alloys into disk alloys is very unlikely because blade alloys display inadequate strength at intermediate temperatures of about 700°C. Further, the blade alloys have been found very difficult to forge and forging has been found desirable in the fabrication of blade from disk alloys. Moreover, the crack growth resistance of disk alloys has not been evaluated. Accordingly, to achieve increased engine efficiency and greater performance, constant demands are made for improvements in the strength and temperature capability of disk alloys as a special group of alloys for use in aircraft engines. Now these capabilities must be coupled with low fatigue crack propagation rates and a low ordr of time-dependency of such rates. In one aspect, the present invention provides a nickel base superalloy composition consisting of the following composition apart from impurities Element Composition in weight % From To Ni balance Cr 8 12 Co 13 18 Mo 3 6 W 3 6 Al 3 4 Ti 2 4 Ta 6 9 Zr 0.01 0.05 B 0.01 0.05 C 0.01 0.05 and which has been supersolvous annealed and cooled at a rate of less than 250°C./min. An alloy embodying the said one aspect has a low fatigue crack propagation rate, has been supersolvous annealed and has been cooled at a rate between 20°C/min and 200°C/min. An alloy embodying the said one aspect has a tensile strength above 133583 Kg/m2 (190 Ksi) at temperatures of (1200°F) 649°C and (1400°F) 760°C. In an alloy embodying the said one aspect, the solvous temperature is below 1200°C. An alloy in accordance with said one aspect has the composition Element Composition in weight % Ni balance Cr 10.0 Co 15.0 Mo 5.0 W 5.0 Al 3.5 Ti 3.0 Ta 7.2 Zr 0.03 B 0.03 C 0.03 The nominal composition is the composition containing percentages of ingredients by weight which were specified and sought in preparing compositions as set out in the Examples below. The range of compositions set forth the percentage of ingredients which are believed to provide novel properties. According to another aspect of the invention, there is provided the method of preparing a nickel base superalloy which comprises preparing a melt consisting of the following composition apart from impurities Element Composition in weight % From To Ni balance Cr 8 12 Co 13 18 Mo 3 6 W 3 6 Al 3 4 Ti 2 4 Ta 6 9 Zr 0.01 0.05 B 0.01 0.05 C 0.01 0.05 cooling the melt to a solid, supersolvous annealing the solid, and cooling the solid at a rate of 250°C./min. or less. In preparing the alloy, the steps of which may be employed include melting the composition to form a melt, atomizing the melt by gas atomization, consolidating the powder by hot isotatic pressing and forging. The forges specimens may then be solution annealed at 1200°C for one hour to provide a supersolvous anneal followed by cooling at a selected cooling rate. The cooling rate may be between 50°C/min and 200°C/min. the samples may then be given an isothermal anneal at about 760°C for about 16 hours, followed by cooling the alloy. Superalloys embodying the invention can be prepared by casting and the cast alloy can be wrought. Superalloys embodying the invention can be prepared by advanced metal processing procedures such as powder metallurgy procedures and spray forming procedures such as by the commercially known OSPREY (Registered Trade Mark) process. A superalloy prepared by such advanced processing procedures can also be effectively wrought or forged. The present invention also encompasses a method for processing a superalloy to produce material with a superior set or combination of properties for use in advanced engine disk applications. Embodiments of the invention may provide nickel-base superalloy products which have high strength at high temperature and are more resistant to cracking. Embodiments of the invention may reduce the tendency of nickel-base superalloys capable of use at temperatures above 649°C (1200°F) to undergo cracking. Embodiments of the invention may provide articles for use under cyclic high stress which have good strength above 649°C (1200°F) and are more resistant to fatigue crack propagation. A composition and method embodying the invention may permit nickel-base superalloys to have imparted thereto high strength above 649°C (1200°F) and resistance to cracking under stress which is applied cyclically over a range of frequencies. It appears that embodiments of the invention provide alloy compositions which have high strength at temperatures above 649°C (1200°F) and which have reduced time dependence of fatigue cracking resistance. Embodiments may also provide extended stress rupture life. It also appears that embodiments may be provided which permit a broad range of variation in the cooling rate but still provide the desired combination of strength at high temperature together with resistance to time dependent fatigue crack propagation. Also, an embodiment may be provided in