CA2091827A1 - Gas turbine blade alloy - Google Patents
Gas turbine blade alloyInfo
- Publication number
- CA2091827A1 CA2091827A1 CA002091827A CA2091827A CA2091827A1 CA 2091827 A1 CA2091827 A1 CA 2091827A1 CA 002091827 A CA002091827 A CA 002091827A CA 2091827 A CA2091827 A CA 2091827A CA 2091827 A1 CA2091827 A1 CA 2091827A1
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- Prior art keywords
- alloy
- titanium
- aluminum
- nickel
- chromium
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- C—CHEMISTRY; METALLURGY
- C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
- C22C—ALLOYS
- C22C19/00—Alloys based on nickel or cobalt
- C22C19/03—Alloys based on nickel or cobalt based on nickel
- C22C19/05—Alloys based on nickel or cobalt based on nickel with chromium
- C22C19/051—Alloys based on nickel or cobalt based on nickel with chromium and Mo or W
- C22C19/056—Alloys based on nickel or cobalt based on nickel with chromium and Mo or W with the maximum Cr content being at least 10% but less than 20%
-
- C—CHEMISTRY; METALLURGY
- C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
- C22C—ALLOYS
- C22C19/00—Alloys based on nickel or cobalt
- C22C19/03—Alloys based on nickel or cobalt based on nickel
- C22C19/05—Alloys based on nickel or cobalt based on nickel with chromium
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- Chemical & Material Sciences (AREA)
- Engineering & Computer Science (AREA)
- Materials Engineering (AREA)
- Mechanical Engineering (AREA)
- Metallurgy (AREA)
- Organic Chemistry (AREA)
- Turbine Rotor Nozzle Sealing (AREA)
Abstract
ABSTRACT
A nickel-based alloy suitable for use in making gas turbine components, such as rotating blades is provided. The alloy comprises the following elements in weight percent:
Chromium 14.75 to 16.0, Cobalt 8.0 to 8.5, Aluminum 3.4 to 4.0, Titanium 3.4 to 4.3, Aluminum plus Titanium 7.7 to 8.3, Tantalum 1.75 to 2.7, Tungsten 2.0 to 4.0, Columbium up to .5, Molybdenum up to 2.0, Carbon .05 to .12 and the balance Nickel. The alloy of the current invention may also comprise impurities and incidental elements generally associated with nickel-based alloys, such as Zirconium up to .06 and Boron up to .015 percent by weight.
A nickel-based alloy suitable for use in making gas turbine components, such as rotating blades is provided. The alloy comprises the following elements in weight percent:
Chromium 14.75 to 16.0, Cobalt 8.0 to 8.5, Aluminum 3.4 to 4.0, Titanium 3.4 to 4.3, Aluminum plus Titanium 7.7 to 8.3, Tantalum 1.75 to 2.7, Tungsten 2.0 to 4.0, Columbium up to .5, Molybdenum up to 2.0, Carbon .05 to .12 and the balance Nickel. The alloy of the current invention may also comprise impurities and incidental elements generally associated with nickel-based alloys, such as Zirconium up to .06 and Boron up to .015 percent by weight.
Description
- 1 - 57,151 BACKGROUND OF THE INVENTION
The current invention relates to an alloy suitable for use in making gas turbine components, such as the rota~ing blades in the turbine section of a gas turbine. More specifically, the current invention concerns a nickel-based alloy having a sufficiently high chromium content for good corrosion resistance yet maintaining high strength when used to make a directionally solidified turbine blade casting.
A gas turbine employs a plurality of rotating blades in its turbine section. Such blades are exposed to gas at temperatures in excess of 1100C (2000F) and subjected to high stress. Consequently, the alloys from which such blades are cast must, after suitable heat treatment, have very high stress rupture skrength and sufficient metallurgical stability ; 15 to maintain this strength for many thousands of hours of operation. Such alloys rnust also have sufficient ductility to withstand the large thermal stresses imposed on turbine blades. In addition, as a result of impurities in the fuel and combustion air, the gases to which the blades are exposed contain corrosive compounds, such as sulfides and chlorides.
Consequently, such blade alloys must also have good hish temperature corrosion resistance, as well as oxidation resistance.
Generally, high temperature corrosion resistance is provided by the incorporation of substantial amounts of chromium into turbine blade alloys. High chromium content inhibits the basic fluxing of the alloy by forming a continuous chromia scale khat is not susceptible to solution _ ~ _ 57,151 and reprecipitation from a Na2SO4 melt, thereby providing an effective barrier for the alloy. Moreover, chromium combines with sulfur to form high melting point sulfides, thereby inhibiting degradation due to sulfidation. As discussed in R. Streiff and D.H. Boone, "Corrosion Resistant Modified Aluminide Coating," Journal of Materials Engineering (1988), a minimum level of 15% chromium is considered necessary for good high temperature corrosion resistance.
One nickel-based alloy, used with some success for a number of years in gas turbine blades, is manufactured by the International Nickel Company and known commercially as IN-738. A typical composition of IN-738 in weight percent, as published in the American Society of Metal Handbook, 9th ed., volume 4, page 244, is as follows: Nickel 61, Chromium 16.0, Cobalt 8.5, Molybdenum 1.7, Tungsten 2.5, Tantalum 1.7, Niobium (aka Columbium) 0.9, Aluminum 3.4, Titanium 3.4, Carbon .17, Boron 0.01, and Zirconium 0.10. In addition, the manufacturer recommends that the electron vacancy number for this alloy not exceed 2.36. This alloy is disclosed in U.S
Patent No. 3,459,545 (Bleber), hereby incorporated by reference in its entirety.
It is known that directional solidification, whereby a uniaxial grain structure is produced, increases both the ability to withstand cyclical thermal stress and the stress rupture strength of many nickel-based alloys -- for example, U.S. Patent No. 4,519,979 (Shaw) discloses that directional solidification of an alloy known commercially as IN-939 and having a composition in weight percent of Carbon 0.15, Chromium 22.5, Cobalt 19, Tungsten 2, Titanium 3.7, Aluminum 1.9, Tantalum 1.4, Niobium 1.0, Zirconium 0.1, Boron 0.01 and the balance Nickel, increased the stress the rupture life of the alloy fxom about 850 hours to 1370 hours at 870C and 200 N/mm2. The inventor has confirmed, however, that directional solidification does little to increase the stress rupture strength of the IN-738 alloy.
Generally, it has been observed that the alloys that derive the most improvement from directional solidification $ ~ 3 _ 3 _ 57,151 have three characteristics -- (i) a relatively low chromium content, (ii) a high gamma prime volume fraction and (iii) a high solution temperature. Consequently, it has previously been thought that the advantages of directional solidification could be obtained for alloys such as IN-738 only by reducing their chromium content -- see, for example, M. McLean, "Directionally Solidified Materials for High Temperature Service," published by Metals Society, London (1983), page 153. Thus, one popular alloy used in d.irectionally solidified blades, known commercially as GTD-111, has a typical composition in weight percent o~ Chromium 14.0, Cobalt 9.5, Aluminum 3, Titanium 4.9, Tantalum 2.8, Tungsten 3.8, Molybdenum 1.5, Boron 0.01, Carbon O.1 and the balance Nickel, as disclosed by R. Viswanathan in "Damage Mechanisms and Life Assessment of High-Temperature Components," published by the American Society of Metals (1989). As can be seen, in GTD-lll the benefits of directional solidification have been obtained by reducing the 16% chromium level used in IN-738 to only 14%.
