GB2235697A

GB2235697A - Nickel-base superalloys

Info

Publication number: GB2235697A
Application number: GB8631037A
Authority: GB
Inventors: Carl Stephen Wukusick; Jr Leo Buchakjian
Original assignee: General Electric Co
Current assignee: General Electric Co
Priority date: 1986-12-30
Filing date: 1986-12-30
Publication date: 1991-03-13
Anticipated expiration: 2006-12-30
Also published as: GB8631037D0; AU621149B2; AU4163989A; CA1327132C; FR2654114A1; FR2654114B1; GB2235697B

Abstract

Gas turbine engine efficiencies are increased through improvements in nickel-base superalloys used to make parts and components for gas turbine engines. These improved superalloys consist essentially of, in percentages by weight, 5-10 Or, 5-10 Co, 0-2 Mo, 3-10W. 3-8 Ta, 0-2 Ti, 5-7 Al, 0-6 Re, 0-0.5 Hf; 0-0.07C, 0-0.015B and 0-0.075Y, the balance being Ni and incidental impurities. They produce single crystal articles having improvements in temperature capability, based on stress rupture strength and low and high cycle fatigue properties. Because of their superior resistance to degradation by cyclic oxidation, and their resistance to hot corrosion, the superalloys possess a unique balance in mechanical and environmental properties.

Description

IMPROVED AND PROPERTY-BALANCED NICKEL-BASE SUPERALLOYS FOR PRODUCING SINGLE CRYSTAL ARTICLES The invention disclosed and claimed herein is related to the invention disclosed and claimed in co-assigned application Ser. No. 595,854 filed on April 2, 1984.

This invention pertains generally to nickel-base superalloys castable as single crystal articles of manufacture, which articles are especially useful as hot-section components of aircraft gas turbine engines, particularly rotating blades.

The efficiency of gas turbine engines depends significantly on the operating temperature of the various engine components with increased operating temperatures resulting in increased efficiencies. One means by which the operating temperature capability can be increased is by casting the components which operate at the highest temperatures, e.g., turbine blades and vanes, with complex hollow passageways therein so that cooling air can be forced through the component and out through holes in the leading and trailing edges. Thus, internal cooling is achieved by conduction and external cooling is achieved by film or boundary layer cooling.

The search for increased efficiencies has also led to the development of heat-resistant superalloys which can withstand increasingly high temperatures yet maintain their basic material properties. Oftentimes, the development of such superalloys has been done in conjunction with the design, development and manufacture of the aforementioned cast components having intricate air cooling passageways therein.

The present invention is directed to the achievement of increased efficiencies through further improvements in nickel-base superalloys. Accordingly, there is provided by the present invention nickel-base superalloys for producing single crystal articles having a significant increase in temperature capability, based on stress rupture strength and low and high cycle fatigue properties, over single crystal articles made from current production nickel-base superalloys. Further, because of their superior resistance to degradation by cyclic oxidation, and their resistance to hot corrosion, the superalloys of this invention possess a balance in mechanical and environmental properties which is unique and has not heretofore been obtained.

According to the present invention, superalloys suitable for making single-crystal castings comprise the elements shown in Table I below, by weight percent (weight 5), with the balance being nickel (Ni) plus incidental impurities: TABLE I ALLOY COMPOSITIONS (weight %) Most Elements Base Preferred Preferred Chromium (Cr): 5 - 10 6.75 - 7.25 7.0 Cobalt (Co): 5 - 10 7.0 - 8.0 7.5 Molybdenum (Mo3: 0 - 2 1.3 - 1.7 1.5 Tungsten (W) : 3 - 10 4.75 - 5.25 5.0 Tantalum (Ta): 3 - 8 6.3 - 6.7 6.5 Titanium (Ti): O - 2 0.02 max 0.0 Aluminum (Al): 5 - 7 6.1 - 6.3 6.2 Rhenium (Re): O - 6 2.75 - 3.25 3.0 Hafnium (Hf):O - 0.50 0.12 - 0.18 0.15 Carbon (C) : 0 - 0.07 0.04 - 0.06 0.05 Boron (B) : 0 - 0.015 0.003 - 0.005 0.004 Yttrium (Y) : 0 - 0.075 0.005 - 0.030 0.01 The invention also includes cast single-crystal articles, such as gas turbine engine turbine blades and vanes, made of an alloy having a composition falling within the foregoing range of compositions.

There are two basic directional solidification (DS) methods now well-known in the art by which single crystal castings may be made. They generally comprise either the use of a seed crystal or the use of a labyrinthine passage which serves to -select a single crystal of the alloy which grows to form the single crystal article ("choke" process).

In order to develop and test alloys of the invention, three series of 3000 gram heats of the alloys listed in Table II were vacuum induction melted and cast into 1-1/2" dia. copper molds to form ingots.

