US3869284A - High temperature alloys - Google Patents

High temperature alloys Download PDF

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US3869284A
US3869284A US346815A US34681573A US3869284A US 3869284 A US3869284 A US 3869284A US 346815 A US346815 A US 346815A US 34681573 A US34681573 A US 34681573A US 3869284 A US3869284 A US 3869284A
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alloy
component
weight
alloys
gas turbine
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James French Baldwin
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FRENCH BALDWIN J
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FRENCH BALDWIN J
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Priority to US346815A priority Critical patent/US3869284A/en
Priority to GB2751073A priority patent/GB1395125A/en
Priority to IL44295A priority patent/IL44295A/en
Priority to CA193,974A priority patent/CA1021604A/en
Priority to FR7408475A priority patent/FR2223470B1/fr
Priority to DE2463065A priority patent/DE2463065C2/de
Priority to DE2463064A priority patent/DE2463064C2/de
Priority to DE2463066A priority patent/DE2463066C2/de
Priority to DE2415074A priority patent/DE2415074C2/de
Priority to SE7404368A priority patent/SE404380B/xx
Priority to JP3671474A priority patent/JPS5716180B2/ja
Priority to IT20814/74A priority patent/IT1012132B/it
Priority to IN1526/CAL/74A priority patent/IN142527B/en
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Priority to US05/598,111 priority patent/USRE28681E/en
Priority to JP56183590A priority patent/JPS5816047A/ja
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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
    • C22C19/055—Alloys based on nickel or cobalt based on nickel with chromium and Mo or W with the maximum Cr content being at least 20% but less than 30%