Unfortunately, as previously discussed, such relatively low levels of chromium result in inadequate corrosion resistance.
Accordingly, it would be desirable to provide a blade alloy that has high strength when directionally solidified yet sufficient chromium to achieve adequate corrosion resistance.
SUMMARY OF THE_INVENTION
It is an object of the current invention to provide a nickel based alLoy suitable for use in casting the rotating blades of a gas turbine and having high stress rupture strength when directionally solidified, at least 15% hy weight chromium for good corrosion resistance, and stability during long term service.
This objectt as well as other objects, is accomplished in a gas turbine having a rotating blade made from an alloy comprising the following elements in weight percent: Chromium 14.75 to 16.0, Cobalt 8.0 to 8.5, Aluminum 3.4 to 4.0, Titanium 3.4 to 4.3, Aluminum plus Titanium 7.7 to 8.3, Tantalum 1.75 to 2.7, Tungsten 2.0 to 4.0, Carbon .05 3`~ ~
- 4 - 57,151 to .12, Columbium up to .5, Molybdenum up to 2.0, and the balance Nickel. The alloy of the current invention may also comprise impurities and incidental elements generally associated with nickel-based alloys, such as Zirconium up to .06 and Boron up to .015 percent by weight.
BRIEE DESCRIPTION OF THE DRAWINGS
Figure 1 is a graph of rupture time t, in hours, versus stress a, in KSI (N/mm2), for both conventionally cast (CC) and directionally solidified (DS) specimens of IN-738, showing the effect of directional solidification on the stress rupture strength of IN-738.
Figure 2 is an isometric view of a gas turbine rotating blade.
Figure 3 is a bar chart of rupture time, t, in thousands of hours, versus stress ~, in KSI (N/mm2), showing the stress rupture life at 870C (1600F) and four stress levels of (i) four heats of the alloy according to the current invention, SAS1-SAS4, as directionally cast, (ii) conventionally cast IN-738 and (iii) directionally solidified IN-6203.
DESCRIPTION OF THE PREFERRED EMBODIMENT
Figure 1 shows a graph of time to rupture versus stress at three temperature levels for both conventionally cast -- that is, having an equia~ed grain structure -- and directionally solidified IN-738 test specimens. As can be seen, data at both the 870C (1600F) and 925C (1700~F) temperature levels indicates that the rupture life of the directionally solidified specimens is worse than that of the conventionally cast specimen above a stress level of about 275 N/mm2 (40 KSI)o There is shown in Figure 2 a rotating blade 1 used in the turbine section of a gas turbine. The current invention is directed to an alloy, referred to as OM 200, from which such blades may be cast, especially using a directional solidification casting process. Generally speaking, the current invention is directed to an alloy comprising the following elements in weight percent: Chromium 14.75 to 16.0, - 5 - 57,151 Cobalt 8.0 to 8.5, Aluminum 3.4 to 4.0, Titanium 3.4 to 4.3, Aluminum plus Titanium 7.7 to 8.3, Tantalum 1.75 to 2.7, Tungsten 2.0 to 4.0l Carbon .05 to .12, Columbium up to .5, Molybdenum up to 2.0, and the balance Nickel. The alloy of the current invention may also comprise impurities and incidental elements generally associated with nickel-based alloys, such as Zirconium up to .06 and Boron up to .015 percent by weight. Prefexably, the alloy of the current invention consists essentially of the following elements in weight percent: Chromium about 15.5, Cobalt about ~.0, Aluminum about 4.0, Titanium abou$ 3.8, Aluminum plus Titanium about 7O8~ Tantalum about 2.7, Tungsten about 2.6, Molybdenum 0.5, Carbon about 0.08 and the balance Nickel.
As previously discussed, both corrosion resistance and strength are important properties for blade alloys.
Corrosion resistance in nickel-based alloys is provided primarily by chromium. Nickel-based alloys used for gas turbine components are strengthened by three mechanisms -- (i~
solid solution strengthening, (ii) strengthening resulting from the presence of carbides and (iii) gamma prime strengthening. Solid solution strengthening is provided by molybdenuml chromium and tungsten and, to a lesser extent, by cobalt, iron and vanadium. Gamma prime strengthening is provided primarily by aluminum and titanium, which strengthen the austenitic matrix through the precipitation of Ni3(Al and/or Ti), an fcc intermetallic compound. In addition, the aluminum in gamma prime can be replaced by tantalum and columbium.
Since the strength of nickel-based super alloys increases with increasing gamma prime volume fraction up to about 60%, in the alloy according to the current invention the content of aluminum, titanium, tantalum and columbium, which tend to increase the gamma prime volume fraction, have been controlled so as to achieve a hi~h gamma prime volume fraction. The amount of gamma prime in an alloy can be determined as discussed below with respect to the determination of the electron vacancy number Nv. The amount 2Pi 6 57,151 of gamma prime, in weight percent, in four heats of the alloy according to the current invention shown in Table I are approximately 52, 54, 56 and 54, respectively. The aforementioned preferred composition of the alloy according to the curren-t invention has approxiamtely 56% gamma prime.
By comparison, the amount of gamma prime in IN-738 is appro~imately 50% by weisht. (It should be noted that the aforementioned variation in the gamma prime content among the four heats of the alloy according to the current invention did not adversely affect corrosion resistance or stability.) As a result of operation at elevated tempexature, nicXel-based alloys undergo microstructural chang~s. Such changes include gamma prime coarsening, which ad~ersely affects the strength of the alloy, and the transformation of gamma prime into unwanted topologically close-packed secondary phases, such as plate or needle-like sigma, eta, etc. The formation of these plate-like phases adversely affects both strength and toughness. Consequently, in order to ensure that high strength and toughness are maintained for many thousands of hours of operation at elevated temperature~ the composition of the strengthening elements must be carefully balanced, as explained below, so that the alloy has a certain degree of microstructural stability.
It is well known that the electron vacancy number Nv is a measure of the microstructural stability of the alloy.
The higher the ~alue of Nv the greater the tendency for the formation of the aforementioned topologically close-packed secondary phases -- see for example, the American Society of Metal Handbook, 9th ed., volume 4, page 278. As disclosed in the brochure entitled ~Alloy IN-738, Technical Data,"
published by The International Nickel Company and hereby incorporated by reference in its entirety, the electron vacancy number may be calculated from the equation Nv = 0 66 Ni + 1.71 Co + 2.66 Fe + 4.66 (Cr + Mo + W) + 6.~6 Zr. The atomic percent of each element in the matri~ to be substituted into this equation is determined by convexting the composition from weight percen~ to atomic percent and assuming that (i) - 7 - 57,151 one-half the carbon forms MC in the order of TaC, CbC, TiC, ~ii) the remaining carhon forms M23C6 with the M comprising twenty three atoms of Cr, (iii) boron is combined as Mo3B2, (iii) gamma prime is Ni3(Al, Ti, Ta, Cb), and (iv) the residual matrix consists of the atomic percent minus those atoms contained in the carbides, the boride and the gamma prime reaction so that the total of the remaining atomic percentages gives the atomic concentration in the matrix.