The ingots were subsequently remelted and cast into 1/2" x 2" x 4" single crystal slabs using the choke process, although the other previously mentioned process could have been used.

In a series of separate experiments, it was determined that yttrium retention in the single crystal slabs was about 30t of that present in the initial ingots. Hence, in preparing the series I, II, and III alloys shown in Table II, sufficient excess yttrium was added to the initially cast alloys so as to achieve the yttrium levels shown in Table II taking into account the 308 retention factor.

The series I alloys were designed to evaluate the interactions between tungsten, molybdenum and rhenium as gamma (Y) matrix alloying elements. The series II alloys were designed somewhat independently from series I in order to accommodate additional variables. Aluminum was maintained at a high level and titanium and tantalum were varied to accomplish a range of gamma prime (Y') levels up to about 63 volume percent and chromium was reduced in order to permit the increased y contents. Since it was determined that the 85 Cr series I alloys as a group were less stable than the series II alloys, the base Cr level was reduced from the 8% in series I to 7t to achieve better stability. Co was varied in alloys 812-814 to evaluate the effect of Co on stability.

The series III alloys were based on evaluations of the series I and II alloys. From series II, the upper limit in .7' content, based on ' solutioning, was about 60 volume percent. Alloys 824-826 were based on alloy 820 which had 5.5t Re and high strength, but was unstable. Thus, the Re content was reduced to achieve stability. Alloys 827-829 were based on alloy 821 (08 Ti), but in which W and Re were varied. Alloys 830-833 were based on alloy 800 (1.5S Ti), but in which Re, W and Mo were varied. Alloys 834 and 835 contained increased Al at the expense of Ta and Ti. In all the series III alloys, the Co content was maintained at 10%, based on the evaluation of alloys 812-814 in series II.