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  • This invention relates to the addition to iron and alloys of iron, hereinafter all referred to as ferrous metals," of an alloying addition, and particularly an alloying addition for the purpose of improving the machinability of the ferrous metals.
  • the alloying addition is intended to improve the machinability of the ferrous metal, it is important that the alloying addition be dispersed as uniformly as possible throughout the ferrous metal ingot so as to improve the machinability substantially uniformly throughout the ingot as well as avoiding the Iocalised segregation which would be liable to impair the mechanical properties of the ferrous metal.
  • the fumes generated when the alloying addition is added to ferrous metals may be toxic, in which case the fumes may be efficiently removed throughout the casting process.
  • the alloying addition may, for example, include one, some or all of the elements lead, selenium and tellurium. in the form of pure metals or in the form of alloys, or as mineral compounds.
  • An object of the invention is the provision of an improved method of adding the alloying addition to ferrous metals, whereby a high recovery of the alloying addition is obtained in the ingot and an ingot is produced in which the alloying addition is present in the form of a uniformly distributed and finely divided mi cro-dispersion, with a reduction in the segregation of the alloying addition as compared with most commonly used methods so that a reduced proportion of the ingot need be discarded.
  • a further object of the invention is the provision of an improved method as described above wherein any toxic fumes generated by the alloying addition may be efficiently and easily removed.
  • a method of adding alloying addition to a ferrous metal comprising the steps of; passing a gas through the ferrous metal, the ferrous metal being heated to a temperature above its customary teeming temperature, the passage ofthe gas being such that turbulence is created in the ferrous metal and a turbulent zone at the surface thereof; simultaneously adding said alloying addition to the ferrous metal at or adjacent to the turbulent zone whilst removing any toxic fumes generated; continuing to pass the gas through the ferrous metal until the temperature of the metal reaches its customary teeming temperature; and'teemin'g the metal into moulds whilst removing any toxic fumes generated by the alloying addition.
  • the ferrous metal may be tapped from a furnace into a tapping ladle and the gas may be passed through the ferrous metal whilst the metal is in the tapping ladle.
  • the alloying addition is added slowly and in a finely divided form.
  • the gas is an inert gas such as argon which will not affect the composition of the metal.
  • ferrous metal as used herein includes alloy and other steels and the method of this invention is particularly applicable to ferrous metals of which the carbon content does not exceed 2.0 percent.
  • Lead has negligible solid solubility in steel, but is slightly soluble in liquid steel; the solubility of lead increases as the temperature of the steel rises.
  • fumes generated when lead is added to steel are highly toxic and must be efficiently removed throughout the casting process.
  • FIG. I is a diagrammatic side elevation, partly in section, of a furnace and ladle for use in carrying out the method of the present invention
  • FIG. 2 is a diagrammatic cross-sectional view. to an enlarged scale, of the ladle of FIG. 1 showing it in operative position with a fume hood,
  • FIG. 3 is a diagrammatic cross-sectional view of the apparatus shown in FIG. 2 but with the cross-section taken so as to show the argon inlet and pressure gun for adding lead,
  • FIG. 4 is a view, to an enlarged scale, of the part of FIG. 3 enclosed in the circle marked 4,
  • This invention relates to nickel-base alloys having relatively great tensile strength at high temperatures and to castings made from such alloys.
  • the nickel-base superalloys of the present invention are particularly useful for fabricating components of gas turbine engines, such as turbine blades, turbine vanes, integral wheels, and the like.
  • thermal fatigue properties are associated with intermediate temperature (1,3OF l,500F) ductility.
  • alloys with high temperature rupture and creep strength have inadequate thermal fatique and hot corrosion resistance.
  • alloys with good hot corrosion resistance show poor high temperature rupture, creep and thermal fatigue properties.
  • Superalloys suitable for fabricating gas turbine components desirably possess good creep-rupture strength, i.e., resist excessive creep or rupture for long periods of time while under stress at high temperatures. Such alloys also desirably possess good creep-rupture ductility. i.e., deform uniformly and predictably while under stress at high temperature, rather than crack and fracture. Alloys that lack ductility will tolerate little deformation before the onset of crack nucleation, rapid crack propagation, and failure. Use of a material lacking adequate ductility can result in unpredicatable and catastrophic engine component failure.
  • a characteristic peculiar to the gamma prime strengthened superalloys is that they are subject to a sharp decrease in creep-rupture ductility and tensile strength at temperatures between about 1,300F. and l,500F.
  • the decrease in ductility is commonly referred to as the ductility trough," as ductility is higher at temperatures below l,300F. and above l,50ll)F.
  • MAR-M200 U.S. Pat. No. 3,164,465 This alloy possesses adequate strength for most advanced gas turbine engine requirements, but lack of l,400F. ductility in the conventionally cast material precludes its usefulness for turbine components.
  • the alloys of the present invention have improved high temperature strength and corrosion resistance. These alloys are capable of withstandingprolonged op eration at temperatures up to about 2,000F. or higher, and may be'formed into highly advantageous castings.
  • alloy compositions have been discovered which possess unique and unusually high creep-rupture strength and ductility in the polycrystalline (non-directionally solidified) form.
  • a previously unrecognized criticality has been discovered in the amounts of two alloying elements (boron and carbon) included in chromium, aluminum, and titanium containing nickel base superalloy compositions.
  • Boron is considered an essential ingredient in superalloys.
  • boron in the form ofcomplex borides is also located at grain boundaries. Grain boundary morphology of superalloys is significant because high temperature creep and rupture failures initiate at and propogate along grain boundaries. Complex borides at grain boundaries reduce the onset of grain boundary tearing under rupture loading.
  • Typical cast superalloys of the prior art preferably contain carbon in an amount of about 0.10% to about 0.25% by weight.
  • the carbon content range is between about 0.03% and about 0.15% by weight.
  • the carbon content is kept as low as 0.05% by weight.
  • the present invention is based, in part, on the discovery of an unusual and unexpected improvement, in both 1,400F. creep-rupture strength and ductility of gamma prime strengthened nickel-base superalloys, obtained by increasing boron content up to about twenty times the accepted optimum level. Maintenance of the boron content within this critical range of the present invention not only eliminates the problem discussed earlier, relating to the ductility trough present at temperatures between about l,300 F. and l,500 F., but results in a marked increase in creep-rupture strength at those temperatures.
  • alloys of the prior art which will exhibit enhanced properties by following the teachings of the present invention are those disclosed in U.S. Pat. Nos. 3,310,399; 3,164,465; 3,061,426; and 3,619,182. While many of the alloy compositions disclosed in these patents are similar to, and generically overlap with, the alloys of the present invention, none of these patents disclose, nor do corresponding commercial alloy have. the unusual and surprisingly advantageous properties and characteristics of the alloys of the present invention. This is because the prior art fails to recognize the critical carbon and boron content ranges of the alloys of the present invention. All of the commercial alloys derived from the patents referred to above contain substantially less than the minimum boron content used in the alloys of the present invention. Additionally, while at least some of these patents suggest broad boron content ranges which overlap the boron content range of the present invention, there is no recognition that high temperature properties will maximize in a narrow range within these broadly disclosed ranges.
  • the alloys of the present invention which have very good stress rupture life at elevated temperatures, contain required minimum amounts of nickel, chromium, aluminum and titanium.
  • the chromium affords primary corrosion resistance while the remaining components are essential to the formation of the gamma prime intermetallic compound, Ni (Al, Ti), which forms the basic superalloy structure of this invention.
  • Ni (Al, Ti) precipitate lends to these alloys their required high temperature strength, and titanium is an important element in providing the strength properties of the present alloys at both room temperature and at elevated temperatures.
  • the presence of significant amounts of Ti strengthener in the present alloys renders them significantly different in character from lower temperature alloys such as those of U.S. Pat. No. 3,005,704, which excludes Ti from its alloys.
