US3920489A - Method of making superalloy bodies - Google Patents

Method of making superalloy bodies Download PDF

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US3920489A
US3920489A US473407A US47340774A US3920489A US 3920489 A US3920489 A US 3920489A US 473407 A US473407 A US 473407A US 47340774 A US47340774 A US 47340774A US 3920489 A US3920489 A US 3920489A
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nickel
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Edward R Buchanan
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General Electric Co
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General Electric Co
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Priority to NLAANVRAGE7102685,A priority Critical patent/NL171309C/en
Priority to DE2109874A priority patent/DE2109874C3/en
Priority to FR7107147A priority patent/FR2084089A5/fr
Priority to US00120289A priority patent/US3821783A/en
Priority to GB2288671A priority patent/GB1318832A/en
Priority to US00402306A priority patent/US3850702A/en
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Priority to US483837A priority patent/US3920492A/en
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    • H—ELECTRICITY
    • H01—ELECTRIC ELEMENTS
    • H01L—SEMICONDUCTOR DEVICES NOT COVERED BY CLASS H10
    • H01L29/00—Semiconductor devices specially adapted for rectifying, amplifying, oscillating or switching and having potential barriers; Capacitors or resistors having potential barriers, e.g. a PN-junction depletion layer or carrier concentration layer; Details of semiconductor bodies or of electrodes thereof ; Multistep manufacturing processes therefor
    • H01L29/02—Semiconductor bodies ; Multistep manufacturing processes therefor
    • H01L29/04—Semiconductor bodies ; Multistep manufacturing processes therefor characterised by their crystalline structure, e.g. polycrystalline, cubic or particular orientation of crystalline planes
    • H01L29/045—Semiconductor bodies ; Multistep manufacturing processes therefor characterised by their crystalline structure, e.g. polycrystalline, cubic or particular orientation of crystalline planes by their particular orientation of crystalline planes
    • B—PERFORMING OPERATIONS; TRANSPORTING
    • B22—CASTING; POWDER METALLURGY
    • B22F—WORKING METALLIC POWDER; MANUFACTURE OF ARTICLES FROM METALLIC POWDER; MAKING METALLIC POWDER; APPARATUS OR DEVICES SPECIALLY ADAPTED FOR METALLIC POWDER
    • B22F3/00—Manufacture of workpieces or articles from metallic powder characterised by the manner of compacting or sintering; Apparatus specially adapted therefor ; Presses and furnaces
    • B22F3/24—After-treatment of workpieces or articles
    • H—ELECTRICITY
    • H01—ELECTRIC ELEMENTS
    • H01L—SEMICONDUCTOR DEVICES NOT COVERED BY CLASS H10
    • H01L21/00—Processes or apparatus adapted for the manufacture or treatment of semiconductor or solid state devices or of parts thereof
    • H—ELECTRICITY
    • H01—ELECTRIC ELEMENTS
    • H01L—SEMICONDUCTOR DEVICES NOT COVERED BY CLASS H10
    • H01L21/00—Processes or apparatus adapted for the manufacture or treatment of semiconductor or solid state devices or of parts thereof
    • H01L21/02—Manufacture or treatment of semiconductor devices or of parts thereof
    • H01L21/02104—Forming layers
    • H01L21/02365—Forming inorganic semiconducting materials on a substrate
    • H01L21/02367—Substrates
    • H01L21/0237—Materials
    • H01L21/02373—Group 14 semiconducting materials
    • H01L21/02381—Silicon, silicon germanium, germanium
    • H—ELECTRICITY
    • H01—ELECTRIC ELEMENTS
    • H01L—SEMICONDUCTOR DEVICES NOT COVERED BY CLASS H10
    • H01L21/00—Processes or apparatus adapted for the manufacture or treatment of semiconductor or solid state devices or of parts thereof
    • H01L21/02—Manufacture or treatment of semiconductor devices or of parts thereof
    • H01L21/02104—Forming layers
    • H01L21/02365—Forming inorganic semiconducting materials on a substrate
    • H01L21/02367—Substrates
    • H01L21/02433—Crystal orientation
    • H—ELECTRICITY
    • H01—ELECTRIC ELEMENTS
    • H01L—SEMICONDUCTOR DEVICES NOT COVERED BY CLASS H10