The series I, II and III alloys were evaluated for stress rupture strength and the results of the tests are set forth in Table III. Prior to testing, the alloys, except for the "R" series noted in Table III, were heat treated as 1/2" thick single crystal slabs according to the following schedule: solutionizing at 2350 - 24000F for two hours to achieve solutioning of at least 958 of the ," phase followed by an intermediate age at 19750F for 4 hrs. and a final age at 16500F for 16 hrs. TABLE II (single crystal analyses) Series I Alloy % Cr Co Mo W Ta Ti Al Re Hf B C Y 800 8 7.5 1.5 4.0 5 1.5 5.8 3.0 0.15 0.004 0.05 0 801 8 7.5 0.5 5.9 5 1.5 5.75 3.0 0.15 0.004 0.05 0 802 8 7.5 0.0 5.9 5 1.5 5.75 3.0 0.15 0.004 0.05 0.015 803 8 7.5 0.0 4.0 5 1.5 5.75 4.5 0.15 0.004 0.05 0 804 8 7.5 0.0 2.0 5 1.5 5.75 6.0 0.15 0.004 0.05 0 805 8 7.5 3.65 0.0 5 1.5 5.9 3.1 0.15 0.004 0.05 0 806 8 7.5 3.0 0.0 5 1.5 5.8 4.5 0.15 0.004 0.05 0 807 8 7.5 1.5 3.0 6 1.0 6.0 3.0 0.15 0.004 0.05 0 808 8 7.5 1.5 3.0 6 0.0 6.5 3.0 0.15 0.004 0.05 0.015 809 8 7.5 0.0 4.0 6 0.0 6.4 4.5 0.15 0.004 0.05 0 810 8 7.5 0.0 2.0 6 0.0 6.4 6.0 0.15 0.004 0.05 0 811 8 7.5 3.0 0.0 6 0.0 6.4 4.5 0.15 0.004 0.05 0 TABLE II (single crystal analyses) Series II Alloy % Cr Co Mo W Ta Ti Al Re Hf C B Y 812 7 5.0 1.5 3 6.0 1.0 6.0 3.0 0.15 0.05 0.004 0.015 813 7 7.5 1.5 3 6.0 1.0 6.0 3.0 0.15 0.05 0.004 0.015 814 7 10.0 1.5 3 6.0 1.0 6.0 3.0 0.15 0.05 0.004 0.015 815 5 7.5 1.5 3 7.5 0.5 6.5 3.0 0.15 0.05 0.004 0.015 816 5 7.5 1.5 3 8.0 0.5 6.5 3.0 0.15 0.05 0.004 0.015 817 5 7.5 0.5 3 8.0 1.0 6.5 3.0 0.15 0.05 0.004 0.015 818 5 7.5 1.5 3 7.0 0.5 6.5 3.5 0.15 0.05 0.004 0.015 819 5 7.5 1.5 3 7.0 0.5 6.5 4.5 0.15 0.05 0.004 0.015 820 5 7.5 1.5 3 7.0 0.0 6.5 5.5 0.15 0.05 0.004 0.015 821 7 7.5 1.5 3 6.5 0.0 6.5 3.0 0.15 0.05 0.004 0.015 822 6 7.5 1.5 3 6.5 0.0 6.5 3.0 0.15 0.05 0.004 0.015 823 5 7.5 1.5 3 6.5 0.0 6.5 3.0 0.15 0.05 0.004 0.015 TABLE II (single crystal analyses) Series III Alloy % Cr Co Mo W Ta Ti Al Re Hf C B Y 824 5.0 10.0 1.5 6.5 7.0 0 6.5 2.0 0.15 0.004 0.05 0.015 825 5.0 10.0 1.5 5.5 7.0 0 6.5 3.0 0.15 0.004 0.05 0.015 826 5.0 10.0 1.5 4.0 7.0 0 6.5 4.0 0.15 0.004 0.05 0.015 827 7.0 10.0 1.5 5.0 6.5 0 6.2 3.0 0.15 0.004 0.05 0.015 828 7.0 10.0 1.5 6.0 6.5 0 6.2 2.5 0.15 0.004 0.05 0.015 829 7.0 10.0 1.5 7.0 6.5 0 6.2 2.0 0.15 0.004 0.05 0.015 830 7.0 10.0 1.5 5.0 5.0 1.5 5.8 3.0 0.15 0.004 0.05 0.015 831 7.0 10.0 1.5 6.0 5.0 1.5 5.8 2.0 0.15 0.004 0.05 0.015 832 7.0 10.0 1.5 7.0 5.0 1.5 5.8 1.0 0.15 0.004 0.05 0.015 833 7.0 10.0 1.5 4.0 5.0 1.5 5.8 2.0 0.15 0.004 0.05 0.015 834 6.0 10.0 1.5 5.5 4.0 0 7.0 3.0 0.15 0.004 0.05 0.015 835 7.0 10.0 1.5 4.0 2.0 0 7.5 3.0 0.15 0.004 0.05 0.015 TABLE III Stress Rupture (parallel to single crystal growth direction) ALLOY ACO LIFE (HRS) 1600 F/ 1800 F/ 2000 F/ 2100 F/ 80 Ksi 40 Ksi 20 Ksi 13 Ksi 800 94.5 68.5 184.5 90.4 801 87.0 44.5 45.5 29.1 802 86.8 63.1 90.0 803 66.7 67.1 70.1 82.9 804 54.2 103.0 52.4 805 56.3 56.3 55.2 30.0 806 85.6 43.8 57.4 22.9 807 75.7 60.1 100.3 66.4 808 55.6 53.5 31.6 39.2 809 N 22.0 69.6 46.0 25.3 810 62.2 64.7 35.4 10.1 811 101.9 161.8 44.1 8.2 812 35.1 49.1 30.2 53.5 813 53.8 51.4 27.0 52.2 814 57.9 63.7 42.8 47.1 815 76.5 65.4 64.8 143.2 816 103.9 83.6 47.6 121.4 817 24.8 55.5 42.3 55.5 818 N 84.3 85.6 68.0 113.3 819 147.4 115.5 100.6 264.3 820 257.2 158.7 153.7 220.1 821 114.3 80.4 98.4 74.3 822 64.2 70.3 43.1 96.9 823 22.3 48.7 - 46.3 TABLE III - continued Stress Rupture ALLOY ACO1 LIFE (HRS) 16000F/ 18000F/ 20000F1 80 Ksi 40 Ksi 20 Ksi 824 97.1 91.9 R 118.8 67.7 825 74.0 94.7 R 128.4 82.3 826 113.1 119.8 R 135.6 122.1 827 N 6.7 76.2 R 108.4 70.0 828 119.1 72.7 R 95.7 119.0 829 110 72.7 R 90.2 88.0 830 -N 59.8 126.5 R 162.2 147.8 831 92.7 68.9 R 82.1 137.6 832 90.1 58.5 R 85.4 107.3 833 96.5 51.6 R 69.1 132.7 834 119.2 80.7 R 105.7 69.1 835 N 43.5 55.3 R 67.4 27.9 1 ACO- Acceptable Crystallographic Orientation single crystal growth orientation within 15 of the tO01] zone axis; N r no, otherwise yes The series III alloys were initially tested at 16000F/80 ksi and 1800"F/40 ksi. Based on other tests, such as those reported in Table VII, additional test specimens were resolutioned at 23900F for two hours, given a more rapid cool and aged at 2050F/4 hours + 16500F/4 hours, the "R" treatment listed in Table III, and stress rupture tested at 1800 F/40 ksi and 20000F/20 ksi. The reheat treatment resulted in an average increase in rupture life at 1800 F/40 ksi of about 30%.At the critical parameter of 1800F/40 ksi for gas turbine engine applications, it is expected that the series I and II alloys would also exhibit a 305 increase in life when given the "R" treatment.