  • the present invention pertains to gamma prime phase strengthened superalloys. These alloys are specifically adapted to be employed in cast shapes under conditions of high stress at high temperature. The invention also concerns cast components for use in gas turbine engines made from such alloys.
  • the alloys of the present invention are predominantly nickel, i.e., at least 35% nickel, and contain in varying amounts, chromium, aluminum, titanium, and boron.
  • One or more ofthe elements carbon, cobalt, zir conium, molybdenum, tantalum, rhenium, columbium, vanadium, and tungsten may also be included in these alloys.
  • the alloys of the present invention may contain minor amounts of other elements ordinarily included in superalloys by those skilled in the art which will not substantially deleteriously effect the important characteristics of the alloy or which are inadvertently included in such alloys by virtue of impurity levels in commercial grades of alloying ingredients.
  • the carbon content of the alloys is maintained below about 0.05% by weight. By additionally maintaining the carbon content below this critical upper limit, it is possible to effect creep-rupture strength and ductility improvement at temperatures around 1,400F. while, at the same time, maintaining or improving creep-rupture strength and ductility at temperatures around l,800 F.
  • Table l sets forth a broad range and two different narrower ranges, in terms of percent by weight, of elements employed in the alloys of the present invention. It should be understood that the tabulation in Table 1 relates to each element individually and is not intended to solely define composite of broad and narrow ranges.
  • a particularly preferred alloy composition in percentages by weight, consists essentially of about 8.0% to about 10.25% chromium, about 4.75 to about 5.5% aluminum, about 1.0% to about 2.5% titanium, about 0.05 to about 0.30% (and more preferably about 0.075% to about 0.2%) boron, up to about 0.17% (and more preferably less than 0.05%) carbon, about 8% to about 12% cobalt, about 0.75% to about 1.8% columbium, about 11% to about 16% tungsten, up to 0.20% zirconium, and the balance essentially nickel and minor amounts of impurities and incidental elements which do not detrimentally affect the basic characteristics of the alloy.
  • FIG. 4 is a reproduction of a photomicrograph at a magnification of 300, ofa commercial alloy outside the ambit of the present invention.
  • FIG. 5 is a reproduction of a photomicrograph, at a magnification of 300, of an alloy (comparable to the alloy ofFlG. 4) within the ambit of the present inventlon.
  • FIG. 6 is a reproduction of a photomicrograph, at a magnification of 7,000, of the same alloy shown in the photomicrograph of FIG. 4.
  • FIG. 7 is a reproduction of a photomicrograph, at a magnification of 7,000, of the same alloy shown in the photomicrograph of FIG. 5.
  • ELEMENT BROAD RANGE NARROWER RANGES Another particularly preferred alloy composition, in percentages by weight, consists essentially of about 7.5% to about 8.5% chromium, about 5.75% to 6.25% aluminum, about 0.8% to about 1.2% titanium, about 0.05% to about 0.30% (and more preferably about 0.075% to about 0.2%) boron, up to about 0.13% (and more preferably less than 0.05%) carbon, about 9.5% to about 10.5% cobalt, about 5.75% to about 6.25% molybdenum, about 4.0% to about 4.5% tantalum, 0.05% to 0.10% zirconium, and the balance essentially nickel and minor amounts of impurities and incidental elements which do not detrimentally affect the basic characteristics of the alloy.
  • Impurities and incidental elements which may be present in the alloys of the present invention include manganese, copper, and silicon in amounts of not more than 0.50%, sulfur and phosphorus in amounts of not more than 0.20%, and iron in amounts of not more than 2.0%. Impurities such as nitrogen, hydrogen, tin, lead, bismuth, calcium, and magnesium should be held to as low a concentration as practical.
  • FIG. 1 is a graphical plot of percent creep elongation against time for two alloys, one within the ambit of the present invention, and the other outside the ambit of the present invention.
  • FIG. 2 is a plot of creep-rupture life in hours against the boron content in weight percent of certain nickel base alloys at both 1,400 F. 94,000 psi and 1,800 F. 29,000 psi.
  • FIG. 3 is similar to FIG. 2 but plots percent creep DESCRIPTION OF EXAMPLES AND PREFERRED EMBODIMENTS
  • the alloys of the present invention containing boron within the critical range of 0.05% to 0.3% by weight. exhibit enhanced creep rupture strength and ductility in the 1,300 F. to l,500 F. temperature range over prior art gamma prime strengthened nickel base superalloys.
  • the alloys of the present invention are capable of withstanding an applied stress of 94,000 psi at 1,400F without rupture for a time in excess of hours. Further, this improvement in strength and ductility properties in the intermediate temperature range (1,300F. 1,500F) is accompanied by a pronounced beneficial effect on high temperature (above 1,700F.)
  • Alloys of the present invention having improved intermediate temperature strength and ductility, demonstrate great advantage in resistance to high temperature thermal fatigue cracking over alloys containing boron in amounts outside the critical range of the present invention.
  • Usual designs involve the mechanical attachment of turbine blades around the periphery of a wheel or disk which rotates at high speed.
  • hot gases pass over the airfoil portion of the blades, causing the blades and disk to rotate at high speed.
  • the hot gases raise the metal temperatures and the high rotational speed of the disk imposes stress due to centrifugal loading.
  • the attachment, or root portion of the blade is heated only to moderate temperatures due to the cooling effect of the massive disk.
  • the temperature to which the root section of the blade is heated is frequently in the ductility trough temperature range (1,300 F. to 1,500 E). It is an essential mechanical property of an alloy being used for such blades that it be capable of deforming predictably in the root section at temperatures around l,400 F. while withstanding mechanically imposed strain without cracking, i.e., the
  • alloy must possess reasonable ductility.
  • the alloys of the present invention containing boron within the critical range of 0.05% to 0.30% by weight, demonstrate great advantage in strength and ductility in the l,400 F. temperature range over prior art alloys intended for use in turbine blades.
  • the rotating turbine disk to which the blade root is attached, also requires high resistance to creep and rupture along with ductility and strength to resist fatigue and crack propagation. Accordingly, the alloys described herein provide enhanced properties desirable in disk alloys.
  • the alloys ofthe present invention enhance the over all performance of integral wheels.
  • example alloys compositionally similar to the commercial alloys of Table 11. but containing boron within the critical range of the present invention, were prepared. Analyses ofthese example alloys (designated A-l, B-l, etc.) are presented in Table 111. Standard cast-to-size test bars (0.25 inch in diameter) of the alloys of Table 11 and the example alloys of Table 111 were prepared by melting and casting under vacuum into shell molds. All example alloy specimens were heat treated under a protective atmosphere at l,975 F. for four hours and then air cooled. The example alloys were also subjected to an aging heat treatment at 1,650 F. for ten hours. Each of the commercial alloys of Table 11 was heat treated in accordance with the practice recommended by the alloy developer.
  • Table IV shows the comparative creep-rupture strength (as measured by timeto rupture) and ductility, (as measured by prior creep) of both commercial alloys A, B, C, and E, and example alloys A-l, B-l, C-l, C-2, C-3 and E-l. All alloys were tested at l,400 F. under a stress of 94,000 psi.
  • Table IV shows a very significant improvement in both l,400 F. creep-rupture strength and ductility for alloys having a boron content within the critical range of the present invention. At 0.20 weight percent boron, the properties of Example C-3, although decreasing from Example C-2, still show a marked improvement over Alloy C.
  • the low carbon aspect of the present invention is particularly important with respect to turbine components requiring enhanced properties at both 1,400 F. and 1,800 F.
  • properties at around l,400 F. are particularly important with respect to the root sections of turbine blades.
  • turbine blades desirably require an alloy having good high temperature properties throughout the temperature range of from about 1,300 F. to about 1.900 F. or higher.
  • example alloys A-2, B-2, (-4 through 13, D-l. E-2 through 9, and F- 1 were prepared by melting under vacuum. Standard test bars (0.25 inch diameter) were cast under vacuum into shell molds and all specimens were heat treated under a protective atmosphere at 1,975 F. for four hours. After air cooling, all specimens were subjected to an aging heat treatment of l,650 F. for ten hours. Analyses for series A,B,D, and F example alloys are shown in Table V. Analyses for series C and E example alloys are shown, respectively, in Tables 1V and VII.
  • Creep-rupture tests were conducted at 1.800 F. under a stress of 29,000 psi and at l.400 F. under a stress of 94,000 psi on all low carbon example alloys. For comparative purposes, the same tests were conducted on the commercial alloys A,B.C,D,E. and F of Table [1. The commercial alloy test bars were heat treated in accordance with the procedures recommended by the producers to achieve maximum mechanical properties. Creep-rupture data for commerical alloys D and F under these conditions were obtained from technical literature provided by the respec- .tive alloy producers.