Conversion of these atomic percentages to a 100 percent basis gives the atomic percentage of each element to be substituted into the equation above. Since it is desirable to maintain a low value of Nv, in the alloy according to the current invention the composition is adjusted so that, in addition to obtaining high strength, the electron vacancy number of the alloy does not exceed about 2.4. In the preferred composition of the alloy the electron vacancy numbex is equal to about 2.4.
As previously discussed, the aluminum in gamma prime can be replaced by columbium and/or tantalum, as well as titanium. However, since replacement of Al by Ti, Cb or ~a in Ni3Al adversely affects the misfit between the austenitic matrix and the gamma prime particles, the stability of these compounds in order of decreasing stability is Ni3Al, Ni3Ti and Ni3Cb(orTa). As a result, the titanium/aluminum ratio plays a major role in gamma prime coarsening. The titanium/aluminum ratio also plays a major role in the transformation of gamma prime into the aforementioned unwanted plate-like phases.
Accordingly, in order to minimize gamma prime coarsening and to improve the stabilit~ of the gamma prime phase, the titanium/aluminum ratio is ~enerally maintained below 2:1 (by weight). The transformation of gamma prime into unwanted needle or plate-like phases can also be retarded by the addition of tungsten.
To assess the stability of the alloy according to the current invention, the shoulder and grip sections of creep rupture specimens tested at 871C (1600F) and 205.9 N/mm2 (30 XSI) were examined metallographically for sigma, eta and mu 2'~
phases. It was found that the aging of these specimens at 871C, under stress/ for up to 4000 hours did not produce any undesirable phases, thereby suggesting that the alloy of the current invention is stable.
In the alloy according to the current invention, the amounts of aluminum, titanium, tantalum, columbium and molybdenum have been balanced so as to attain high strength when the alloy is directionally solidified while maintaining good microstructural stability. Significantly, this result has been achieved without the need to reduce the chromium content, and, -therefore, without impairing corrosion xesistance, as had heretofore been thought necessary by those skilled in the art. Specifically, in the alloy according to the current invention, the content of aluminum and titanium has been increased, when compared with IN-738, to a minimum of 3.4% for each, with the minimum combined aluminum plus titanium content being 7.7%. The maximum amounts of aluminum and titanium are 4.0% and 4.3%, respectively, with the maximum aluminum plus titanium content being 8.3%. In addition, the amounts of columbium and molybdenum have been reduced so that the optimum composition includes no columbium and only 0.5%
molybdenum. At most only 0.5% columbium and 2.0~ molybdenum are permitted in the alloy of the current invention.
Chromium, however, has been maintained in the range of 14.75%
to 16%, so that adequate corrosion resistance is maintained.
Tungsten is maintained in the range of 2.0% to 4.0% and tantalum in the range of 1.75% to 2.7%. Zirconium and boron are limited to 0.06% and 0.015~, respectively, with none of either of these elements being present in the most preferred composition. Carbon is maintained in the 0.05% to 0.12%
range. Moreover, as previously discussed, in the alloy according to the current invention, the elements are adjusted wîthin the aforementioned ranges so that the eleckron vacancy number is maintained at no more than about 2.4, thereby ensuring that adequate microstructural stability is achieved.
As a result of the foregoing, the strength of the alloy according to the current invention when cast by a 9 ~ 57,151 directional solidification process is high, despite its relatively high chromium content. Significantly, good microstructural stability of the alloy according to the current invention has been achieved by considerably reducing the levels of columbium and molybdenum, when compared to IN-738, so that the amounts of aluminum and titanium can be increased without driving the electron vacancy number too high.
Four heats of the alloy according to the current invention, designated SASl through SAS4, were prepared using the compositions in weight percent shown in Table I.
Specimens were cast from these heats 80 as to be directionally solidified and were given the standard heat treatment suggested for IN-738 -- that is, a 1121C (2050F~ solution heat treatment for two hours followed by a 843C (1550F) aging treatment for twenty four hours. These specimens were subjected to a variety of tests and the results compared to conventionally cast IN-738 and directionally solidified IN-6203. IN-6203 is a nickel-based alloy having a nominal composition in weight percent of Chromium 22.0, Cobalt 19.0, Aluminum 2.3, Titanium 3.5, Tantalum 1.10, Columbium 0.80:
Tungsten 2.00, Boron 0.01, Carbon 0.15, Zirconium 0.10 Hafnium 0.75 and the balance Nickel.
TABLE I
Element SASl SAS2 SAS3 AS4 Cr 16.00 16.00 14.75 16.00 Co 8.50 8.00 8.00 8.00 Al 3.40 3.40 3.60 4.00 Ti 4.30 4.30 4.30 3.40 Al+Ti 7.70 7.70 7.90 7.40 W 2.60 2.60 2.60 2.60 Mo 1.75 0.90 1.00 0.60 Ta 1.75 2.70 2.70 2.70 Cb 0.00 0.00 0.50 0.00 C 0.09 0.09 0.09 0.09 Zr 0.04 0.04 0.04 0.04 B 0.01 0.01 0.01 0.01 Ni Bal Bal Bal Bal The results of low fatigue testing at 871C (1600F) of specimens from heats SAS1, SAS3 and SAS4 are shown in Table II and indicate that the cyclic life of the alloy according - 10 - 57,151 to the current invention is superior to IN-738. The results of impact testing of three specimens from heat SASl -- taken from the tip, mid-section and bottom of th~ cast test bar, respectively - are shown in Table III and reveal that the impact strength of the alloy of the current invention is comparable to that of IN-738. Note that IN-738 is provided in both low and high carbon forms and the data is reported accordingly where appropriate. In its low carbon form the carbon content is reduced from 0.17% to 0.11% and zirconium is reduced from 0.10~ to 0.05%t be weight, compared to conventional high carbon IN-738.
The results of stress rupture tests are shown in Tables IV and V and Figure 3, wherein the rupture life of specimens from each of the four heats are compared with those of conventionally cast IN-738 and directionally solidified IN-6203. As can be seen, the data shows that the stress rupture life of the alloy according to the curren-t invention, when directionally solidified, is significantly superior to both that of both IN-738 and IN-6203 -- for example, at 69 N/mm2 (10 KSI), the alloy of the current invention can provide the same rupture life as IN-738 and IN-6203 at a temperature about 56C (100F) higher. When applied to a gas turbine blade, such increased metal temperature capability allows the turbine to operate at higher gas temperatures, ~hereby significantly improving its performance.
~ABLE I I
Total Strain Cycles to Fail~lre Range IN-738LC SAS1 SAS3 SAS4 0.012 - 1412 1442 1061 0.01 317 2990 3973 2852 0.0085 862 5342 6673 11200 0.0075 - 16220 50173+ 10504 0.007 1510 0.00~ 1819 0.005 2503 0.004 10153 - - -2. ~ ~
- 11 - 57~151 T~BLE III
Temp., C _ Impact Energy~ Joules IN-738HC IN-738LC SASl Tip Mid Bot 5 24 6.3 8.8 6.8 - 8.1 260 6.9 - 7.9 11.316.7 538 11.7 - 7.5 9.612.1 649 7.6 10.0 6.3 8.714.1 760 7.1 8.5 6.0 7.39.9 10871 8.1 9.8 9.6 9.114.4 Table IV
Temp., C Time to Rupture~ Hours @ 69 N/mm2 ~10 KSI) _N-738 IN-6203 SASl 968 2~96 982 1437 2752 8400+
996 820 1359 6450+
Table V
S$resæ Time to Rupture, Hours @ 871C ~
N/mm2 (KSI ) IN-738 IN-6203SASl SAS3 SAS4 275.8 (40)188 169 403 633 317 24~.3 (35)420 683 1165 1397 776 25206.9 (30) 1143 1~44 3272 4211 2676 172.4 (25)3177 - 8~00~8800~8247 Turbine blades cast from the alloy of the current invention are advantageously made by vacuum-induction melting and vacuum casting using a directional solidification process.