Other experiments have shown that cooling rates from the solutionizing temperature to 2000 F in the range of 100-6000F/min have only a slight effect on the stress rupture properties of the alloys of the invention with higher rates tending to improve the life at 18000F/40 ksi slightly. The data are shown in Table IV.

TABLE IV Cooling Rate Stress Rupture Life, Hours F/Min 18000F/40 ksi 2000 F/20 ksi 600 91 107 300 85 123 100-150 75 120 Average of all - 79 100 prior data (vari ous cooling rates) Thus, for the superalloys of the invention, the presently preferred heat treatment is as follows: solutionize in a temperature range sufficient to achieve solutioning of at least 95% of the y' phase, preferably 2385-23950F, for 2 hrs., cool to 20000F at 1000F/min. minimum, furnace cool to 12000F in 60 min.

or less and thereafter cool to room temperature; heat to 2050 + 250F for 4 hrs., furnace cool to below 12000F in 6 min. or less and thereafter to room temperature; and heat to 1650 + 250F for 4 hrs. and thereafter furnace cool to room temperature. All heat treatment steps are performed in vacuum or an inert atmosphere, and in lieu of the steps calling for cooling to room temperature the treatment may proceed directly to the next heating step.

The stress rupture data from the series I, II, and III alloys indicates that about 5t Re provides the highest rupture strength at 18000F/40 ksi. The data also show, when rupture life is graphed as a function of rhenium content at constant tungsten contents, that high rupture life at 1800 F/40 ksi can be obtained with rhenium plus tungsten levels in the (3Re + 7W) to (SRe + 3W) ranges. In the most preferred embodiment, Alloy 821, the presently preferred (Re + W) combination is (3Re + SW) due to the present relative costs of rhenium and tungsten.

All the alloys were evaluated for microstructural stability. Specimens were heat treated by solutionizing at 2375-2400 F/2 hrs. and aging at l9750F/4 hrs. and at 16500F/16 hrs. Thereafter, different sets of specimens were heated for 1000 hrs.

at 1800"F and for 1000: hrs. at 20000F. After preparation, including etching with diluted Murakami's electrolyte, the specimens were examined metallographically and the relative amount of topologically close packed phase (TCP) was determined visually. The series II alloys, except alloys 818 and 819, showed either no TCP precipitation or only traces of precipitation (821) and, as a group, were less prone to microstructural instability than the series III alloys and much less prone than the series I alloys at both 18000F and 20000F.

Table V presents the results of cyclic oxidation tests on uncoated 1/4" dia. x 3",long pin specimens conducted at 21SOOF using a natural gas flame at Mach 1 gas velocity. The specimens were rotated for uniform exposure and cycled out of the flame once per hour to cool the specimens to room temperature.

External petal loss was measured on a section cut transverse to the length dimension of the specimen.

Metal loss per side was found by dividing the difference between the pin diameter before and after test by two. The data in the table are the average of two such measurements at 90 to each other across the diameter of the specimen.

The two series 1 alloys that contained yttrium (802 and 808) had exceptional oxidation resistance.

The series II alloys, all of which were yttrium-bearing, exhibited no metal loss after 200 hours of high velocity oxidation (Mach I) at 215O0F and only 2-3 oils y' depletion, demonstrating that a synergistic Y + Hf effect was operating. These data also demonstrate that Re improves the oxidation resistance or at least is less detrimental than W which it has replaced in the alloys and, from metallographic studies, also results in improved 7' stability.

TABLE V Oxidation, 21500F, Mach 1.0 ,,' Depletion Metal Loss (mils/side) (mils/side) Time (hrs) 18.6 62 6 127.6 169.6 214.6 800 1.0 1.8 4.8 6.8 9.3 6-12 801 0.8 1.5 4.8 8.0 9.8 4-8 802 0 0 0.3 0 0.3 2-3 803 0.8 1.5 2.8 4.0 5.8 8-10 804 0.8 1.3 1.5 4.0 5.3 8-12 805 . 1.0 1.5 10.3 7.5 9.5 8-10 806 0.8 1.8 4.8 6.3 9.8 8-10 807 1.0 2.0 4.5 6.3 8.0 6-8 808 0.3 0 0.3 0.3 0.5 2-3 809 1.0 1.8 2.8 2.0 4.3 10-14 810 1.0 1.8 2.8 3.5 4.0 10-16 811 1.3 2.5 3.0 4.5 5.5 12-16 812 0 0 0 0 0 1-2 813 0 0 0 0 0 1-2 814 0 0 0 0 0 1-2 815 0 0 0 0 0 1-2 816 0 0 0 0 0 1 817 0 0 0 0 0 1 818 0 0 0 0 0 2 819 0 0 0 0 0 2 820 0 0 0 0 0 2 821 0 0 0 O 0 1-2 822 0 0 0 0 0 1-2 823 0 0 0 0 0 1-2 R125 - - - - 80 R80 - - - - 90 MA754 - - - - 12 The hot corrosion resistance of the alloys of the invention was evaluated alongside three alloys used to produce production turbine blades, Rene' 125, B1900, and MM200(Hf), in tests wherein specimens of the alloys were exposed to a JP-5 fuel-fired flame at 1600eF with a nominal 1 ppm salt added to the combustion products.