  • the data of Table VII demonstrates the applicability of the present invention to a wide range of superalloys.
  • the four example alloys corresponding to the four commercial alloys designated A,B,D, and F had boron and carbon levels approaching the target compositions, i.e., 0.01 weight percent carbon and 0.10 to 0.12 weight percent boron.
  • the comparative test results between the commercial alloys A, B, D, and F and corresponding series A, B, D and F example alloys set forth in Table VIII shows in all cases that very significant improvements are effected at both 1,400 F. and 1,800 F in rupture life and ductility.
  • Example alloy C-4 shows very good EXAMPLE rupture life at 1,400 E, but the low level of both boron A4 [)4 F4 and carbon causes low ductility in the l,800 F. test.
  • the combination oflow boron and low carbon 8 822 9-22 contents causes poor castability and a tendency for c6 13,15 11,76 10. 943 20 castings to crack on cooldown during solidification.
  • w 9.66 4.18 M0 m9 310 2.43 204 The minimum boron required to clrcumvent these Ta 3. 53/0 "50 423 problems in the low carbon alloys 15 about 0.05 welght Ti 0.96 0.99 1.38 3.69 percent.
  • the improvement over commercial alloy E is significant.
  • the l,800 F. properties are maintained within a boron range of about 0.05 to 0.15 weight percent. At 0.22 weight percent boron in Example alloy E9. the l.800 F. strength is about sixty percent that of commercial alloy E.
  • Example Alloy C-7 cast-to-size test bars were subjected for 1,000 hours and examined microstructurally. No deleterious phase formation was observed and subsequent creep-rupture testing was conducted at 1.400F and 94,000 psi for comparison with the same to creep testing at 1,500F under a stress of 40,000 psi TABLE XII-Continued Creep-Rupture Properties I500F exposure for I000 hours under stress of 40,000psi alloy in the as-heat treated condition. Results shown in Table XII reveal essentially no change in rupture life and an improvement in l,400F ductility.
  • FIG. 1 shows the creep characteristics of typical Alloy C and one of the example alloy C-7 test bars in the l,400F test.
  • percent creep elongation is plotted against time. The improved results obtained with the alloys of the present invention are dramatically demonstrated.
  • FIGS. 2 and 3 further demonstrate the critical relationship between boron content and strength and ductility.
  • FIG. 2 is a plot of creep rupture life in hours against the boron content in weight percent ofC series, low carbon (less than 0.05% by weight), alloys at both l,400F 94,000 psi and l',800F 29,000 psi.
  • the creep rupture life for commercial alloy C at 1,400F 94,000 psi and 1,800F 29,000 psi for commercial alloy C is noted on the plot at, respectively, points A and B.
  • substantial improvements in creep rupture life are obtained at l,400F by maintaining the boron content within the critical range of the present invention.
  • FIG. 3 is a plot of percent creep elongation against boron content for C series, low carbon alloys at both l,400F 94,000 psi and 1,800F 29,000 psi.
  • the percent creep elongation for commercial alloy C at both l,400F 94,000 psi and l,800F 29,000 psi is also noted on this plot, respectively, at points A and B. Again substantial improvements are apparent at l,400F. with respect to alloys containing boron within the critical range of the present invention. While the percent creep elongation obtained at 1,800F with alloys within the ambit of the present invention is not as high as that ofthe commercial alloy, highly acceptable levels are achieved.
  • FIG. 4 shows the normal microstructure of commercial Alloy C in the as-cast condition at 300 magnifications.
  • the light etching dendrite arms or branch-like areas indicate tungsten segregation.
  • a few titanium rich carbides are visible in the lower center portion of the photomicrograph.
  • FIG 5 shows the profound microstructural change resulting from the added boron and reduced carbon of example alloy C-7. Reducing carbon to less than 0.02 weight percent frees titanium previously tied up as a stable carbide. The increased available titanium in the alloy results in the formation of gamma-gamma prime eutectic in the grain boundaries, a microstructural effect known to enhance l,400F ductility.
  • the boron addition results in the formation of discrete grain bondary particles, identified by electron-beam micro-probe analysis as an M 8 type boride where M (in the C alloy series) is chromium and tungsten. These grain boundary particles are responsible for restoring l,800F creep-rupture ductility to low carbon alloys.
  • FIGS. 6 and 7 Electron photomicrographs of commercial alloy C and example alloy C-7, at 7,000 magnifications, are shown, respectively, in FIGS. 6 and 7.
  • FIG. 6 shows, as previously stated to be the general case, borides located at the grain boundaries.
  • a boride precipitate within each gamma prime particle may be observed, a phenomenon absent in superalloys of the more conventional compositions.
  • the presence of the very fine boride particles appears to retard dislocation movement through the gamma-prime particles and, in
  • alloys of the present invention may be extruded and hot forged. Wrought, high strength nickel-base superalloys are generally employed in applications where ductility and fracture toughness in the 1,000F to 1,500F temperature range are'of prime concern. Such applications include gas turbine engine turbine and compressor disks.
  • the series E alloys of the present invention may be hot forged, using conventional techniques, into shaped articles having the characteristics considered to be essential in advanced wrought alloys. For example, alloys E-l and E-5 have responded very satisfactorily to extrusion and forging in the 2,000F to 2,200F temperature range in anticipation of the requirements for advanced wrought disk and blade materials.
  • the present invention also anticipates the use of powder metallurgy for controlling the size, morphology and distribution of the boride microconstituents previously described.
  • a nickel base alloy for use at relatively high temperatures consisting essentially of the following elements in the weight percent ranges set forth:
  • a cast component for use in a gas turbine engine formed of the alloy of claim 3.
  • a shaped object of the alloy of claim 1 capable of withstanding an applied stress of 94,000 psi at 1,400 F. without rupture for a time in excess of 120 hours.
  • a shaped object of the alloy of claim 1 capable of withstanding an applied stress of 29,000 psi at 1,800 F. without rupture for a time in excess of 40 hours.
  • the alloy of claim 1 which contains, on a weight basis, about 6.0% to about 17% chromium, about 2% to about 8% aluminum, about 0.75% to about 3% titanium, about 2% to about 17% cobalt, and about 40% to 80% by weight nickel.
  • a cast component for use in a gas turbine engine formed of the alloy of claim l4.
  • the alloy of claim 1 which contains, on a weight basis, about 5% to 12% chromium, about 4% to about 8% aluminum. about 0.75% to about 2.5% titanium, about 5% to about 15.5% cobalt, and about 40% to 80% by weight nickel.
  • a nickel base alloy for use at relatively high temperatures consisting essentially of the following elements in the weight percent ranges set forth:
  • nickel being present in an amount of from about 40% to by weight.
  • a nickel base alloy for use at relatively high temperatures consisting essentially of the following elements in the weight percent ranges set forth:
  • the balance of the alloy being essentially nickel and minor amounts of impurities and incidental elements which do not detrimentally affect the basic characteristics of the alloy.
  • a shaped object of the alloy of claim 27 capable of withstanding an applied stress of 94,000 psi at 1,400" F. without rupture for a time in excess of hours.
  • a shaped object of the alloy of claim 29 capable of withstanding an applied stress of 29,000 psi at l.800 F. without rupture for a time in excess of 40 hours.
  • a nickel base alloy for use at relatively high tem peratures consisting essentially of the following elements in the weight percent ranges set forth:
  • Elements Percent 65 the balance of the alloy being essentially nickel and minor amounts of impurities and incidental elements which do not detrimentally affect the basic characteristhe balance of the alloy being essentially nickel and minor amounts of impurities and incidental elements which do not detrimentally affect the basic characteristics of the alloy, said nickel being present in an amount of from about 40% to 80% by weight.
  • a nickel base alloy for use at relatively high temperatures consisting essentially of the following elements in the weight percent ranges set forth:
  • the balance of the alloy being essentially nickel and minor amounts of impurities and incidental elements which do not detrimentally affect the basic characteristics of the alloy.
  • a shaped object of the alloy of claim 41 capable of withstanding an applied stress of 94,000 psi at 1.400 F. without rupture for a time in excess of hours.
  • a shaped object of the alloy of claim 43 capable of,withsta'nding an applied stress of 29,000 psi at l,800 F. without rupture for a time in excess of 40 hours.
  • a nickel base alloy for use at relatively high temperatures consisting essentially of the following elements in the weight percent ranges set forth:

