NICKEL BASE ALLOY AND TURBINE ENGINE BLADE CAST THEREFROM

Abstract: Improved nickel base alloys for use in vacuum directional solidification casting which alloys include a plurality of ele­ ments but which, in particular, include the following elements in the percentages by weight specified: carbon .07-.09, zirco­ nium .007-.015, tungsten 9.3-9,7, and molybdenum 0.4-0.6. The new ultra high strength casting alloys are free of the prob­ lem of grain boundary cracking which has been encountered in the production of directional solidification castings of this type of alloy and minimize M₆C acicular phase formation in turbine blades under high temperature conditions over extended ope­ rating periods. The invention also includes a turbine engine blade cast from the alloy.

Description: NICKEL BASE ALLOY AND TURBINE ENGINE BLADE CAST THEREFROM This invention is in the field of nickel-base casting alloys particularly intended for vacuum casting, by directional solidification, of thin wall products of complex shape for use under extremely demanding service conditions. Particularly are alloys in accordance with this invention suitable for use in the casting of turbine blades (including vanes) for jet engines where the castings will be subjected to extended periods of operation under very high temperatures. For some time equiaxed grain castings were used for such parts. However, these proved unsatisfactory as the operating conditions for the castings became more severe. In response to these problems the gas turbine industry developed the technique of directional solidification (referred to as DS) with significant improvement in high temperature alloy properties and blade performance. However, parts cast by this technique are susceptible to grain boundary cracking. Designs to increase the power and efficiency of gas turbine engines have required reduction of the wall section of such turbine blade castings to improve blade cooling efficiency. This necessitates the use of special, high strength, cast alloys. Extremely thin wall sections are required to provide rapid thermal transfer because the blades may be operating in a gaseous environment at temperatures above the melting point of the alloy. In such an environment it is critical to provide sufficiently rapid heat transfer to ensure that the actual temperature of the blades is kept below that at which they lose structural integrity. As the wail thickness of the blades has been decreased, the strength of the alloys has had to be increased. However, as the strength of the alloy has been increased, the problem of grain boundary cracking during DS casting of the blades has also increased. The response to this has been the development of improved alloys such as those described in United States Patents 3 526 499, 3 677 747 and 3 720 509. Despite these improvements in ultra high strength alloys suitable for directional solidification, grain boundary cracking has continued to be a problem. According to the present invention, there is provided an improved nickel base alloy for vacuum directional solidification casting of thin wall, complex cored, high core-to-metal volume ratio castings for use under high stress, high temperature conditions and characterized by the absence of grain boundary cracking, consisting of the following elements in percentages by weight (other than those expressed as ppm): C .07-.09 Si .03 max. Mn .10 max. P .005 max. S 20 ppm max. Cr 7.5-8.5 Co 9.0-9.5 Mo 0.4-0.6 W 9.3-9.7 Ti 0.6-0.9 Al 5.4-5.7 Ta 3.1-3.3 Hf 1.4-1.6 B 0.01-0.02 Zr .007-.015 Fe .2 max. Cu .05 max. Nb .10 max. Mg 40 ppm max. Pb 2 ppm max. Se 1.0 ppm max. Bi .3 ppm max. Te .5 ppm max. T1 .5 ppm max. Ni Balance After a long period of investigation and experimentation, it has been learned that by reducing the maximum amount of carbon, zirconium, tungsten and molybdenum, the problem of grain boundary cracking can be significantly reduced or eliminated, without significant sacrifice of strength while at the same time substantially eliminating the formation of acicular M6C under subsequent operating conditions. Acicular phase formation in cast turbine blades is undesirable since it acts as a nucleant for forming premature cracks due to complex fatigue/creep interactions. The problem of grain boundary cracking has been particularly evident in the thin wall, first stage turbine blade and vane designs for advanced turbine engines. The blades have a wall thickness of 0.635 - 0.762 mm (0.025 0.030 inches) with very complex coring and a high core-tometal volume ratio. This type of casting requires a very high strength alloy. This type of casting has been impractical to manufacture from some existing alloys with the required high temperature strength because of potential failure resulting from grain boundary cracking, occurring during the DS casting of the parts. This is a very serious problem in high performance turbine blades and vanes. The operation conditions are such that structural failure is intolerable. Thus, complete and thorough inspection of each part is required. Because of the intricacy of the designs and the minuteness of the passages, these parts can only be inspected internally by suitable probes inserted through each passage and even then the inspector is frequently required to exercise a high degree of judgement as to whether what is seen through the probe is in fact a crack. This is very time consuming