Directional solidification causes th~ grain boundaries to be oriented substantially parallel to the principal stress axis of the blade with almost no grain boundaries orient~d normal to the principal stress axis. Techniques for directional solidification are well known in the art -- see, for example, U.S. Patent Nos. 3,260,505 (Ver Snyder)l 3,494,709 (Piearcey) and 3,897,815 (Smashey), hereby incorporated by reference in their en~irety.
The gamma prime distribution depends on heat treatment, as well as composition. The standard heat treatment for nickel-based alloys such as IN-738 -~ i.e., a solution treatment followed by an aging treatment -- produces duplex gamma prime comprised of coarse, cuboidal primary gamma prime and fine, spherical gamma prime in approximately equal - 12 - 57,151 amounts. The coarse gamma prime is undesolved gamma prime that did not go into solution during the solution treatment.
Hence, the amount of coarse gamma prime present in the alloy depends on the degree by which the solution temperature is below the ga~a prime solvus temperature, at which all of the gamma prime goes into solution -- that is, the lower the solution temperature r the greater the amount of coarse gamma prime. The fine gamma prime forms during the aging treatment, the amount depending on the amount of gamma prime that did not go into solution during solution treatment.
Directionally solidified components are cooled at a slower rate than conventionally cast components. As a result, directional solidification produces coarser primary gamma prime so that the volume fraction of cuboidal gamma prime is higher and that of spherical gamma prime is lower than in directionally cast components. Unfortunately, a high volume fraction of coarse gamma prime has a deleterious effect on strength. This discrepancy in gamma prime distribution can be minimized by optimizing the heat treatment, particularly the solution treatment, based on the gamma prime solvus and incipient melting temperatures, both of which depend on alloy composition. Generally, it is desirable that the heat treatment be performed at a temperature between the solvus temperature and the incipient melting temperature. Hence, in or~er to obtain the full benefits of heat treatment, the solvus temperature should be considerably below the incipient melting temperature. Unfortunately, although aluminum, titanium, tantalum and columbium increase the gamma prime volume fraction, and therefore, strength, as previously discussed, they also have the effect of raising the solvus temperature and decreasing the incipient melting temperature, thereby narrowing the heat treatment window.
The gamma prime solvus and incipient melting temperatures for three melts of the alloy according to the current invention are shown in Table VI. These temperatures were determined using the differential thermal analysis and gradient bar -- i.e., metallography -- method, in which the - 13 - 57,151 bar was exposed to various temperatures in the 1066C (1950F) ~o 1427C t2300F) temperature range for four hours and then fan cooled. As can be seen, the solvus temperature varies from 1211C to 1229C. By comparison, the solvus temperature for IN-738 is approximately 1204C (2200F). As previously mentioned, IN-738 is typically given a 1121C (2050F) solution heat treatment for two hours followed by a 843C
(1550F) aging treatment for twenty four hours. Various heat treatment regimes for the alloy according to the current invention, based on its incipient melting and solvus temperatures, are discussed further below.
TABLE VI
SASl SAS3 SAS4 Incipient melting temp. C 1241 1229 1229 15Gamma prime solvus temp., C 1211 1211 1229 According to the current invention, the as~cast blades may be heat treated in any of four ways -- (i) solution treating for ~ hours at 1121C (2050F), followed by aging for 24 hours at 843C (1550F), (ii) solution treating for 4 hours at 1149C (2100F), followed by agins for 24 hours at 843C
(1550F), (iii) solution treating Eor 4 hours at 1204C
(2200F), resolution treating for 2 hours at 1121C (2050F) followed by aging for 24 hours at 843C (1550F), and (iv) solution treating for 4 hours a~ 1204C (2200F), resolution treating for 4 hours at 1149C (2100F), followed by aging for 24 hours at 843C (1550F). These heat treatment optimizes mechanical strength by creating a duplex gamma prime structure in a gamma matrix and discrete chrome-carbides at the grain boundaries. The use of higher solution temperatures will decrease the amount of coarse gamma prime and increase the amount of fine gamma prime, hence, further strengthening the alloy.
Although the current invention has been described with reference to an alloy for casting turbine blades, it should be understood that the alloy according to the current invention may also be used to form other components that require high strength at elevated temperature and good corrosion resistance. Accordingly, the present invention may ~ 182r~
- 14 - 57, 151 be embodied in other specific forms without departing from the spirit or essential attributes thereof and, accordingly, reference should be made to the appended claims, rather than to the foregoing specification, as indicating the scope of the invention.
The current invention relates to an alloy suitable for use in making gas turbine components, such as the rota~ing blades in the turbine section of a gas turbine. More specifically, the current invention concerns a nickel-based alloy having a sufficiently high chromium content for good corrosion resistance yet maintaining high strength when used to make a directionally solidified turbine blade casting.
A gas turbine employs a plurality of rotating blades in its turbine section. Such blades are exposed to gas at temperatures in excess of 1100C (2000F) and subjected to high stress. Consequently, the alloys from which such blades are cast must, after suitable heat treatment, have very high stress rupture skrength and sufficient metallurgical stability ; 15 to maintain this strength for many thousands of hours of operation. Such alloys rnust also have sufficient ductility to withstand the large thermal stresses imposed on turbine blades. In addition, as a result of impurities in the fuel and combustion air, the gases to which the blades are exposed contain corrosive compounds, such as sulfides and chlorides.
Consequently, such blade alloys must also have good hish temperature corrosion resistance, as well as oxidation resistance.
Generally, high temperature corrosion resistance is provided by the incorporation of substantial amounts of chromium into turbine blade alloys. High chromium content inhibits the basic fluxing of the alloy by forming a continuous chromia scale khat is not susceptible to solution _ ~ _ 57,151 and reprecipitation from a Na2SO4 melt, thereby providing an effective barrier for the alloy. Moreover, chromium combines with sulfur to form high melting point sulfides, thereby inhibiting degradation due to sulfidation. As discussed in R. Streiff and D.H. Boone, "Corrosion Resistant Modified Aluminide Coating," Journal of Materials Engineering (1988), a minimum level of 15% chromium is considered necessary for good high temperature corrosion resistance.
One nickel-based alloy, used with some success for a number of years in gas turbine blades, is manufactured by the International Nickel Company and known commercially as IN-738. A typical composition of IN-738 in weight percent, as published in the American Society of Metal Handbook, 9th ed., volume 4, page 244, is as follows: Nickel 61, Chromium 16.0, Cobalt 8.5, Molybdenum 1.7, Tungsten 2.5, Tantalum 1.7, Niobium (aka Columbium) 0.9, Aluminum 3.4, Titanium 3.4, Carbon .17, Boron 0.01, and Zirconium 0.10. In addition, the manufacturer recommends that the electron vacancy number for this alloy not exceed 2.36. This alloy is disclosed in U.S
Patent No. 3,459,545 (Bleber), hereby incorporated by reference in its entirety.
