AU621149B2

AU621149B2 - Improvements in or relating to alloys

Info

Publication number: AU621149B2
Application number: AU41639/89A
Authority: AU
Inventors: Leo Buchakjian Jr.; Carl Stephen Wukusick
Original assignee: General Electric Co
Current assignee: General Electric Co
Priority date: 1986-12-30
Filing date: 1989-09-20
Publication date: 1992-03-05
Anticipated expiration: 2009-09-20
Also published as: GB2235697B; CA1327132C; FR2654114A1; FR2654114B1; GB8631037D0; AU4163989A; GB2235697A

Description

2 9 307 LIn ohm <0 1,> 621149 7~~ Our Ref: 293910

AUSTRALIA

Patents Act COMPLETE SPECIFICATION FORM

(ORIGINAL)

Application Number: Lodged: Complete Specification Lodged: Accepted: Published: a00 0°0 0 o0 0 0 00 a o o 0000 0 00 0 0f 0 00*0 0 &0 0 0 Priority: Related Art: Applicant(s): GENERAL ELECTRIC COMPANY 1 River Road, Schenectady, New York, U.S.A.

ARTHUR S. CAVE CO.

Patent Trade Mark Attornerys Level 10, 10 Barrack Street SYDNEY NSW 2000 Address for Service: O 00 «o a oe f9 o 0° 0 Complete specification for the invention entitled 0"Improvements in or relating to alloys".

"Improvements in or relating to alloys".

@00 0080 00 0 0 0 O 00 The following statement is a full description best method of performing it known to me:of this invention, including the 1 3075U/gs I( j -j 7- i I'l; 3063U:AB r I L~ a

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t -lao This invention pertains generally to nickel-base superalloys castable as single crystal narticles 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 So 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, turbine blades and Svanes, with complex hollow passageways therein so that SL 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, I- I r ir L 13DV-8623 2the 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-bese 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 o QO resistance to degradation by cyclic oxidation, and their resistance to hot corrosion, the superalloys of this invention possess a balance in mechanical and oo 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 o i percent (weight with the balance being nickel (Ni) Splus incidental impurities: 10 hvin a sgniican inreas intempratre 13DV-8623 -3- TABLE I ALLOY COMPOSITIONS (weight a 0 a 0 Most Elements Base Preferred Preferred Chromium 5 10 6.75 7.25 Cobalt 5 10 7.0 8.0. Molybdenum 0 2 1.3 1.7 Tungsten 3 10 4.75 5.25 Tantalum 3 8 6.3 6.7 Titanium 0 2 0.02 max 0.0 Aluminum S 7 6.1 6.3 6.2 Rhenium 0 6 2.75 3.25 Hafnium 0 0.50 0.12 0.18 0.15 Carbon 0 0.07 0.04 0.06 0'.05 Boron 0 0.015 0.003 0.005 0.004 Yttrium 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 fallingV 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"1 dia. copper molds to form ingots.

The ingots were subsequently remelted and cast into 1/2"1 x 2 g1 x 4"1 single crystal slabs using the choke process, although the otheT previously mentioned process could have been used.

13DV-8623 -4- In a series of separate experiments, it was determined that yttrium retention in the single crystal slabs was about 30% 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 30% retention factor.

The series I alloys were designed to evaluate the interactions between tungsten, molybdenum and rhenium as gamma matrix alloying elements. The series II alloys were designed somewhat independently 6 from series I in order to accommodate additional Svariables. Aluminum was maintained at a high level and o 4o a 15 titanium and tantalum were varied to accomplish a range SS of gamma prime 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 8% 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 7% to achieve better o 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 y' content, based on 7' solutioning, was about 60 volume percent. Alloys 824-826 were based on alloy 820 which had 5.5% Re and high strength, but was unstable. Thus, the Re content was reduced to achieve stability. Alloys 827-829 were based on alloy 821 Ti), but in which W and Re were varied. Alloys 830-833 were based on alloy 800 Ti), but in which Re, W and Me 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 i 4.

j 13DV-8623 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, excc,-pt for the "R11 series noted in Table III, were heat treated as 1/211 thick single crystal slabs according to the following schedule: solutionizing at 2350 2400 0 F for two nours to ach 'ieve solutioning of at least 95% of the y' phase followed by an intermediate age at 1975*F for 4 hrs. and a final age at 1650*F for 16 hrs.