    • H01L21/00—Processes or apparatus adapted for the manufacture or treatment of semiconductor or solid state devices or of parts thereof
    • H01L21/02—Manufacture or treatment of semiconductor devices or of parts thereof
    • H01L21/02104—Forming layers
    • H01L21/02365—Forming inorganic semiconducting materials on a substrate
    • H01L21/02518—Deposited layers
    • H01L21/02521—Materials
    • H01L21/02524—Group 14 semiconducting materials
    • H01L21/02532—Silicon, silicon germanium, germanium
    • H—ELECTRICITY
    • H01—ELECTRIC ELEMENTS
    • H01L—SEMICONDUCTOR DEVICES NOT COVERED BY CLASS H10
    • H01L29/00—Semiconductor devices specially adapted for rectifying, amplifying, oscillating or switching and having potential barriers; Capacitors or resistors having potential barriers, e.g. a PN-junction depletion layer or carrier concentration layer; Details of semiconductor bodies or of electrodes thereof ; Multistep manufacturing processes therefor
    • Y—GENERAL TAGGING OF NEW TECHNOLOGICAL DEVELOPMENTS; GENERAL TAGGING OF CROSS-SECTIONAL TECHNOLOGIES SPANNING OVER SEVERAL SECTIONS OF THE IPC; TECHNICAL SUBJECTS COVERED BY FORMER USPC CROSS-REFERENCE ART COLLECTIONS [XRACs] AND DIGESTS
    • Y10—TECHNICAL SUBJECTS COVERED BY FORMER USPC
    • Y10S—TECHNICAL SUBJECTS COVERED BY FORMER USPC CROSS-REFERENCE ART COLLECTIONS [XRACs] AND DIGESTS
    • Y10S117/00—Single-crystal, oriented-crystal, and epitaxy growth processes; non-coating apparatus therefor
    • Y10S117/901—Levitation, reduced gravity, microgravity, space
    • Y10S117/902—Specified orientation, shape, crystallography, or size of seed or substrate
    • Y—GENERAL TAGGING OF NEW TECHNOLOGICAL DEVELOPMENTS; GENERAL TAGGING OF CROSS-SECTIONAL TECHNOLOGIES SPANNING OVER SEVERAL SECTIONS OF THE IPC; TECHNICAL SUBJECTS COVERED BY FORMER USPC CROSS-REFERENCE ART COLLECTIONS [XRACs] AND DIGESTS
    • Y10—TECHNICAL SUBJECTS COVERED BY FORMER USPC
    • Y10S—TECHNICAL SUBJECTS COVERED BY FORMER USPC CROSS-REFERENCE ART COLLECTIONS [XRACs] AND DIGESTS
    • Y10S148/00—Metal treatment
    • Y10S148/049—Equivalence and options
    • Y—GENERAL TAGGING OF NEW TECHNOLOGICAL DEVELOPMENTS; GENERAL TAGGING OF CROSS-SECTIONAL TECHNOLOGIES SPANNING OVER SEVERAL SECTIONS OF THE IPC; TECHNICAL SUBJECTS COVERED BY FORMER USPC CROSS-REFERENCE ART COLLECTIONS [XRACs] AND DIGESTS
    • Y10—TECHNICAL SUBJECTS COVERED BY FORMER USPC
    • Y10S—TECHNICAL SUBJECTS COVERED BY FORMER USPC CROSS-REFERENCE ART COLLECTIONS [XRACs] AND DIGESTS
    • Y10S148/00—Metal treatment
    • Y10S148/097—Lattice strain and defects
    • Y—GENERAL TAGGING OF NEW TECHNOLOGICAL DEVELOPMENTS; GENERAL TAGGING OF CROSS-SECTIONAL TECHNOLOGIES SPANNING OVER SEVERAL SECTIONS OF THE IPC; TECHNICAL SUBJECTS COVERED BY FORMER USPC CROSS-REFERENCE ART COLLECTIONS [XRACs] AND DIGESTS
    • Y10—TECHNICAL SUBJECTS COVERED BY FORMER USPC
    • Y10S—TECHNICAL SUBJECTS COVERED BY FORMER USPC CROSS-REFERENCE ART COLLECTIONS [XRACs] AND DIGESTS
    • Y10S148/00—Metal treatment
    • Y10S148/115—Orientation
    • Y—GENERAL TAGGING OF NEW TECHNOLOGICAL DEVELOPMENTS; GENERAL TAGGING OF CROSS-SECTIONAL TECHNOLOGIES SPANNING OVER SEVERAL SECTIONS OF THE IPC; TECHNICAL SUBJECTS COVERED BY FORMER USPC CROSS-REFERENCE ART COLLECTIONS [XRACs] AND DIGESTS
    • Y10—TECHNICAL SUBJECTS COVERED BY FORMER USPC
    • Y10S—TECHNICAL SUBJECTS COVERED BY FORMER USPC CROSS-REFERENCE ART COLLECTIONS [XRACs] AND DIGESTS
    • Y10S148/00—Metal treatment
    • Y10S148/118—Oxide films
    • Y—GENERAL TAGGING OF NEW TECHNOLOGICAL DEVELOPMENTS; GENERAL TAGGING OF CROSS-SECTIONAL TECHNOLOGIES SPANNING OVER SEVERAL SECTIONS OF THE IPC; TECHNICAL SUBJECTS COVERED BY FORMER USPC CROSS-REFERENCE ART COLLECTIONS [XRACs] AND DIGESTS
    • Y10—TECHNICAL SUBJECTS COVERED BY FORMER USPC
    • Y10S—TECHNICAL SUBJECTS COVERED BY FORMER USPC CROSS-REFERENCE ART COLLECTIONS [XRACs] AND DIGESTS
    • Y10S438/00—Semiconductor device manufacturing: process
    • Y10S438/973—Substrate orientation

Definitions