It is known that directional solidification, whereby a uniaxial grain structure is produced, increases both the ability to withstand cyclical thermal stress and the stress rupture strength of many nickel-based alloys -- for example, U.S. Patent No. 4,519,979 (Shaw) discloses that directional solidification of an alloy known commercially as IN-939 and having a composition in weight percent of Carbon 0.15, Chromium 22.5, Cobalt 19, Tungsten 2, Titanium 3.7, Aluminum 1.9, Tantalum 1.4, Niobium 1.0, Zirconium 0.1, Boron 0.01 and the balance Nickel, increased the stress the rupture life of the alloy fxom about 850 hours to 1370 hours at 870C and 200 N/mm2. The inventor has confirmed, however, that directional solidification does little to increase the stress rupture strength of the IN-738 alloy.
Generally, it has been observed that the alloys that derive the most improvement from directional solidification $ ~ 3 _ 3 _ 57,151 have three characteristics -- (i) a relatively low chromium content, (ii) a high gamma prime volume fraction and (iii) a high solution temperature. Consequently, it has previously been thought that the advantages of directional solidification could be obtained for alloys such as IN-738 only by reducing their chromium content -- see, for example, M. McLean, "Directionally Solidified Materials for High Temperature Service," published by Metals Society, London (1983), page 153. Thus, one popular alloy used in d.irectionally solidified blades, known commercially as GTD-111, has a typical composition in weight percent o~ Chromium 14.0, Cobalt 9.5, Aluminum 3, Titanium 4.9, Tantalum 2.8, Tungsten 3.8, Molybdenum 1.5, Boron 0.01, Carbon O.1 and the balance Nickel, as disclosed by R. Viswanathan in "Damage Mechanisms and Life Assessment of High-Temperature Components," published by the American Society of Metals (1989). As can be seen, in GTD-lll the benefits of directional solidification have been obtained by reducing the 16% chromium level used in IN-738 to only 14%.
Unfortunately, as previously discussed, such relatively low levels of chromium result in inadequate corrosion resistance.
Accordingly, it would be desirable to provide a blade alloy that has high strength when directionally solidified yet sufficient chromium to achieve adequate corrosion resistance.
SUMMARY OF THE_INVENTION
It is an object of the current invention to provide a nickel based alLoy suitable for use in casting the rotating blades of a gas turbine and having high stress rupture strength when directionally solidified, at least 15% hy weight chromium for good corrosion resistance, and stability during long term service.
This objectt as well as other objects, is accomplished in a gas turbine having a rotating blade made from an alloy comprising the following elements in weight percent: Chromium 14.75 to 16.0, Cobalt 8.0 to 8.5, Aluminum 3.4 to 4.0, Titanium 3.4 to 4.3, Aluminum plus Titanium 7.7 to 8.3, Tantalum 1.75 to 2.7, Tungsten 2.0 to 4.0, Carbon .05 3`~ ~
- 4 - 57,151 to .12, Columbium up to .5, Molybdenum up to 2.0, and the balance Nickel. The alloy of the current invention may also comprise impurities and incidental elements generally associated with nickel-based alloys, such as Zirconium up to .06 and Boron up to .015 percent by weight.
BRIEE DESCRIPTION OF THE DRAWINGS
Figure 1 is a graph of rupture time t, in hours, versus stress a, in KSI (N/mm2), for both conventionally cast (CC) and directionally solidified (DS) specimens of IN-738, showing the effect of directional solidification on the stress rupture strength of IN-738.
Figure 2 is an isometric view of a gas turbine rotating blade.
Figure 3 is a bar chart of rupture time, t, in thousands of hours, versus stress ~, in KSI (N/mm2), showing the stress rupture life at 870C (1600F) and four stress levels of (i) four heats of the alloy according to the current invention, SAS1-SAS4, as directionally cast, (ii) conventionally cast IN-738 and (iii) directionally solidified IN-6203.
DESCRIPTION OF THE PREFERRED EMBODIMENT
Figure 1 shows a graph of time to rupture versus stress at three temperature levels for both conventionally cast -- that is, having an equia~ed grain structure -- and directionally solidified IN-738 test specimens. As can be seen, data at both the 870C (1600F) and 925C (1700~F) temperature levels indicates that the rupture life of the directionally solidified specimens is worse than that of the conventionally cast specimen above a stress level of about 275 N/mm2 (40 KSI)o There is shown in Figure 2 a rotating blade 1 used in the turbine section of a gas turbine. The current invention is directed to an alloy, referred to as OM 200, from which such blades may be cast, especially using a directional solidification casting process. Generally speaking, the current invention is directed to an alloy comprising the following elements in weight percent: Chromium 14.75 to 16.0, - 5 - 57,151 Cobalt 8.0 to 8.5, Aluminum 3.4 to 4.0, Titanium 3.4 to 4.3, Aluminum plus Titanium 7.7 to 8.3, Tantalum 1.75 to 2.7, Tungsten 2.0 to 4.0l Carbon .05 to .12, Columbium up to .5, Molybdenum up to 2.0, and the balance Nickel. The alloy of the current invention may also comprise impurities and incidental elements generally associated with nickel-based alloys, such as Zirconium up to .06 and Boron up to .015 percent by weight. Prefexably, the alloy of the current invention consists essentially of the following elements in weight percent: Chromium about 15.5, Cobalt about ~.0, Aluminum about 4.0, Titanium abou$ 3.8, Aluminum plus Titanium about 7O8~ Tantalum about 2.7, Tungsten about 2.6, Molybdenum 0.5, Carbon about 0.08 and the balance Nickel.
As previously discussed, both corrosion resistance and strength are important properties for blade alloys.
Corrosion resistance in nickel-based alloys is provided primarily by chromium. Nickel-based alloys used for gas turbine components are strengthened by three mechanisms -- (i~
solid solution strengthening, (ii) strengthening resulting from the presence of carbides and (iii) gamma prime strengthening. Solid solution strengthening is provided by molybdenuml chromium and tungsten and, to a lesser extent, by cobalt, iron and vanadium. Gamma prime strengthening is provided primarily by aluminum and titanium, which strengthen the austenitic matrix through the precipitation of Ni3(Al and/or Ti), an fcc intermetallic compound. In addition, the aluminum in gamma prime can be replaced by tantalum and columbium.
Since the strength of nickel-based super alloys increases with increasing gamma prime volume fraction up to about 60%, in the alloy according to the current invention the content of aluminum, titanium, tantalum and columbium, which tend to increase the gamma prime volume fraction, have been controlled so as to achieve a hi~h gamma prime volume fraction. The amount of gamma prime in an alloy can be determined as discussed below with respect to the determination of the electron vacancy number Nv. The amount 2Pi 6 57,151 of gamma prime, in weight percent, in four heats of the alloy according to the current invention shown in Table I are approximately 52, 54, 56 and 54, respectively. The aforementioned preferred composition of the alloy according to the curren-t invention has approxiamtely 56% gamma prime.