#tit r 4 4 0640 0060 a a, a a, a a S a to a a tot a a o -s a 0 CC 0 TABLE II (single crystal analyses) Series I Alloy f Cr Co Mo W la Ti Al Re Hf B Y 800 801 802 803 804 805 806 807 808 809 810 811 7.5 7.5 7.5 7.5 7.5 7.5 7.5 7.5 7.5 7.5 7.5 7.5 1.5 4.0 0.5 5.9 0.0 5.9- 0.0 4.0 0.0 2.0 3.65 .0.0 3.0 0.0 1.5 3.0 1.5 3.0 0.0 4.0 0.0 2.0 3.0 0.0 1.5 1.5 1.5 1.5 1.5 1.5 1.5 1.0 0.0 0.0 0.0 0.0 5.8 5.75 5.75 5.75 5.75 5.8 6.0 6.5 6.4 6.4 6.4 3.0 3.0 3.0 4.5 6.0 3.1 4.5 3.0 3.0 4.5 6.0 4.5 0.15 0.15 0.15 0.15 0.15 0.15 0.15 0.15 0.15 0.15 0.15 0.15 0.004 0.004 0.004 0.004 0.004 0.004 0.004 0.004 0.004 0.004 0.004 0.004 0.05 0.05 0.05 0.05 0.05 0.05 0.05 0.05 0.05 0.05 0.05 0.05 0 0 0.015 0 0 0 0 0 0.015 0 0 0 *1 tj Ip 8 9.89 09 *8 9 *9 8 9 8 0 CP 4 00 690 9 8 8 8~ 92 8 0 0 8 ~88 8 0 8 8 8 8 09 8 89 8 C TABLE 11 (single crystal analyses) Series II Alloy 812 813 814.

815 816 817 818 819 820 821 822 823 Cr 7 7 7 5 5.

5 5 5 5 7 6 5 Co 5.0 7.5 10.0 7.5 7.5 7.5 7. 5 7.5 7.5 7.5 7.5 7.5 Mo W Ta Ti 1.5 3 6.0 1.0 1.5 3 6.0 1.0 1.5 3 6.0 1.0 1.5 3 7.5 0.5 1.5 3 8.0 0.5 0.5 3 8.0 1.0 1.5 3 7.0 0.5 1.5 3 7.0 0.5 1.5 3 7.0 0.0 1.5 5 6.5 0.0 1.5 5 6.5 0.0 1.5 5 6.5 0.0 Al 6.0 6.0 6.0 6.5 6.5 6.5 6.5 6.5 6.5 6.2 6.2 6.2 Re Hf C B Y 3.0 0.15 0.05 0.004 0.015 3.0 3.0 3.0 3.0 3.0 3.5 4.5 5.5 3.0 3 0 3.0 0.15 0.15 0.15 0.15 0.15 0.15 0.15 0.15 0.15 0.15 0.15 0.05 0.05 0.05 0.05 0.05 0.05 0.*05 0.05 0.05 0.05 0.05 0.004 0.004 0.004 0.004 0.004 0.004 0.004 0.004 0.004 0.004 0.004 0.015 0.015 0.015 0.*015 0.015 0.015 0.015 0.015 0.015 0.015 0.015

I-,

k IW- r 0 0 0 00 0 0 00 000 a c-c o 0 0 0 0 0 00 0 0 0 00* 0 0 0 0 00 0 00 0 0 0 000 0 6 00 TABLE II (single crystal analyses) Series III Alloy f Cr Co Mo W Ta Ti SA1 Re Hf B C Y 824 825 826 827 828 829 830 831 832 833 834 835 5.0 5.0 5.0 7.0 7.0 7.0 7.0 7.0 7.0 7.0 6.0 7.0 10.0 10.0 10.0 10.0 10.0 10.0 10.0 10.0 10.0 10.0 10.0 10.0 1.5 6.5 1.5 5.5 1.5. .4.0 1.5 5.0 1.5 6.0 1.5 7.0 1.5 5.0 1.5 6.0 1.5 7.0 2.5 4.0 1.5 5.5 1.5 4.0 7.0 7.0 7.0 6.5 6.5 6.5 5.0 5.0 5.0 5.0 4.0 2.0 0 0 0.