The test was first run at -1 ppm for 1040 hrs., and then at -2 ppm, for 578 hrs. The chemical determination of NaCl on calibration pins at every 200 hours indicated that the salt level was between 0 and 1 ppm during the first 1000 hours, between 1 and 2 ppm during the next 300-400 hours and above 2 ppm during the remaining 300 hours.The following conclusions were drawn from these hot corrosion tests: 1) B1900 was least resistant to hot corrosion at all salt levels, 2) MM200(Hf) was the next least resistant alloy at all salt levels, 3) the alloys of the invention, especially alloy 821, and Rene' 125 exhibit similar hot corrosion behavior, with the alloys of the invention being slightly less resistant than Rene 125, and 4) as is the case for Rene' 125 and other alloys, the alloys of the invention appear to be sensitive to salt level in the corrosion test with increased salt level resulting in poorer corrosion resistance. Thus, the difference between B1900, MM200(Hf), Rene' 125,;ant the alloys of the invention narrows at high salt levels.These results are consistent with prior experience and indicate that the hot corrosion resistance of the alloys of the invention will be adequate for applications where Rene' 125 equivalency is required.

Alloy 821 was scaled up as a 300 lb master heat having the composition given in Table VI. No yttrium was added to the master heat; rather, yttrium was added when the master heat material was remelted and molten prior to DS'ing to produce single crystal slabs and turbine blades. For the test specimens used to obtain the data of Tables VII, VIII, IX, and X, yttrium in the amount of 400 ppm was added. Stress rupture strength data for alloy 821 from the 300 lb master heat and the 12 ib. laboratory heat are presented in Table IX.

TABLE VI 300 Lb Alloy 821 Master Heat Cr 6.79 Ti O Co 7.30 Re 2.95 Mo 1.48 Hf 0.17 W 4.95 C 0.05 Ta 6.40 B 0.004 Al 6.15 Y O TABLE VII Stress Rupture Data T Stress Life El RA Heat H. Treat ( F (ksi) CHrsJ (t C%) 12 ib 2390/2 + 1975/4 1600 80 114.3 Lab.Ht. + 1650/16 1800 40 80.4 H.T. as slabs 2000 20 98.4 7.7 43.1 2100 13 74.3 16.8 6.8 300 lb 2390/2 + 1975/4 1400 130 1.9 19.5 26.9 Master Ht. + 1650/16 1400 110 351.6 14.8 24.4 Alloy 821 H.T. as slabs 1600 80 155.4 20.1 26.8 1800 40 72.7 39.4 29.9 1800 40 75.8 20.6 33.2 1800 35 227.8 17.5 27.3 1800 30 509.2 16.8 28.7 1900 25 120.2 10.1 23.4 1900 22 357.2 13.9 28.6 2000 20 81.3 13.6 38.5 2000 17.5 391.9 13.1 23.3 2100 13 80.5 3.4 48.6 300 lb Reheat treated* 1600 80 115.8 19.0 25.0 Alloy 821 + 1900/4 age + 1800 40 68.4 17.0 30.5 1650/4 age 2000 20 82.7 13.9 35.2 Reheat treated* 1600 80 155.2 19.0 26.2 + 1975/4 age + 1800 40 85.2 25.5 39.0 1650/4 age 2000 20 101.2 14.7 34.4 Reheat treated* 1600 80 160.0 18.9 27.5 + 2050/4 age + @ 1800 40 103.8 18.1 28.3 1650/4 age @ 2000 20 125.7 11.6 40.3 Reheat treated* 1600 80 139.9 19.3 24.0 + 2125/1 age + 1800 40 97.4 23.2 28.6 + 1975/4 (coating 2000 20 126.9 12.8 32.9 simulation) + 1650/4 age Reheat treated* 1600 80 131.0 17.8 24.7 + 2200/1 age + 1800 40 90.5 20.6 29.8 + 1975/4 (coating 2000 20 97.2 12.4 31.1 simulation) + 1650/4 age * All resolutioned in test specimen form at 23900F/2 hr + fast cool to 20000 F.

Tensile strength, low cycle fatigue and high cycle fatigue tests were performed on single crystal material from the 300 lb heat of alloy 821 solutioned at 2390 F/2 hrs. and aged at 19750F/4 hrs. and 16500F/16 hrs., with the results shown in Tables VIII, IX, and X, respectively, where UTS is ultimate tensile strength; YS is yield strength at 0.28 strain offset; El is elongation; and RA is reduction in area.