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US346815A 1973-04-02 1973-04-02 High temperature alloys Expired - Lifetime US3869284A (en)

Priority Applications (15)

Application Number Priority Date Filing Date Title
US346815A US3869284A (en) 1973-04-02 1973-04-02 High temperature alloys
GB2751073A GB1395125A (en) 1973-04-02 1973-06-08 High temperature alloys
IL44295A IL44295A (en) 1973-04-02 1974-02-26 Nickel-based alloys for high temperatures
CA193,974A CA1021604A (en) 1973-04-02 1974-03-04 High temperature nickel based alloys
FR7408475A FR2223470B1 (nl) 1973-04-02 1974-03-13
DE2415074A DE2415074C2 (de) 1973-04-02 1974-03-28 Verwendung einer Superlegierung auf Nickelbasis zur Herstellung von Gasturbinenteilen
DE2463064A DE2463064C2 (de) 1973-04-02 1974-03-28 Verwendung einer Superlegierung auf Nickelbasis zur Herstellung von Gasturbinenteilen
DE2463066A DE2463066C2 (de) 1973-04-02 1974-03-28 Verwendung einer Superlegierung auf Nickelbasis zur Herstellung von Gasturbinenteilen
DE2463065A DE2463065C2 (de) 1973-04-02 1974-03-28 Verwendung einer Superlegierung auf Nickelbasis zur Herstellung von Gasturbinenteilen
SE7404368A SE404380B (sv) 1973-04-02 1974-04-01 Nickelbaserad gamma-primforsterkt hogtemperaturlegering med forbettrad kryphallfasthet och duktilitet mellan 700-980Ÿc
JP3671474A JPS5716180B2 (nl) 1973-04-02 1974-04-02
IT20814/74A IT1012132B (it) 1973-04-02 1974-04-08 Leghe per alte temperature
IN1526/CAL/74A IN142527B (nl) 1973-04-02 1974-07-08
US05/598,111 USRE28681E (en) 1973-04-02 1975-07-29 High temperature alloys
JP56183590A JPS5816047A (ja) 1973-04-02 1981-11-16 ニツケル基合金