0 0 0 1.5 1.5 1.5 1.5 0 0 6.5 6.5 6.5 6.2 6.2 6.2 5.8 5.8 5.8 5.8 7.0 7.5 2.0 3.0 4.0 3.0 2.5 2.0 3.0 2.0 1.0 2.0 3.0 3.0 0.15 0.15 0.15 0.15 0.15 0.15 0.15 0.15 0.15 0.15 0.15 0.15 0.004 0.004 0.004 0.004 0.004 0.004 0.004 0.004 0.004 0.004 0.004 0.004 0.05 0.05 0.05 0.05 0.05 0.05 0.05 0.05 0.05 0.05 0.05 0.05 0.015 0.015 0.015 0.015 0.015 0.015 0.015 0.015 0.015 0.015 0.015 0.015

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13DV- 8623 TABLE III stress Ruptur' (parallel to single crystal growzti direction) ALLOY AC0 1 LI'FE (HRS) 160OFT Ksi t'000 0 o eQ 4~0 0 o 0~' 0e C 0 1, C 0~~Q~e *00~ 0 00 @0 0- It I I I LI It I I 800 801 802 803 804 805 806 807 808 15 809 810 811 812 813 20 814 815 816 817 818 819 820 821 822 823 94.5 87.0 86.8 66.7 54.2 56.3 85.6 75.7 55.6 N 22.0 62.2 101.9 35.1 53.8 57.9 76.5 103.9 24.8 N 84.3 147. 4* 257.2 114.3 64. 2 22.3 180 0 "F/ 40 Ksi 68.5 44.5 63.1 67.1 103.0 56.3 43.8 60.1 53,5 69.6 64.7 161.8 49.1 51.4 63.7 65.4 83.6 55.5 85.6 115.5 158.7 80.4 70.3 48.7 184.5 45.5 90.0 70.1 52.4 55.2 57.4 100.3 31.6 46.0 35.4 44.1 30.2 27.0 42.8 64.8 47.6 42.3 68.0 100.6 153.7 98.4 43.1 29.1 82.9 30.0 22.9 66.4 39.2 25.3 10.1 8.2 53.5 52.2 47.1 143.2 121.4 55.5 113.3 264.3 Z20. 1 74.3 96.9 46.3 (continued) 2000 0 F/ 21000F/ 20 Ksi 13 Ksi

LI

I I I

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13DV-8623 TABLE III continued Stress Rupture ALLOY ACO 1 LIFE (HRS) 1600OF Ksi 97.1 824 82S 826

R

827 74.0 113.1 6.7 119.1 828 04 o 0 4 829 110 180 0

FI

40 Ksi 91.9 118.8 94.7 128.4 119.8 135.6 76.2 108.4 72.7 95.7 72.7 90.2 126.5 162.2 68.9 82.1 58.5 85.4 51.6 69. 1 80.7 105.7 55.3 67.4 67.7 82.3 122.1 70.0 119.0 20 000FI 20 IKsi 88.0 830 20 R 832 59.8 92.7 90.1 96.5 119.2 43.5 147.8 137.6 107.3 132.7 833 834 835 69.1 27.9 4 a4 4 41 *4 0 1ACOn Acceptable Crystallographic Orientation 30 single crystal growth orientation within 15* of the [001] zone axis; N m no, otherwise yes The series III alloys were initially tested at 1600*F/80 ksi and 1800OF/40 ksi.. Based on other tests, such as those reported in Table VII, additional test specimens were resolutioned. at 2390*F for two hours, given a more rapid cool and aged at 20500F/4 hours 1650OF/4 hours, the "R11 treatment listed in Table III, vsm.