TABLE VIII Tensile Data (Master Heat Alloy 821) Temp UTS 0.2t YS 0.02t YS El RA (OF) (Ksi) (Ksi) (Ksi) (%) (%) 1000 128.6 113.4 110.7 11.6 18.9 1200 129.6 112.4 106.5 14.2 19.9 1400 142.8 112.8 102.6 9.9 13.3 1600 143.3 129.4 103.5 18.0 30.8 1800 110.1 94.7 71.9 10.0 28.1 2000 64.1 51.2 39.2 19.1 21.6 TABLE IX Low Cycle Fatigue (Master Heat Alloy 821) Alternating Pseudostress Cycles to Failure (ksi)@ Nf 21 4.9 x 103 31 2.3 x 103 37 2.5 x 103 1 2 min. compressive strain hold, 2000 F TABLE X High Cycle Fatiguel (Master Heat Alloy 821) Alternating Stress Cycles to Failure (ksi) Nf 10 9.6 x 106 11 4.4 x 106 13 1.4 x 106 15 0.5 x 106 1 20500F; A r 0.67, 30 Hz As discussed at greater length in co-pending co-assigned application Ser. No. 595,854, the superalloys of this invention break with the long-standing wisdom of the single crystal superalloy arts that grain boundary strengthening elements such as B, Zr and C are to be avoided, i.e., kept to the lowest levels possible consistent with commercial melting and alloying practice and technology.One general reason given for restricting such elements is to increase the incipient melting temperature in relation to the y' solvus temperature thus permitting solutionizing heat treatments to be performed at temperatures where complete solutionizing of the y' phase is possible in reasonable times without causing localized melting of solute-rich regions. Another is to minimize or preclude the formation of deleterious TCP phases.

As noted in the Ser. No. 595,854 application, single crystal articles are not necessarily wholly of a single crystal as there may be present therein grain boundaries referred to as low angle grain boundaries wherein the crystallographic mismatch across the boundary is generally accepted to be less than about 5 to 6 degrees. Low angle grain boundaries are to be distinguished from high angle grain boundaries which are generally regarded as boundaries between adjacent grains whose crystallographic orientation differs by more than about 5-6 degrees. High angle grain boundaries are regions of high surface energy, i.e., on the order of several hundreds of ergs/cm2, and of such high random misfit that the structure cannot easily be described or modelled.

As also noted therein, the discovery that small, but controlled, amounts of such previously prohibited elements can be tolerated resulted in the single crystal superalloys of the Ser. No. 595,854 application which have improved tolerance to low angle grain boundaries, i.e., have greater grain boundary strength than the state-of-the-art single crystal superalloys. As one result of this increased grain boundary strength, grain boundary mismatches far greater than the 60 limit for prior art single crystal superalloy articles can be tolerated in single crystal articles made from the nickel-base superalloys of that invention. This translates, for example, into better in-service reliability, lower inspection costs and higher yields as,grain boundaries over a broader range can be accepted by the usual inspection techniques.

The novel features of that invention have been embodied in the novel superalloys of the present invention; thus, the superalloys of the present invention also exhibit improved tolerance to low angle grain boundaries and also have the above-described benefits.

The superalloys of this invention are also alloyed with yttrium which renders them more highly reactive with respect to ceramic molds and cores used in the investment casting process than nickel-base superalloys not alloyed with yttrium. Ceramic/metal instability is controlled by the bulk thermodynamic condition of the system. The more negative the free energy of formation, r?rGOf, the greater the affinity for oxygen. It has been found that the free energy of formation for oxides becomes more negative as more reactive elements, such as yttrium, are added resulting in a greater potential for metal/ceramic reaction than when typical Si02 and ZrO2 ceramic mold and core systems are used. Based on thermodynamic considerations and the work reported in U.S.Department of the Air Force publication AFML-TR-77-211, "Development of Advanced Core and Mold Materials for Directional Solidification of Eutectics" (1977), alumina is less reactive and is, therefore, a preferred material for molds, cores and face coats when casting superalloys containing reactive elements.

It has also been found that melt/mold and core interactions are decreased, the retention of yttrium increased and the uniformity of yttrium distribution improved by the use of low investment casting parameters and temperatures. This translates to the use of the lowest possible superheat and mold preheat and a high withdrawal rate in the casting of the single crystal articles of this invention Several uncored small turbine blades were investment cast using alloy 821 material from the previously mentioned 300 lb scale-up master heat.

Those blades measured about 1.5" from tip to root with a span of approximately 0.75". Blade tip to platform distance was 1". As noted earlier, yttrium was added to the master heat material while molten and prior to DS'ing -- in this case the amount was 2000 ppm. In general, most blades exhibited acceptable crystal structure and, as shown in Table XI, those cast using low casting parameters had better yttrium retention.