which this broad range of cooling rates does not detract from other properties of the alloy as the alloy itself does have a good combination of strength and rupture life. The combination of high temperature strength and other properties appears to be achieved at cooling rates over any part of a broad range where such cooling is responsible for the resulting properties of the time dependent fatigue crack propagation resistant alloy. An embodiment of the invention may provide a disk alloy which has superior strength at temperatures above 649°C (1200°F) and which has a lower or minimum time dependence of fatigue crack propagation and moreover a high resistance to fatigue cracking. Further, it may have high strength and long stress rupture life. An embodiment of the present invention, given by way of non-limiting example will now be described with reference to the accompanying drawings, in which: FIGURES 1 and 2 are graphic (log-log plot) representations of fatigue crack growth rates (da/dN) obtained at 694°C (1200°F) at various stress intensities (ΔK) for the same alloy, after different cooling rates from elevated supersolvus solution annealing temperatures, under cyclic stress applications at a series of frequencies one of which cyclic stress applications includes a hold time at maximum stress intensity. FIGURE 3 is a bar graph in which fatigue crack propogation rate is plotted on a comparative basis for the three different cycles which were studied. The properties which are conventionally needed for materials used in engine disk applications include high tensile strength and high stress rupture strength. The ability to resist crack growth is essential for the component low cycle fatigue life or LCF. Example 1 An alloy powder was prepared by vacuum melting and gas atomization. The alloy was first melted by vacuum induction melting and was then gas atomized using argon as the atomizing gas. The ingredient content of the alloy was prescribed according to the following composition: NOMINAL COMPOSITION OF CH-88A wt.% at.% Ni bal. bal. Co 15.0 15.56 Cr 10.0 11.75 Mo 5.0 3.19 W 5.0 1.66 Al 3.5 7.93 Ti 3.0 3.83 Ta 7.2 2.43 Zr 0.03 0.02 B 0.03 0.17 C 0.03 0.15 The alloy of Table I is a superalloy which forms a strengthening γ' precipitate. The alloy was designated alloy CH-88A. A master heat having a composition as set forth in Table I was prepared by vacuum induction melting as described above to form an 18 lb. ingot. Powder atomization was performed in a gas atomizer using argon gas. Screened powder of -140 mesh of the US series were collected. The collected powder was placed in a metal can as conventionally used for hot isostatic pressing and the can was evacuated and sealed. The sealed can was HIPped at 1175°C (2147°F) at 15 ksi (103 MPa) for four hours. The as-HIPped can had a rectanguloid form and dimensions of 38mm by 76mm by 101mm (1.5 inches by 3.0 inches by 4.0 inches). The as-HIPped can was hot pressed several times at 1150°C, the last pressing of which reduced one dimension by 37.5%. The specimen which had been forged was cut into smaller specimens designated as blanks and the blanks were subjected to controlled heat treatments including a cooling at a controlled rate. The precipitate solvous of the sample was determined by metallographic technique to be 1185°C. Solution annealing was carried out at 1200°C for one hour for all the blanks followed by cooling at two different rates. Half the solution annealed samples were cooled at 150°C/min. and the other half were cooled at 75°C/min. The cooling at 150°C/min. was accomplished by chamber cooling. The cooling at 75°C/min was accomplished by steel-can cooling. After the solution treatment and cooling the alloy specimens received a single aging treatment at 760°C for 16 hours (1400°F/16 hours). Standard round tensile bars of 2.54mm (0.10 in.) gauge diameter were machined and low-stress ground for tensile testing. The tensile properties as functions of the testing temperature are shown in Table 2. Alloy CH-88A exhibits good strength at 649°C (1200°F) and also at 760 (1400°F). Both tensile and yield strength start to decrease above 760°C (1400°F). Samples of the two materials CH-88A and Rene' 95 which had been solution annealed and cooled at the rate indicated in Table 2 were tested at the two temperatures 649°C (1200°F) and 760°C (1400°F). Test results are given in Table 2 below. (RENE is a Registered Trade Mark.) High Temperature Tensile Properties Alloy Cooling Rate (C/min) Test Temp °C(F) Yield Strength (ksi) Tensile Strength (ksi) CH-88A 150 649(1200) 173 240 760(1400) 169 203 CH-88A 75 649(1200) 164 234 760(1400) 157 192 Rene' 95 150 649(1200) 158 223 760(1400) 152 173 Rene' 95 75 649(1200) 152 215 760(1400) 146 165 It was observed that a slow cooling rate results in a lower strength at both testing temperatures and for both alloys but that the CH-88A