^M

1

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13DV-8623 -11and stress rupture tested at 1800°F/40 ksi and 2000°F/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 1800*F/40 ksi for gas turbine engine applications, it is expected that the series I and II alloys would also exhibit a increase in life when given the treatment.

Other experiments have shown that cooling rates from the solutionizing temperature to 2000F in the range of 100-600*F/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 1800°F/40 ksi slightly. The data are shown in Table IV.

0000 o o o 00 00 6" 0 0 00 0000 9 06 0 TABLE IV Cooling Rate °F/Min Stress Rupture Life, Hours 1800°F/40 ksi 2000*F/20 ksi 0000 0 0 00 0 0 00 0 8 0 0 00 0000 0 0 0 00 0 600 300 20 100-150 Average of all prior data (various cooling rates) 91 85 75 79 107 123 120 100 eo eo 0 0 0 0 0 0 00 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 95t of the y' phase, preferably 2385-2395°F, for 2 hrs., cool to 2000"F at 100°F/min. minimum, furnace cool to 1200*F in 60 min.

or less and thereafter cool to room temperature; heat

A.

I

13DV-8623 -12o o 0 o 0 0 0 0 0 9 e 00t 0 0 0 o 00o 0 00 6 0 0 0 o os it I t 6 t t to 2050 25°F for 4 hrs., furnace cool to below 1200°F in 6 min. or less and thereafter to room temperature; and heat to 1650 25°F 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 5% Re provides the highest rupture strength at 1800°F/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 0 F/40 ksi can be obtained with 15 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 5W) 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 1975°F/4 hrs. and at 1650°F/16 hi, Thereafter, different sets of specimens were heated for 1000 hrs.

at 1800°F and for 1000- hrs. at 2000°F. 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 1800°F and 2000°F.

I

I i i i i' E 1 i 13DV-8623 -13- Table V presents the results of cyclic oxidation tests on uncoated 1/4" dia. x 3".long pin specimens conducted at 2150OF 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 metal 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.

S0 The two series I alloys that contained yttrium 15 (802 and 808) had exceptional oxidation resistance.

4 BThe series II alloys, all of which were yttrium-bearing, exhibited no metal loss after 200 Shrs of high velocity oxidation (Mach I) at 2150°F and e? only 2-3 mils 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 tstudies, also results in improved y' stability.

St A I i, a 1- 307

I

.4 13DV-8673 14- TABLE V Oxidation, 21500F, Mach Metal Loss (mils/side) a0 0 *0 :0 09 0 a a 0 0 11 00 00 0 0o 0 00000 .00..

0 Go0 'time Curs) 800 801 802 803 804 805 806 807 808 809 810 811 812 813 20 814 815 816 817 818 819 820 821 822 823 R125 MA7 54 18.662.6 127.6 169.6 214.0 1A~ 0.

0 0.8 0.8 1.0 0.8 1.0 0.3 1.0 1.0 1.3 0 0 0 0 0 0 0 0 0 0 0 0 1.8 1.5 0 *1.5 1.3 1.5 1.8 2.0 0 1.8 1.8 2.5 0 0 0 0 0 0 0- 0 0 0 0 0 4.8 4.8 0.3 2.8 1.5 10.3 4.8 4.5 0.3 2.8 2.8 0 0 0 0 0 0 0 0 0 0 0 6.8 8.0 0 4.0 4.0 7.5 6.3 6.3 0.3 2.0 3.5 4.5 0 0 0 0 0 0 0 9.3 9.8 0.3 5.8 5.3 9.5 9.8 8.0 0.5 4.3 4.0 5.5 0 0 0 0 0 0 0 6-12 4-8 2-3 8-10 8-12 8-10 8'-10 6-8 2-3 10-14 10-16 12-16 1-2 1-2 1-2 1-2 1 1 2 2 2 1-2 1-2 1-2 Depletion (mils/side) 0 0 0 0 0 0 0 0 0 0 -12 ,1~ I

V

maintain their basic material properties. Oftentimes, I 91 13DV-8623 Q 44 44 O0 0 o e 04 o 0 0 001*8 *40944 0 o *8.fo .444 00 0

I

The hot corrosion resistance of the alloys of the invention was evaluated alongs-ide 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-S fuel-fired fl~we at 1600*F with a nominal 1 ppm salt added to the combustion products.