and expensive, especially when reject rates involve a substantial percentage of the total castings poured. It is another important facet of this invention that the alloy remain metallurgically stable under the severe temperature of its operating environment. The fact that alloys in accordance with the present invention are stable and will develop minimum, if any, acicular M6C is shown by their chemistry and their low NV3B value. In these alloys the M is mainly W and Mo. While the presence of W and Mo increases the tendency to formation of the acicular M6C phase, these elements along with Ta cannot be eliminated because of their essential role in strengthening at high temperatures. The low value for N is indicative that no deleterious sigma phase forma v3B tio-n will occur. The reduction in C is important in accomplishing the minimized M6C acicular phase formation by reducing the initial as-cast formation of interdendritic and grain boundary MC (Hf, Ta, Ti) carbides. The fact that the amount of C could be reduced while improving alloy characteristics is significant because C, while essential to the formation of interdendritic and grain boundary carbides, is essential to the strength and stability of the alloy In the development of these alloys, it was discovered that Zr, a known source of increased alloy creep strength, was also a causative of grain boundary crack sensitivity, although the reason for this is not fully understood. It is known that Zr migrates to the grain boundaries in these cast Nickel-base alloys. Even very minor changes in the percentage of Zr present were found to produce significant changes in alloy DS castability. It is also important to keep the S at extremely low levels ( < 20 parts per million (ppm)). S migrates to the grain boundary and causes a significant reduction in grain boundary strength and ductility. Alloys in accordance with this invention not only exhibit virtual or total elimination of the grain boundary cracking problem. without unacceptable sacrifice of high temperature strength, but also have dependable repeatability under commercially acceptable manufacturing and foundry use conditions. A specific composition for the alloy of this invention is tabulated in Column A of Table I. In this tabulation it is compared with a commercially used alloy utilizing the teachings of U.S. Patent 3 720 509, set out in Column B: TABLE I A B C .08 .13-.17 Si .03 max. .15 max. Mn .10 max. .20 P .005 max. .015 max. S 20 ppm max. 100 ppm max. Cr 8.0 8.00-8.80 Co 9.3 9.00-11.00 Mo .55 .50-.80 W 9.5 9.50-10.50 Ti .7 .90-1.20 Al 5.60 5.30-5.70 A B Ta 3.2 2.80-3.30 Hf 1.5 1.20-1.60 B .015 .010-.020 Zr .01 .03-.08 Fe .2 max. .25 max. Cu .05 max. .10 max. Nb .10 max. .10 max. Mg 40 ppm max. Not Recorded Pb 2 ppm max. 5 ppm max. Se 1.0 ppm max. 1 ppm max. Bi .3 ppm max. .3 ppm max. Te .5 ppm max. .5 ppm max. T1 .5 ppm max. .5 ppm max. (N) ppm 15 max. Not Rccorded (0) ppm 10 max. Not Recorded N v3B 2.15 max. Not Recorded Ni Balance Balance Many alloys can be formulated in accordance with the invention but some specific examples will now be described. These alloys have been cast in the test device shown in the accompanying drawings, in which: Figure 1 is a partially schematic, sectional elevational view of the test device for determining susceptibility to grain boundary cracking; Figure 2 is a sectional view taken along the plane II-II of Figure 1; and Figure 3 is a photograph of a pair of castings made using the test device of Figures 1 and 2. EXAMPLE I A 113.4 kg (250 pound) heat No. VF-165 utilizing the Column A composition of Table I was manufactured by vacuum induction melting and was tested by vacuum DS casting it as a tube around an alumina core. The equipment for the test is illustrated in Figures 1 and 2. This equipment includes an outer mould shell 10 and an inner tubular mould shell 11 of alumina. The shells were spaced apart at top and bottom and held in concentric relationship by spacer rings 12. Communication is provided between the interior of the inner shell and the annular space between the shells by a pair of vertical channels 13 at the bottom of the inner shell. In this particular test the radial spacing between the shells 10 and 11 forming the annular channel 14 was 0.15 mm and the outside diameter of the shell 11 was 1.9 mm. The test unit was placed on a chill plate 15. The unit was surrounded by an induction heater 16. A rod of the alloy from heat VF-165 was placed within the inner shell 11. The induction heater was turned on to melt the alloy causing it to flow through the channels 13 into the annular space 14. After the annular channel was filled, the induction heater, while still activated, was gradually withdrawn upwardly. This procedure results in directional solidification of the alloy except at the bottom adjacent the chill plate where some equiaxe solidification occurs. This can be seen in the photographs of Figure 3. Using the identical test equipment and procedure, a second tubular casting was made utilizing the alloy of the composition of Column B from heat No. VE-930. The castings, after solidification and cooling to room temperature and removal of the outer mould shell were inspected for evidence of cracks. As clearly appears on the left hand side of Figure 3, numerous longitudinal cracks resulting from grain boundary cracking were observed in the casting made from the alloy set out in