Also, it appeared that surface to volume ratio Influences yttrium retention; as the ratio increases, the yttrium retention decreases. This is illustrated by comparison of yttrium retention at the leading and trailing edges; the surface to volume ratio is lower in the leading edge compared to the trailing edge, and the yttrium retention in the leading edge is consistently higher than at the trailing edge.

TABLE XI Yttrium Content (ppm) Airfoil Tip Airfoil Near Platform Blade Root Casting Condition LE(1) TE (1) LE TB ROOT@ Low 130 100 160 100 130 Superheat 90 60 80 50 160 190 120 190 150 190 170 90 180 150 200 410 330 470 360 380 310 120 270 160 280 High 80 60 120 70 100 Superheat 80 80 100 70 130 100 90 90 150 100 80 60 100 100 100 130 150 190 150 120 170 200 240 210 170 1 LE " leading edge 2 TE s trailing edge 3 ROOT = rqot, center Additional single crystal investment castings of large solid turbine blades (4-3/4" tip-to-root) and small and large turbine blades having cores therein to define serpentine passageways for the provision of cooling air were also made. The large solid turbine blades required late yttrium additions of up to 2400 ppm in order to obtain yttrium distributions within the desired 50-300 ppm level. Similar such levels, coupled with the use of low investment casting parameters, were required to obtain acceptable yttrium levels in the cored blades. As was the case with the uncored small turbine blades, the effect of surface to volume ratio was evident; the leading edge retained higher yttrium levels compared to the trailing edge.

Although the present invention has been described in connection with specific examples, it will be understood by those skilled in the art that the present invention is capable of variations and modifications within the scope of the invention as represented by the appended claims.

Claims

CLAIMS What is claimed is:

1. A nickel-base superalloy consisting essentially of, in percentages by weight, 5-10 Cr, 5-10 Co, 0-2 Mo, 3-10 W, 3-8 Ta, 0-2 Ti, 5-7 Al, 0-6 Re, 0-0.5 Hf, or0.0? C, O-0.015 B, and 0-0.075 Y, the balance being nickel and incidental impurities.

2. The superalloy of claim 1 consisting essentially of, in percentages by weight, 6.75-7.25 Cr, 7.0-8.0 Co, 1.3-1.7 Mo, 4.75-5.25 W, 6.3-6.7 Ta, 0.02 max. Ti, 6.1-6.3 Al, 2.75-3.25 Re, 0.12-0.18 Hf, 0.04-0.06 C, 0.003-0.005 B, and 0.005-0.030 Y, the balance being nickel and incidental impurities.

3. The superalloy of claim 2 consisting essentially of, in percentages by weight, 7 Cr, 7.5 Co, 1.5 Mo, 5 W, 6.5 Ta, O Ti, 6.2 Al, 3 Re, 0.15 Hf, 0.05 C, 0.004 B, and 0.01 Y, the balance being nickel and incidental impurities.

4. A nickel-base superalloy for producing single crystal articles consisting essentially of, in percentages by weight, 5-10 Cr, 5-10 Co, 0-2 Mo, 3-10 W, 3-8 Ta, 0-2 Ti, 5-7 A1, 0-6 Re, 0-0.5 Hf, 0-0.07 C, 0-0.015 B, and 0-0.075 Y, the balance being nickel and incidental impurities.

5. The superalloy of claim 4 consisting essentially of, in percentages by weight, 6.75-7.25 Cr, 7.0-8.0 Co, 1.3-1.7 Mo, 4.75-5.25 W, 6.3-6.7 Ta, 0.02 max Ti, 6.1-6.3 Al, 2.75-3.25 Re, 0.12-0.18 Hf, 0.04-0.06 C, 0.003-0.005 B, and 0.005-0.030 Y, the balance being nickel and incidental impurities.

6. The superalloy of claim 5 consisting essentially of, in percentages by weight, 7 Cr, 7.5 Co, 1.5 Mo, 5 W, 6.5 Ta, O Ti, 6.2 Al, 3 Re, 0.15 of, 0.05 C, 0.004 B, and 0.01 Y, the balance being nickel and incidental impurities.

7. The article of claim 4 which is an airfoil member for a gas turbine engine.

8. A nickel-base superalloy for producing single crystal articles having improved tolerance to low angle grain boundaries consisting essentially of, in percentage by weight, 5-10 Cr, 5-10 Co, 0-2 Mo, 3-10 W,3-8 Ta, 0-2 Ti, 5-7 Al, 0-6 Re, 0-0.5 Hf, 0-0.07 C, 0-0.015 B, and 0-0.075 Y, the balance being nickel and incidental impurities.