The test was f irst run at -l1 ppm f or 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 t~he next 300-400 hours and above 2 ppm during the remaining 300 hours. The followipg conclusions were drawnt from these hot covrrosion tests: 1) B1900 was least resistant to 15 hot corrosion at all salt l'evels, 2) MMZOO(Hf) was the next l east resistant alloy at all salt levels, x) the alloys of the invention, especially allpy 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, 1M200(Hf), Rene' 125, :and 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 adde.d when the master heat material-was remelted and molten

IA

-~13DV-8623 -16prior to WSing to produce single crystal slabs and turbine blades. For the test specimens used to obtain the data of Tables VII, VIII, IX. and X, yttrim in the, amount of 400 ppm was added. Stress rupture strength data for alloy 821 from the 300 lb master heato, and the 12 lb. laboratory heat are presented in Table IX.

TABLE VI 0 300 Lb Alloy 821 Master Heat Cr 6.79 Ti 0 010 Co 7.30 Re 2.95 Ta 6.40 B 0.004 Al 6.15 Y 0 L4 process could have been used.

'1 13DV-8623 rl7- TABLE VII Stress Rupture Data Tem Stress Life El RA eFT JHrs 7TT T7 Heat H. Treat 12 lb 2390/2 1975/4 Lab. Ht. 1650/16 H.T. as slabs 300 lb Master Ht.

(Alloy 821 P 0* u0 o Or, 00 0o 9 15 o 0 00 0 20 300 lb Alloy 821 2390/2 1975/4 1650/16 H.T. as slabs 1600 1800 2000 2100 1400 1400 1600 1800 1800 1800 1800 1900 1900 2000 2000 2100 130 110 80 40 40 35 30 25 22 20 17.5 13 0 ad O w 0 0 0P al4 4 t4 C Reheat treated* 1900/4 age 1650/4 age Reheat treated* 1975/4 age 1650/4 age Reheat treated* 2050/4 age 1650/4 age 1600 1800 2000 1600 1800 2000 1600 18OO 2000 114.3 80.4 98.4 74.3 1.9 351.6 155.4 72.7 75.8 227.8 509.2 120.2 35 7.2 81.3 391.9 80.5 115.8 68.4 82.7 155.2 85.2 101.2 160.0 103.8 125.7 139.9 97.4 126.9 131.0 90.5 97.2 19.5 14.8 20.1 39.4 20.6 17.5 16.8 10.1 13.9 13.6 13.1 3.4 19.0 17.0 13.9 19.0 25.5 14.7 18.9 18.1 11.6 19.3 23.2 12.8 17.8 20.6 12.4 26.9 24.4 26.8 29.9 33.2 27.3 28.7 23.4 28.6 38.5 23.3 48.6 25.0 30.5 35.2 26.2 39.0 34.4 27.5 28.3 40.3 24.0 28.6 32.9 24.7 29.8 31.1 7.7 43.1 16.8 6.8 Reheat treated* 2125/1 age 1975/4 (coating simulation) 1650/4 age Reheat treated* 2200/1 age 1975/4 (coating simulation) 1650/4 age 1600 1800 2000 1600 1800 2000 All resolutioned in test specimen form at to 20000 F.

2390°F/2 hr fast cool a iJ AL i L Q&A. I III I 13 DV-86Z3 -18 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 23900F/2 hrs. and aged at 1975*F/4 hrs. and 16509F/16 hrs., with the results shown in Tables VIII, IX, and X, respectively, where UTS is ultimate tensile strength; YS is yield strength at 0.2t strain offset; El. is elongation; and RA is reduction in area.

0000 o 00 00 0 00 00 00 0 0 0 0 I~9 00 0 0,00 0 000000 0 0 0000 0000 TABLE VIII Tensile Data (Master Heat Alloy 821) Temp

(OF)

1000 1200

UTS

(Ksi) 128.6 129.6 142.8 143.3 110.1 64.1 (Ksl) (Ksi).