sample retained good strength at 760°C (1400°F) even where the slower cooling rate was employed. Rene' 95 is a conventional superalloy which is known to have high strength properties. It is evident from the data of Table 2 that the CH-88A alloy has high temperature strength properties which are superior to those of the Rene' 95 alloy both at 649°C (1200°F) and at 760°C (1400°F). Fatigue crack propagation (FCP) tests were performed by employing a single-edge notched (SEN) specimen and a dc electric potential drop technique. Different cyclic waveforms, as described above with reference to the NASA study, were employed in the testing. The cycles included standard 20 cycles per minute (cpm) sinusoidal cycle (3 second per cycle); 60 times slower at 0.33 cpm sinusoidal cycle (180 second per cycle); and 20 cpm sinusoidal cycle with a 177 second hold at peak load (180 second per cycle). The minimum to maximum load ratio R was kept at 0.05 in all waveforms. The da/dN curves for alloy CH-88A alloy samples are plotted in Figures 1 and 2. The CH-88A alloy displays no time dependance of this crack growth rate even for the fast cooling rate of 150°C per minute after a supersolvus anneal. To provide a basis for comparison of crack growth rate of the CH-88A alloy relative to a conventional alloy data is plotted in Figure 3 to display their comparative similarities and differences. The comparative data plotted is for a crack growth rate of ΔK = 0.106 (Kg/mm2) √ mm (30 ksi √ in ) . This data was obtained after both the CH-88A alloy and the Rene' 95 had undergone a supersolvous anneal and had been cooled at the game cooling rate. What is noteworthy about the CH-88A alloy is that it can be subjected to a range of cooling rates extending from 75°C/min to 150°C/min without very substantial change in the resultant strength of the alloy. Measurements were also made of the relative strengths of each of the samples prepared, as indicated above, and the relationship or the function of the cooling rate relative to the strength is listed in Table 2. It is evident from Table 2 that generally the higher cooling rates favor higher tensile strength and yield strength. However, the tensile and yield strength achieved at the lower heating rate is still very substantial. As indicated in the table, the heating of the tensile and yield strengths were measured 649°C (1200°F) and at 700°C (1400°F). For comparison, similar measurements were made with samples of Rene' 95 which were also supersolvous annealed and cooled at different rates. Rene' 95 is known to be the strongest commercially available superalloy. The CH-88A alloy has an advantage of 66°C (150°F) to 93°C (200°F) over the strengths measured for the Rene' 95 alloy sample.

Claims: A nickel base superalloy composition consisting of the following composition apart from impurities Element Composition in weight % From To Ni balance Cr 8 12 Co 13 18 Mo 3 6 W 3 6 Al 3 4 Ti 2 4 Ta 6 9 Zr 0.01 0.05 B 0.01 0.05 C 0.01 0.05 and which has been supersolvous annealed and cooled at a rate of less than 250°C./min. The alloy of claim 1 which has a low fatigue crack propagation rate, which has been supersolvous annealed and which has been cooled at a rate between 20°C./min and 200°C./min. The alloy of claim 1 in which the tensile strength is above 133583 Kg/m2 (190 ksi) at temperatures of 649°C (1200°F) to 760°C (1400°F). The alloy of claim 1 in which the solvous temperature is below 1200°C. The alloy of claim 1 in which the composition is as follows: Element Composition in weight % Ni balance Cr 10.0 Co 15.0 Mo 5.0 W 5.0 Al 3.5 Ti 3.0 Ta 7.2 Zr 0.03 B 0.03 C 0.03 The method of preparing a nickel base superalloy which comprises preparing a melt consisting of the following composition apart from impurities. Element Composition in weight % From To Ni balance Cr 8 12 Co 13 18 Mo 3 6 W 3 6 Al 3 4 Ti 2 4 Ta 6 9 Zr 0.01 0.05 B 0.01 0.05 C 0.01 0.05 cooling the melt to a solid, supersolvous annealing the solid, and cooling the solid at a rate of 250°C./min. or less. The method of claim 6 in which the cooling rate is between 50°C./min and 200°C./min. The method of claim 6 in which the solid is subjected to heat aging following the cooling. The method of claim 6 in which the solid is subjected to heat aging at about 760°C. for about 16 hours. The method of claim 6 in which the composition is as follows: Element Composition in weight % Ni balance Cr 10.0 Co 15.0 Mo 5.0 W 5.0 Al 3.5 Ti 3.0 Ta 7.2 Zr 0.03 B 0.03 C 0.03 The method of claim 6 in which the alloy is forged after cooling from the melt. The method of claim 6 in which the cooling of the melt to a solid is by powder atomization. The method of claim 6 in which the cooling of the melt to a solid is by spray forming.

Inventors: CHANG, KEH-MINN
Applicants: GENERAL ELECTRIC COMPANY
Application year: 1988
Main classification: C22C19
IPC codes: C22C19, C22F1