9. The heat treatment of claim 6 wherein the temperature range in b) is from about 2385 to 2395 -F.

9. The superalloy of claim 8 consisting essentially of, in percentages by weight, 6.75-7.25 Cr, 7.0-8.0 Co, 1.3-1.7 Mo, 4.75-5.25 W, 6.3-6.7 Ta, 0.02 max. Ti, 6.1-6.3 Al, 2.75-3.25 Re, 0.12-0.18 Hf, 0.04-0.06 C, 0.003-0.005 B, and 0.005-0.030 Y, the balance being nickel and incidental impurities.

10. The superalloy of claim 9 consisting essentially of, in percentages by weight, 7 Cr, 7.5 Co, 1.5 Mo, 5 W, 6.5 Ta, 0 Ti, 6.2 Al, 3 Re, 0.15 Hf, 0.05 C, 0.004 B, and 0.01 Y, the balance being nickel and incidental impurities.

11. The article of claim 8 which is an airfoil member for a gas turbine engine.

12. A heat treatment conducted in a vacuum or inert atmosphere comprising the steps of: a) selecting a single crystal article having a composition consisting essentially of, in percentages by weight, 5-10 Cr, 5-10 Co, 0-2 Mo, 3-10 W, 3-8 Ta, 0-2 Ti, 5-7 Al, 0-6 Re, 0-0.5 Hf, 0-0.07 C, 0-0.015 B, and 0-0.075 Y, the balance being nickel and incidental impurities; b) heating for 2 hours in a temperature range sufficient to achieve solutioning of at least about 95% of the y' phase; c) cooling to 20000F at 1000F/minute minimum and then to 12000F in 60 minutes or less; d) heating to 2050 + 25 F for 4 hours; e) cooling to below 1200 F in 6 minutes or less; and f) heating to 1650 + 250F for 4 hours.

13. The heat treatment of claim 12 wherein the composition of the article consists essentially of, in percentages by weight, 6.75-7.25 Cr, 7.0-8.0 Co; 1.3-1.7 Mo, 4.75-5.25 W, 6.3-6.7 Ta, 0.0Z max. Ti, 6.1-6.3 Al, 2.75-3.25 Re, 0.12-0.18 Hf, 0.04-0.06 C, 0.003-0.005 B, and 0.005-0.030 Y, the balance being nickel and incidental imp@rities.

14. The heat treatment of claim 12 wherein the composition of the article consists essentially of, in percentages by weight; 7 Cr, 7.5 Co, 1.5 Mo, 5 W, 6.5 Ta, O Ti, 6.2 Al, 3 Re, 0.15 Hf, 0.05 C, 0.004 B, and 0.01 Y, the balance being nickel and incidental impurities.

15. The heat treatment of claim 12 wherein the temperature range in ti, is from about 2385 to 2395 F.

Amendments to the claims have been filed as follows 1. A nickel-base superalloy consisting essentially of, in percentages by weight, 5-10 Cr, 5-lOCo, 0-2 Mo, 3-10 W, 38 Ta, 0-2 Ti, 5-7 Al, 0-6 Re, 0-0.5 Hf, 0-0.07 C, 0-0.015 B, and 0-0.075 Y, the balance being nickel and incidental impurities.

4. A single crystal article produced from a nickel-base superalloy having a composition according to any one of the preceding claims.

5. An article according to claim 4 which is an airfoil member for a gas turbine engine.

6. A heat treatment conducted in a vacuum or inert atmosphere comprising the steps of: (a) selecting a single crystal article having a composition consisting essentially of, in percentages by weight, 5-10 Cr, 5-10 Co, O 2 Mo, 3-10 W, 3-8 Ta, 0-2 Ti, 5-7 Al, 0-6 Re, 0-0.5 Hf, 0-0.07 C, 0-0.015 B, and 0 0.075 y, the balance being nickel and incidental impurities; b) heating for 2 hours in a temperature range sufficient to achieve solutioning of at least 95t of the y' phase; c) cooling to 20000F at 100 F/minute minimum and then to 12000F in 60 minutes or less; d) heating to 2050 + 25 F for 4 hours; e) cooling to below 12000F in 6 minutes or less; and fj heating to 1650 + 25 F for 4 hours.

7. The heat treatment of claim 6 wherein the composition of the article consists essentially of, in percentages by weight, 6.75-7.25 Cr, 7.0-8.0 Co, 1.3-1.7 Mo, 4.75-5.25 W, 6.3-6.7 Ta, 0.02 max. Ti, 6.1-6.3 Al, 2.75-3.25 Re, 0.12-0.18 Hf, 0.04-0.06 C, 0.003-0.005 B, and 0.005-0.030 Y, the balance being nickel and incidental impürities.

8. The heat treatment cf claim 6 wherein the composition of the article consists essentially of, in percentages by weight , 7 Cr, 7.5 Co, 1.5 Mo, 5 W, 6.5 Ta, 0 Ti, 6.Z Awl, 3 Re, 0.15 Hf, 0.05 C, 0.004 B,and 0.01 Y, the balance being nickel and incidental impurities.