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GB (1) GB1395125A (nl)
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US4108647A (en) * 1975-07-17 1978-08-22 The International Nickel Company, Inc. Alloys of nickel, chromium and cobalt
US4127410A (en) * 1976-03-24 1978-11-28 The International Nickel Company, Inc. Nickel based alloy
US4610736A (en) * 1983-03-23 1986-09-09 The United States Of America As Represented By The Administrator Of The National Aeronautics And Space Administration Nickel base coating alloy
US4629521A (en) * 1984-12-10 1986-12-16 Special Metals Corporation Nickel base alloy
GB2198143A (en) * 1986-11-28 1988-06-08 Korea Advanced Inst Sci & Tech Heat resistant nickel alloy.
US5156808A (en) * 1988-09-26 1992-10-20 General Electric Company Fatigue crack-resistant nickel base superalloy composition
US6521175B1 (en) * 1998-02-09 2003-02-18 General Electric Co. Superalloy optimized for high-temperature performance in high-pressure turbine disks
US20040187973A1 (en) * 2003-03-24 2004-09-30 Noritaka Takahata Nickel base heat resistant cast alloy and turbine wheels made thereof
US20040208777A1 (en) * 2001-09-18 2004-10-21 Jacinto Monica A. Burn-resistant and high tensile strength metal alloys
US20100008790A1 (en) * 2005-03-30 2010-01-14 United Technologies Corporation Superalloy compositions, articles, and methods of manufacture
US20110194971A1 (en) * 2004-12-02 2011-08-11 Hiroshi Harada Heat-resistant superalloy
EP3086899B1 (en) 2013-12-24 2020-04-15 Liburdi Engineering Limited Precipitation strengthened nickel based welding material for fusion welding of superalloys