By comparison, the amount of gamma prime in IN-738 is appro~imately 50% by weisht. (It should be noted that the aforementioned variation in the gamma prime content among the four heats of the alloy according to the current invention did not adversely affect corrosion resistance or stability.) As a result of operation at elevated tempexature, nicXel-based alloys undergo microstructural chang~s. Such changes include gamma prime coarsening, which ad~ersely affects the strength of the alloy, and the transformation of gamma prime into unwanted topologically close-packed secondary phases, such as plate or needle-like sigma, eta, etc. The formation of these plate-like phases adversely affects both strength and toughness. Consequently, in order to ensure that high strength and toughness are maintained for many thousands of hours of operation at elevated temperature~ the composition of the strengthening elements must be carefully balanced, as explained below, so that the alloy has a certain degree of microstructural stability.
It is well known that the electron vacancy number Nv is a measure of the microstructural stability of the alloy.
The higher the ~alue of Nv the greater the tendency for the formation of the aforementioned topologically close-packed secondary phases -- see for example, the American Society of Metal Handbook, 9th ed., volume 4, page 278. As disclosed in the brochure entitled ~Alloy IN-738, Technical Data,"
published by The International Nickel Company and hereby incorporated by reference in its entirety, the electron vacancy number may be calculated from the equation Nv = 0 66 Ni + 1.71 Co + 2.66 Fe + 4.66 (Cr + Mo + W) + 6.~6 Zr. The atomic percent of each element in the matri~ to be substituted into this equation is determined by convexting the composition from weight percen~ to atomic percent and assuming that (i) - 7 - 57,151 one-half the carbon forms MC in the order of TaC, CbC, TiC, ~ii) the remaining carhon forms M23C6 with the M comprising twenty three atoms of Cr, (iii) boron is combined as Mo3B2, (iii) gamma prime is Ni3(Al, Ti, Ta, Cb), and (iv) the residual matrix consists of the atomic percent minus those atoms contained in the carbides, the boride and the gamma prime reaction so that the total of the remaining atomic percentages gives the atomic concentration in the matrix.
Conversion of these atomic percentages to a 100 percent basis gives the atomic percentage of each element to be substituted into the equation above. Since it is desirable to maintain a low value of Nv, in the alloy according to the current invention the composition is adjusted so that, in addition to obtaining high strength, the electron vacancy number of the alloy does not exceed about 2.4. In the preferred composition of the alloy the electron vacancy numbex is equal to about 2.4.
As previously discussed, the aluminum in gamma prime can be replaced by columbium and/or tantalum, as well as titanium. However, since replacement of Al by Ti, Cb or ~a in Ni3Al adversely affects the misfit between the austenitic matrix and the gamma prime particles, the stability of these compounds in order of decreasing stability is Ni3Al, Ni3Ti and Ni3Cb(orTa). As a result, the titanium/aluminum ratio plays a major role in gamma prime coarsening. The titanium/aluminum ratio also plays a major role in the transformation of gamma prime into the aforementioned unwanted plate-like phases.
Accordingly, in order to minimize gamma prime coarsening and to improve the stabilit~ of the gamma prime phase, the titanium/aluminum ratio is ~enerally maintained below 2:1 (by weight). The transformation of gamma prime into unwanted needle or plate-like phases can also be retarded by the addition of tungsten.
To assess the stability of the alloy according to the current invention, the shoulder and grip sections of creep rupture specimens tested at 871C (1600F) and 205.9 N/mm2 (30 XSI) were examined metallographically for sigma, eta and mu 2'~
phases. It was found that the aging of these specimens at 871C, under stress/ for up to 4000 hours did not produce any undesirable phases, thereby suggesting that the alloy of the current invention is stable.
In the alloy according to the current invention, the amounts of aluminum, titanium, tantalum, columbium and molybdenum have been balanced so as to attain high strength when the alloy is directionally solidified while maintaining good microstructural stability. Significantly, this result has been achieved without the need to reduce the chromium content, and, -therefore, without impairing corrosion xesistance, as had heretofore been thought necessary by those skilled in the art. Specifically, in the alloy according to the current invention, the content of aluminum and titanium has been increased, when compared with IN-738, to a minimum of 3.4% for each, with the minimum combined aluminum plus titanium content being 7.7%. The maximum amounts of aluminum and titanium are 4.0% and 4.3%, respectively, with the maximum aluminum plus titanium content being 8.3%. In addition, the amounts of columbium and molybdenum have been reduced so that the optimum composition includes no columbium and only 0.5%
molybdenum. At most only 0.5% columbium and 2.0~ molybdenum are permitted in the alloy of the current invention.
Chromium, however, has been maintained in the range of 14.75%
to 16%, so that adequate corrosion resistance is maintained.
Tungsten is maintained in the range of 2.0% to 4.0% and tantalum in the range of 1.75% to 2.7%. Zirconium and boron are limited to 0.06% and 0.015~, respectively, with none of either of these elements being present in the most preferred composition. Carbon is maintained in the 0.05% to 0.12%
range. Moreover, as previously discussed, in the alloy according to the current invention, the elements are adjusted wîthin the aforementioned ranges so that the eleckron vacancy number is maintained at no more than about 2.4, thereby ensuring that adequate microstructural stability is achieved.
As a result of the foregoing, the strength of the alloy according to the current invention when cast by a 9 ~ 57,151 directional solidification process is high, despite its relatively high chromium content. Significantly, good microstructural stability of the alloy according to the current invention has been achieved by considerably reducing the levels of columbium and molybdenum, when compared to IN-738, so that the amounts of aluminum and titanium can be increased without driving the electron vacancy number too high.
Four heats of the alloy according to the current invention, designated SASl through SAS4, were prepared using the compositions in weight percent shown in Table I.
Specimens were cast from these heats 80 as to be directionally solidified and were given the standard heat treatment suggested for IN-738 -- that is, a 1121C (2050F~ solution heat treatment for two hours followed by a 843C (1550F) aging treatment for twenty four hours. These specimens were subjected to a variety of tests and the results compared to conventionally cast IN-738 and directionally solidified IN-6203. IN-6203 is a nickel-based alloy having a nominal composition in weight percent of Chromium 22.0, Cobalt 19.0, Aluminum 2.3, Titanium 3.5, Tantalum 1.10, Columbium 0.80:
Tungsten 2.00, Boron 0.01, Carbon 0.15, Zirconium 0.10 Hafnium 0.75 and the balance Nickel.
TABLE I
Element SASl SAS2 SAS3 AS4 Cr 16.00 16.00 14.75 16.00 Co 8.50 8.00 8.00 8.00 Al 3.40 3.40 3.60 4.00 Ti 4.30 4.30 4.30 3.40 Al+Ti 7.70 7.70 7.90 7.40 W 2.60 2.60 2.60 2.60 Mo 1.75 0.90 1.00 0.60 Ta 1.75 2.70 2.70 2.70 Cb 0.00 0.00 0.50 0.00 C 0.09 0.09 0.09 0.09 Zr 0.04 0.04 0.04 0.04 B 0.01 0.01 0.01 0.01 Ni Bal Bal Bal Bal The results of low fatigue testing at 871C (1600F) of specimens from heats SAS1, SAS3 and SAS4 are shown in Table II and indicate that the cyclic life of the alloy according - 10 - 57,151 to the current invention is superior to IN-738. The results of impact testing of three specimens from heat SASl -- taken from the tip, mid-section and bottom of th~ cast test bar, respectively - are shown in Table III and reveal that the impact strength of the alloy of the current invention is comparable to that of IN-738. Note that IN-738 is provided in both low and high carbon forms and the data is reported accordingly where appropriate. In its low carbon form the carbon content is reduced from 0.17% to 0.11% and zirconium is reduced from 0.10~ to 0.05%t be weight, compared to conventional high carbon IN-738.