113.4 110.7 El 11.6

RA

M%

18.9 00 OI 0 00 00 0 0 *0 00 0 o (040 1 0 0 £4 1400 1600 1800 2000 112.4 112.8 129.4 94.7 51.2 106.5 102.6 103. 5 71.9 39. 2 18.0 14.2 19.9 9.9 13.3 30.8 10.0 28.1 19.1 21.6 .TABLE

IX

Low Cycle Fatigue (Master Heat Alloy 821) Alternating Pseudostress (ksi) 1 Cycles to Failure Nf 4.9 x 103 2.3 x 103 1 2 t in. compressive strain hold, 2000*F f 13DV-8623 -19- TABLE X High Cycle Fatigue 1 (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 6 0.5 x 10 0o 1 2050°F; A 0.67, 30 Hz 0 0 As discussed at greater length in co-pending co-assigned application Ser. No. 595,854, the superalloys of this invention break with the longestanding wisdom of the single crystal superalloy arts that grain boundary strengthening elements such as a 15 B, Zr and C are to be avoided, 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 7' solvus temperature thus permitting solutionizing heat treatments to be performed at temperatures where complete solutionizing of the 7' 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 ;i a tI

I

I,

13DV-8623 00 o o 00 o 0 0 00 j 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 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, on the order of i;veral hundreds of ergs/cmz, 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, 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 rellabili> 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.

a

I

21- 1. 13DV-8623 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 net alloyed with yttrium. Ceramic/metal instability is controlled by the bulk thermodynamic condition of the system. The more negative the free energy of formation, AG*f, 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 o" when typical SiO 2 and Zr02 ceramic mold and core ooo0 systems are used. Based on thermodynamic o 15 considerations and the work reported in U.S. Department of the Air Force publication AYML-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.

S° It has also been found that melt/mold and core o, interactions are decreased, the retention of yttrium increased and the uniformity of yttrium distribution 25 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 As noted earlier, yttrium was added s 13DV-8623 I- 22general, 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 co.-istently higher than at the trailing edge.

4 4 TABLE

XI

Yttrium Content (ppm) Airfoil Tip Airfoil Near Platform Blade Root Casting Condition LE( 1

TE(

2 LE TE ROOT 3 Low 130 100 160 100 130 S* 20 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 S 100 90 90 150 100 60 100 100 100 130 150S 190 150 120 170 200 240 210 170 1 LE leading edge 2 TE trailing edge 3 ROOT rQot, center :1 i 1 i'K

(IT

S13DV-8623 -23- Additional single crystal investment castings of large solid turbine blades 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 asditions of up to Z400 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 .o was evident; the leading edge Tetained higher yttrium levels compared to the trailing edge.

o.

A Although the present invention has been f described in connection with specific examples, it will 9000o o be understood by those skilled in the art that the oo: present invention is capable of variations and modifications within the scope of the invention as represented by the appended claims.

D o ooe 9 0 o o 9 9 In a

Claims

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, 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. 4. A single crystal article produced from a nickel-base superalloy having a composition according to any one of the preceding claims. 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: 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; heating for 2 hours in a temperature range sufficient to achieve solutioning of at least 95% ,I of the y' phase; S(c) cooling to 2000OF at 100°F/minute minimum and then to 1200°F in 60 minutes or less; heating to 2050 ±25"F for 4 hours; cooling to below 1200°F in 6 minutes or less; and heating to 1650 25 0 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, TI 1 i alloys ana mucn iesb y^ LU-C L U U both 1800*F and 2000*F. 0133h:AB 25 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.
8. The heat treatment of 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, 6.2 Al, 3 Re, 0.15 Hf, 0.05 C, 0.004 B, and 0.01 Y, the balance being nickel and incidental impurities.
9. The heat treatment of claim 6 wherein the temperature range in is from about 2385 to 2395 0 F. DATED this 27th day of June, 1991. GENERAL ELECTRIC COMPANY By Its Patent Attorneys ARTHUR S. CAVE CO. e o 0 0ou4 I ;I i I 1 a a* II', b 1 1-o