Column B of Table I. This cracking was observed to have occurred along grain boundary lines entirely around the test casting. In contrast to this, the casting poured from the alloy composition of Column A of Table I showed no evidence of any cracking except some of about 0.5 mm in length at the extreme upper end and some in the equiaxe zone where little or no directional solidification occurred because of the effect of the chill plate 15. This test is a particularly severe one for alloys susceptible to grain boundary cracking since it develops about a 2% plastic strain. Based upon a rating of A to F wherein: A - no cracks apparent B - only minor cracks at tips C to E - progressively increasing severity of grain boundary cracking F - gross grain boundary cracking the test casting from heat VF-165 (Column A, Table I) was rated A-B. EXAMPLE II Approximately 20 directionally solidified, complex cored, thin wall blades (hereinafter referred to as blades C) were vacuum cast from the alloy of the same heat (VF-165) as that tested in Example I. These were cast using a European-developed, directional solidification, casting process. Identical blades (hereinafter referred to as blades D) were cast from heats V-5134 and V-5224 of the alloy set out in Column B of Table I, except for the following differences: Si .10 max. Mn .10 max P .010 max. S 20 ppm max. Mg 40 ppm max. Pb 2 ppm max. (N) ppm 15 max. (0) ppm 10 max. The blades cast from both alloys were aged by heat 0 treating at 1600 F. for 20 hours. From the blades of both castings, test specimens of 2.26 mm (0.089 inches) diameter were machine cut and tested for stress-rupture with the following results: Blades C LONGITUDINAL Life Hours Mean Life Heat No. 0 500 MPa/850 C 103, 100, 102 (Mean 102) VF-165 (72.5 ksi/l5600F) 145 MPa/10400C 42, 44, 49 (Mean 45) VF-165 (21.0 ksi/19000F) Blades D LONGITUDINAL Life Hours Mean Life Heat No. 500 MPa/8500C 134, 143, 155 (Mean 144) V-5134 0 02.5 ksi/1560 F) 115, 128, 120 (Mean 121) V-5224 0 145 MPa/850 C 53, 55, 64 (Mean 57) V-5134 (21.0 ksi/19000F) 47, 48, 55 (Mean 50) V-5224 TRANSVERSE 465 MPa/8500C 130, 138 (Mean 134) V-5134 (67.5 ksi/156OOF) V-5224 0 108 MPa/1040 C 142, 189 (Mean 165) V-5134 (15.7 ksi/19000F) V-5224 Tensile r'oer"ties Blade temp. MPa (ksi) % elongation C 6490C. (12000F) 917 (133) 11.5 D 649 C. (12000F) 855 (124) 6.4 C 749 C. (13800F) 1000 (11+5) 8.1 D 7490C. (13800F) 938 (136) 5.5 C 849 0C. (15600F) 869 (126) 10.5 D 8490C. (15600F) 862 (125) 6.9 In the case of all of these tests from which this data was compiled the test continued until failure by rupture occurred. Careful examination of the blades C, including X-ray and fluorescent dyeenetrant inspection, revealed no cracks in any of the blades. This is despite the fact that these blades were of a design considered to be "crack prone". These tests indicate that the alloy has overcome the problem of grain boundary cracking with improved tensile properties while retaining acceptable stress rupture properties. The blades also were cut up and microscopically inspected for microporosity. The blades were found to have internal microporosity of p0.2% with most of the fields exhibiting a microporosity of about 0.1%. These performance capability results for the alloy of Column A are at least equal to those for the alloy of Column B while almost entirely eliminating part rejection because of grain boundary cracking and represent a substantial improvement over other conventional alloys which have been used for casting of high performance turbine blades.

Claims: CLAIMS 1. An improved nickel base alloy for vacuum directional solidification casting of thin wall, complex cored, high core-to-metal volume ratio castings for use under high stress, high temperature conditions and characterized by the absence of grain boundary cracking, consisting of the following elements in percentages by weight (other than those expressed as ppm): C .07-.09 Si .03 max. Mn .10 max. P .005 max. S 20 ppm max. Cr 7.5-8.5 Co 9.0-9.5 Mo 0.4-0.6 W 9.3-9.7 Ti 0.6-0.9 Al 5.4-5.7 Ta 3.1-3.3 Hf 1.4-1.6 B 0.01-0.02 Zr .007-.015 Fe .2 max. Cu .05 max. Nb .10 max. Mg 40 ppm max. Pb 2 ppm max. Se 1.0 ppm max. Bi .3 ppm max. Te .5 ppm max. T1 .5 ppm max. Ni Balance 2. An alloy as claimed -in Claim 1 in which the following elements are present in the weight percentages specified: C .08 Cr 8.0 Co 9.3 Mo .55 W 9.5 Ti .7 Al 5.60 Ta 3.2 Hf 1.5 B .015 Zr .01 3. An alloy as claimed in Claim 1 or Claim 2 in which the metallurgical stability factor of the alloy expressed by the formula N is 2.15 max. v3B 4. An alloy as claimed in Claim 1 or Claim 2 or Claim 3 in which the maximum of each of the elements nitrogen and oxygen is 15 parts per million and 10 parts per million, respectively. 5. An alloy as claimed in any of the preceding claims having a mean longitudinal machined-from-blade (MFB) stress-rupture life of 45 hours at 1040 C. (l9000F) under an applied load of 145 MPa (21.0 ksi). 6. An alloy as claimed in any of the preceding claims having a mean longitudinal MFB stress-rupture life of 102 hours at 8500C. (15600F) under an applied load of 500 MPa (72.5 ksi). 7. A turbine engine blade of an alloy according to any of the preceding claims cast under vacuum and directionally solidified.

Inventors: HARRIS, KENNETH
Applicants: CANNON-MUSKEGON CORPORATION
Application year: 1981
Main classification: C22C19
IPC codes: C30B21, C22C19, F01D5