The results of stress rupture tests are shown in Tables IV and V and Figure 3, wherein the rupture life of specimens from each of the four heats are compared with those of conventionally cast IN-738 and directionally solidified IN-6203. As can be seen, the data shows that the stress rupture life of the alloy according to the curren-t invention, when directionally solidified, is significantly superior to both that of both IN-738 and IN-6203 -- for example, at 69 N/mm2 (10 KSI), the alloy of the current invention can provide the same rupture life as IN-738 and IN-6203 at a temperature about 56C (100F) higher. When applied to a gas turbine blade, such increased metal temperature capability allows the turbine to operate at higher gas temperatures, ~hereby significantly improving its performance.
~ABLE I I
Total Strain Cycles to Fail~lre Range IN-738LC SAS1 SAS3 SAS4 0.012 - 1412 1442 1061 0.01 317 2990 3973 2852 0.0085 862 5342 6673 11200 0.0075 - 16220 50173+ 10504 0.007 1510 0.00~ 1819 0.005 2503 0.004 10153 - - -2. ~ ~
- 11 - 57~151 T~BLE III
Temp., C _ Impact Energy~ Joules IN-738HC IN-738LC SASl Tip Mid Bot 5 24 6.3 8.8 6.8 - 8.1 260 6.9 - 7.9 11.316.7 538 11.7 - 7.5 9.612.1 649 7.6 10.0 6.3 8.714.1 760 7.1 8.5 6.0 7.39.9 10871 8.1 9.8 9.6 9.114.4 Table IV
Temp., C Time to Rupture~ Hours @ 69 N/mm2 ~10 KSI) _N-738 IN-6203 SASl 968 2~96 982 1437 2752 8400+
996 820 1359 6450+
Table V
S$resæ Time to Rupture, Hours @ 871C ~
N/mm2 (KSI ) IN-738 IN-6203SASl SAS3 SAS4 275.8 (40)188 169 403 633 317 24~.3 (35)420 683 1165 1397 776 25206.9 (30) 1143 1~44 3272 4211 2676 172.4 (25)3177 - 8~00~8800~8247 Turbine blades cast from the alloy of the current invention are advantageously made by vacuum-induction melting and vacuum casting using a directional solidification process.
Directional solidification causes th~ grain boundaries to be oriented substantially parallel to the principal stress axis of the blade with almost no grain boundaries orient~d normal to the principal stress axis. Techniques for directional solidification are well known in the art -- see, for example, U.S. Patent Nos. 3,260,505 (Ver Snyder)l 3,494,709 (Piearcey) and 3,897,815 (Smashey), hereby incorporated by reference in their en~irety.
The gamma prime distribution depends on heat treatment, as well as composition. The standard heat treatment for nickel-based alloys such as IN-738 -~ i.e., a solution treatment followed by an aging treatment -- produces duplex gamma prime comprised of coarse, cuboidal primary gamma prime and fine, spherical gamma prime in approximately equal - 12 - 57,151 amounts. The coarse gamma prime is undesolved gamma prime that did not go into solution during the solution treatment.
Hence, the amount of coarse gamma prime present in the alloy depends on the degree by which the solution temperature is below the ga~a prime solvus temperature, at which all of the gamma prime goes into solution -- that is, the lower the solution temperature r the greater the amount of coarse gamma prime. The fine gamma prime forms during the aging treatment, the amount depending on the amount of gamma prime that did not go into solution during solution treatment.
Directionally solidified components are cooled at a slower rate than conventionally cast components. As a result, directional solidification produces coarser primary gamma prime so that the volume fraction of cuboidal gamma prime is higher and that of spherical gamma prime is lower than in directionally cast components. Unfortunately, a high volume fraction of coarse gamma prime has a deleterious effect on strength. This discrepancy in gamma prime distribution can be minimized by optimizing the heat treatment, particularly the solution treatment, based on the gamma prime solvus and incipient melting temperatures, both of which depend on alloy composition. Generally, it is desirable that the heat treatment be performed at a temperature between the solvus temperature and the incipient melting temperature. Hence, in or~er to obtain the full benefits of heat treatment, the solvus temperature should be considerably below the incipient melting temperature. Unfortunately, although aluminum, titanium, tantalum and columbium increase the gamma prime volume fraction, and therefore, strength, as previously discussed, they also have the effect of raising the solvus temperature and decreasing the incipient melting temperature, thereby narrowing the heat treatment window.
The gamma prime solvus and incipient melting temperatures for three melts of the alloy according to the current invention are shown in Table VI. These temperatures were determined using the differential thermal analysis and gradient bar -- i.e., metallography -- method, in which the - 13 - 57,151 bar was exposed to various temperatures in the 1066C (1950F) ~o 1427C t2300F) temperature range for four hours and then fan cooled. As can be seen, the solvus temperature varies from 1211C to 1229C. By comparison, the solvus temperature for IN-738 is approximately 1204C (2200F). As previously mentioned, IN-738 is typically given a 1121C (2050F) solution heat treatment for two hours followed by a 843C
(1550F) aging treatment for twenty four hours. Various heat treatment regimes for the alloy according to the current invention, based on its incipient melting and solvus temperatures, are discussed further below.
TABLE VI
SASl SAS3 SAS4 Incipient melting temp. C 1241 1229 1229 15Gamma prime solvus temp., C 1211 1211 1229 According to the current invention, the as~cast blades may be heat treated in any of four ways -- (i) solution treating for ~ hours at 1121C (2050F), followed by aging for 24 hours at 843C (1550F), (ii) solution treating for 4 hours at 1149C (2100F), followed by agins for 24 hours at 843C
(1550F), (iii) solution treating Eor 4 hours at 1204C
(2200F), resolution treating for 2 hours at 1121C (2050F) followed by aging for 24 hours at 843C (1550F), and (iv) solution treating for 4 hours a~ 1204C (2200F), resolution treating for 4 hours at 1149C (2100F), followed by aging for 24 hours at 843C (1550F). These heat treatment optimizes mechanical strength by creating a duplex gamma prime structure in a gamma matrix and discrete chrome-carbides at the grain boundaries. The use of higher solution temperatures will decrease the amount of coarse gamma prime and increase the amount of fine gamma prime, hence, further strengthening the alloy.
Although the current invention has been described with reference to an alloy for casting turbine blades, it should be understood that the alloy according to the current invention may also be used to form other components that require high strength at elevated temperature and good corrosion resistance. Accordingly, the present invention may ~ 182r~
- 14 - 57, 151 be embodied in other specific forms without departing from the spirit or essential attributes thereof and, accordingly, reference should be made to the appended claims, rather than to the foregoing specification, as indicating the scope of the invention.
Claims (12)
1. A gas turbine having a turbine blade made from an alloy comprising the following elements in weight percent:
Chromium 14. 75 to 16.0 Cobalt 8.0 to 8.5 Aluminum 3.4 to 4.0 Titanium 3.4 to 4.3 Aluminum plus Titanium 7 .7 to 8.3 Tantalum 1.75 to 2.7 Tungsten 2.0 to 4.0 Carbon .05 to .12 Nickel Balance.