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GB1544720A (en) * 1977-01-13 1979-04-25 Inco Europ Ltd Nickel-base superalloys
CA1117320A (en) * 1977-05-25 1982-02-02 David N. Duhl Heat treated superalloy single crystal article and process
US4207098A (en) 1978-01-09 1980-06-10 The International Nickel Co., Inc. Nickel-base superalloys
DE2830396A1 (de) * 1978-07-11 1980-01-24 Inco Europ Ltd Nickel-chrom-superlegierung
BE880399A (fr) * 1978-12-18 1980-04-01 United Technologies Corp Article en superalliage a base de nickel et procede de fabrication
US4292076A (en) * 1979-04-27 1981-09-29 General Electric Company Transverse ductile fiber reinforced eutectic nickel-base superalloys
US4222794A (en) * 1979-07-02 1980-09-16 United Technologies Corporation Single crystal nickel superalloy
IL65677A0 (en) * 1981-06-12 1982-08-31 Special Metals Corp Nickel base cast alloy
DD231225A3 (de) * 1982-12-28 1985-12-24 Mai Edelstahl Verwendung einer warmverformbaren aushaertbaren nickellegierung fuer warmverschleissbestaendige warmarbeitswerkzeuge
US4731221A (en) * 1985-05-06 1988-03-15 The United States Of America As Represented By The United States Department Of Energy Nickel aluminides and nickel-iron aluminides for use in oxidizing environments
US4814023A (en) * 1987-05-21 1989-03-21 General Electric Company High strength superalloy for high temperature applications
US5087305A (en) * 1988-07-05 1992-02-11 General Electric Company Fatigue crack resistant nickel base superalloy
JP2778705B2 (ja) * 1988-09-30 1998-07-23 日立金属株式会社 Ｎｉ基超耐熱合金およびその製造方法
DE4323486C2 (de) * 1992-07-23 2001-09-27 Abb Research Ltd Ausscheidungshärtbare Nickelbasis-Superlegierung und Verwendung der Legierung als Werkstoff bei der Herstellung eines gerichteten erstarrten Bauteils, wie insbesondere einer Gasturbinenschaufel
US6354799B1 (en) * 1999-10-04 2002-03-12 General Electric Company Superalloy weld composition and repaired turbine engine component
CN102162049B (zh) * 2011-04-07 2012-12-19 上海大学 一种超超临界汽轮机用镍基合金材料及其制备方法
CN103540801A (zh) * 2013-10-17 2014-01-29 常熟市良益金属材料有限公司 一种高温合金
CN103789576B (zh) * 2014-01-15 2016-03-02 常州大学 一种高晶界强度镍基合金及其制备方法
DE102017129218A1 (de) * 2017-12-08 2019-06-13 Vdm Metals International Gmbh Schweisszusatzwerkstoff
EP3572541B1 (en) 2018-05-23 2023-05-17 Rolls-Royce plc Nickel-base superalloy
WO2024058101A1 (ja) * 2022-09-14 2024-03-21 株式会社プロテリアル 熱間鍛造用金型およびその製造方法

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US3061426A (en) * 1960-02-01 1962-10-30 Int Nickel Co Creep resistant alloy
US3164465A (en) * 1962-11-08 1965-01-05 Martin Metals Company Nickel-base alloys
US3310399A (en) * 1964-07-10 1967-03-21 Baldwin James French Alloys for use at high temperatures
US3486887A (en) * 1964-01-31 1969-12-30 Nat Res Inst Metals Nickel base heat-resisting alloy

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US3155501A (en) * 1961-06-30 1964-11-03 Gen Electric Nickel base alloy

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US3061426A (en) * 1960-02-01 1962-10-30 Int Nickel Co Creep resistant alloy
US3164465A (en) * 1962-11-08 1965-01-05 Martin Metals Company Nickel-base alloys
US3486887A (en) * 1964-01-31 1969-12-30 Nat Res Inst Metals Nickel base heat-resisting alloy
US3310399A (en) * 1964-07-10 1967-03-21 Baldwin James French Alloys for use at high temperatures

Cited By (18)

* Cited by examiner, † Cited by third party
Publication number Priority date Publication date Assignee Title
US4108647A (en) * 1975-07-17 1978-08-22 The International Nickel Company, Inc. Alloys of nickel, chromium and cobalt
US4127410A (en) * 1976-03-24 1978-11-28 The International Nickel Company, Inc. Nickel based alloy
US4610736A (en) * 1983-03-23 1986-09-09 The United States Of America As Represented By The Administrator Of The National Aeronautics And Space Administration Nickel base coating alloy
US4629521A (en) * 1984-12-10 1986-12-16 Special Metals Corporation Nickel base alloy
GB2198143A (en) * 1986-11-28 1988-06-08 Korea Advanced Inst Sci & Tech Heat resistant nickel alloy.
GB2198143B (en) * 1986-11-28 1990-09-05 Korea Advanced Inst Sci & Tech Heat resistance ni-cr-w-al-ti-ta alloy
US5156808A (en) * 1988-09-26 1992-10-20 General Electric Company Fatigue crack-resistant nickel base superalloy composition
US6521175B1 (en) * 1998-02-09 2003-02-18 General Electric Co. Superalloy optimized for high-temperature performance in high-pressure turbine disks
US20100266442A1 (en) * 2001-09-18 2010-10-21 Jacinto Monica A Burn-resistant and high tensile strength metal alloys
US20040208777A1 (en) * 2001-09-18 2004-10-21 Jacinto Monica A. Burn-resistant and high tensile strength metal alloys
US20040187973A1 (en) * 2003-03-24 2004-09-30 Noritaka Takahata Nickel base heat resistant cast alloy and turbine wheels made thereof
US20110194971A1 (en) * 2004-12-02 2011-08-11 Hiroshi Harada Heat-resistant superalloy
US8734716B2 (en) * 2004-12-02 2014-05-27 National Institute For Materials Science Heat-resistant superalloy
US20100008790A1 (en) * 2005-03-30 2010-01-14 United Technologies Corporation Superalloy compositions, articles, and methods of manufacture
US20100158695A1 (en) * 2005-03-30 2010-06-24 United Technologies Corporation Superalloy Compositions, Articles, and Methods of Manufacture
US8147749B2 (en) 2005-03-30 2012-04-03 United Technologies Corporation Superalloy compositions, articles, and methods of manufacture
EP3086899B1 (en) 2013-12-24 2020-04-15 Liburdi Engineering Limited Precipitation strengthened nickel based welding material for fusion welding of superalloys
EP3086899B2 (en) † 2013-12-24 2023-05-31 Liburdi Engineering Limited Precipitation strengthened nickel based welding material for fusion welding of superalloys

Also Published As

Publication number Publication date
FR2223470A1 (nl) 1974-10-25
IN142527B (nl) 1977-07-23
IT1012132B (it) 1977-03-10
DE2463064C2 (de) 1984-06-14
GB1395125A (en) 1975-05-21
SE404380B (sv) 1978-10-02
JPS49128818A (nl) 1974-12-10
JPS5816047A (ja) 1983-01-29
DE2463065C2 (de) 1984-09-06
IL44295A (en) 1976-08-31
CA1021604A (en) 1977-11-29
DE2415074A1 (de) 1974-10-17
DE2415074C2 (de) 1983-12-15
DE2463066C2 (de) 1984-07-05
FR2223470B1 (nl) 1978-12-01
JPS5716180B2 (nl) 1982-04-03
IL44295A0 (en) 1974-05-16

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