Chromium 14. 75 to 16.0 Cobalt 8.0 to 8.5 Aluminum 3.4 to 4.0 Titanium 3.4 to 4.3 Aluminum plus Titanium 7 .7 to 8.3 Tantalum 1.75 to 2.7 Tungsten 2.0 to 4.0 Carbon .05 to .12 Nickel Balance.
2. The gas turbine according to claim 1, wherein said alloy has an electron vacancy number of no more than about 2.4.
3. The gas turbine according to claim 1, wherein said alloy further comprises the following elements in weight percent:
Columbium 0.0 to .5 Molybdenum 3.0 to 2.0 Zirconium 0.0 to .06 Boron 0.0 to .015.
Columbium 0.0 to .5 Molybdenum 3.0 to 2.0 Zirconium 0.0 to .06 Boron 0.0 to .015.
4. The gas turbine according to claim 1, wherein said blade is made from said alloy by a directional solidification process, whereby the grain boundaries within said blade are oriented substantially parallel.
- 16 - 57,151
- 16 - 57,151
5. A nickel-based alloy comprising the following elements in weight percent:
Chromium 14.75 to 16.0 Cobalt 8.0 to 8.5 Aluminum 3.4 to 4.0 Titanium 3.4 to 4.3 Aluminum plus Titanium 7.7 to 8.3 Tantalum 1.75 to 2.7 Tungsten 2.0 to 4.0 Carbon .05 to .12 Nickel Balance.
Chromium 14.75 to 16.0 Cobalt 8.0 to 8.5 Aluminum 3.4 to 4.0 Titanium 3.4 to 4.3 Aluminum plus Titanium 7.7 to 8.3 Tantalum 1.75 to 2.7 Tungsten 2.0 to 4.0 Carbon .05 to .12 Nickel Balance.
6. The alloy according to claim 5, wherein the electron vacancy number of said alloy is no more than about 2.4.
7. The alloy according to claim 5, wherein the weight percents of tungsten and carbon are about 2.6 and about 0.09, respectively.
8. The alloy according to claim 5, further comprising the following elements in weight percent:
Columbium 0.0 to .5 Molybdenum 0.0 to 2.0 Zirconium 0.0 to .06 Boron 0.0 to .015.
Columbium 0.0 to .5 Molybdenum 0.0 to 2.0 Zirconium 0.0 to .06 Boron 0.0 to .015.
9. The alloy according to claim 8, wherein the weight percent of tungsten is about 2.6, carbon is about 0.09, molybdenum is 0.60 to 1.75, zirconium is about 0.04 and boron is about 0.01.
10. The alloy according to claim 5, wherein the alloy comprises the following elements in weight percent:
Chromium about 15.5 Cobalt about 8.0 - 17 - 57,151 Aluminum about 4.0 Titanium about 3.8 Aluminum plus Titanium about 7.8 Tantalum about 2.6 Tungsten about 2.6 Carbon about 0.08 Molybdenum about 0.5 Nickel Balance.
Chromium about 15.5 Cobalt about 8.0 - 17 - 57,151 Aluminum about 4.0 Titanium about 3.8 Aluminum plus Titanium about 7.8 Tantalum about 2.6 Tungsten about 2.6 Carbon about 0.08 Molybdenum about 0.5 Nickel Balance.
11. A nickel-based alloy consisting of essentially of the following elements in weight percent:
Chromium 14.75 to 16.0 Cobalt 8.0 to 8.5 Aluminum 3.4 to 4.0 Titanium 3.4 to 4.3 Aluminum plus Titanium 7.7 to 8.3 Tantalum 1.75 to 2.7 Tungsten 2.0 to 4.0 Carbon .05 to .12 Columbium 0.0 to .5 Molybdenum 0.0 to 2.0 Zirconium 0.0 to .06 Boron 0.0 to .015 Nickel Balance.
Chromium 14.75 to 16.0 Cobalt 8.0 to 8.5 Aluminum 3.4 to 4.0 Titanium 3.4 to 4.3 Aluminum plus Titanium 7.7 to 8.3 Tantalum 1.75 to 2.7 Tungsten 2.0 to 4.0 Carbon .05 to .12 Columbium 0.0 to .5 Molybdenum 0.0 to 2.0 Zirconium 0.0 to .06 Boron 0.0 to .015 Nickel Balance.
12. The alloy according to claim 11, wherein the electron vacancy number of said alloy is no more than about 2.40.
Applications Claiming Priority (2)
| Application Number | Priority Date | Filing Date | Title |
|---|---|---|---|
| US85414292A | 1992-03-18 | 1992-03-18 | |
| US854,142 | 1992-03-18 |
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| CA2091827A1 true CA2091827A1 (en) | 1993-09-19 |
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|---|---|---|---|
| CA002091827A Abandoned CA2091827A1 (en) | 1992-03-18 | 1993-03-17 | Gas turbine blade alloy |
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| EP (1) | EP0561179A3 (en) |
| JP (1) | JPH0617171A (en) |
| KR (1) | KR930019844A (en) |
| CN (1) | CN1076508A (en) |
| AU (1) | AU3380093A (en) |
| CA (1) | CA2091827A1 (en) |
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| US5882586A (en) * | 1994-10-31 | 1999-03-16 | Mitsubishi Steel Mfg. Co., Ltd. | Heat-resistant nickel-based alloy excellent in weldability |
| JP2862487B2 (en) * | 1994-10-31 | 1999-03-03 | 三菱製鋼株式会社 | Nickel-base heat-resistant alloy with excellent weldability |
| FR2949234B1 (en) * | 2009-08-20 | 2011-09-09 | Aubert & Duval Sa | SUPERALLIAGE NICKEL BASE AND PIECES REALIZED IN THIS SUPALLIATION |
| US10227678B2 (en) | 2011-06-09 | 2019-03-12 | General Electric Company | Cobalt-nickel base alloy and method of making an article therefrom |
| US10266926B2 (en) | 2013-04-23 | 2019-04-23 | General Electric Company | Cast nickel-base alloys including iron |
| GB201309404D0 (en) * | 2013-05-24 | 2013-07-10 | Rolls Royce Plc | A nickel alloy |
| WO2018148110A1 (en) * | 2017-02-08 | 2018-08-16 | Borgwarner Inc. | New alloys for turbocharger components |
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| US3459545A (en) * | 1967-02-20 | 1969-08-05 | Int Nickel Co | Cast nickel-base alloy |
| US4152488A (en) * | 1977-05-03 | 1979-05-01 | United Technologies Corporation | Gas turbine blade tip alloy and composite |
| GB2024858B (en) * | 1978-07-06 | 1982-10-13 | Inco Europ Ltd | Hightemperature nickel-base alloys |
| CN1045607A (en) * | 1989-03-15 | 1990-09-26 | 中国科学院金属研究所 | A kind of method that improves the superalloy performance |
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1993
- 1993-02-23 EP EP19930102779 patent/EP0561179A3/en not_active Withdrawn
- 1993-02-23 TW TW082101251A patent/TW222017B/zh active
- 1993-02-26 AU AU33800/93A patent/AU3380093A/en not_active Abandoned
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| EP0561179A2 (en) | 1993-09-22 |
| KR930019844A (en) | 1993-10-19 |
| AU3380093A (en) | 1993-09-23 |
| EP0561179A3 (en) | 1993-11-10 |
| TW222017B (en) | 1994-04-01 |
| MX9301280A (en) | 1993-09-01 |
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