  • Superalloys are heat resistant materials having superior strength and oxidation resistance at high tempera tures. Many of these alloys contain iron, nickel or C- balt alone or in combination as the principal element together with chromium to impart surface stability, and usually containing one or more minor constituents, such as molybdenum, tungsten, columbium, titanium and aluminum for the purpose of effecting strengthening. The physical properties of the superalloys make them particularly useful in the manufacture of gas turbine components.
  • the strength of superalloys is determined in part by their grain size. At low temperatures, fine grained equiaxed structures are preferred. At high temperatures (generally above 1600F.), large grain size structures are usually found to be stronger than fine-grained. This is believed related to the fact that failure generally originates at grain boundaries oriented perpendicular to the direction of the induced stress.
  • An improved technique for cast superalloys used in gas turbine engines was developed by Ver Snyder, U.S. Pat. No. 3,260,505 which discloses the preparation of a blade having an elongated columnar structure with unidirectional crys tals aligned substantially parallel to the long axis of the blade.
  • a method of making superalloy bodies characterized by having aligned elongated grains or a monocrystalline grain structure by employing powder metallurgy techniques.
  • the method involves hot compacting a nickel-base superalloy powder to form a relatively dense solid, solution heat treating the solid at a temperature below the incipient melting point, subjecting the material to a sufficient strain below the recrystallization temperature to permit subsequent recrystallization, and unidirectionally recrystallizing the strained material in a temperature gradient at a maximum temperature below the incipient melting temperature and above the 'y solvus temperature, to form a body having an elongated columanar grain or monocrystalline grain structure.
  • the product produced by my process is substantially similar in macrostructure to the cast superalloy articles prepared by directional solidification.
  • inhomogeneities such as eutectic modules are avoided and it is possible to prepare a material having a uniform structure throughout without this segregation characteristic of cast bodies.
  • Nickel-base superalloys are strong, high temperature materials which are particularly useful in gas turbine engines. A substantial listing of these materials is set positions and Rupture Strengths 0f Superalloys, ASTM Data Series Publication No. DS9E, and may be represented by the nominal compositions in weight percent of the following superalloys:
  • the nickel-base superalloy is in the form of a fine metal powder which is prepared in such a way that each powder particle is substantially of the same nominal composition as the final alloy composition.
  • a conventional technique for preparing such powders is by atomization of a melt of the alloy.
  • a powder ensures alloy homogeneity and overcomes the problems resulting from alloy segregation which occurs in large ingot sections and causes variations in physical properties within a single large part or from separate parts made from the same ingot.
  • powder materials to disperse a chemically inert phase, e.g., alumina or yttria, uniformly through the alloy by various milling techniques toachieve additional high temperature strength.
  • a chemically inert phase e.g., alumina or yttria
  • the next step involves hot compacting the metal powder into a dense solid.
  • the hot compaction is performed either by extrusion or by hot pressing. It is preferred that during hot compacting, a protective atmosphere or vacuum be used to prevent oxidation of some of the reactive elements in the alloy.
  • the alloy powder 3 may be extruded by canning it in a steel jacket and then hot extruding the billet to finished size or to stock which is machined to the desired final dimensions.
  • the dense solid consists essentially of a 'y' precipitate phase with a 'y matrix.
  • the dense solid is annealed to dissolve a substantial portion of the 'y phase.
  • the reason for the anneal is related to the fact that the 7 phase appears to impede elongated grain growth.
  • the annealing temperature should be above the y solvus and below the incipient melting temperature of the alloy.
  • Rene 120 has a 'y solvus temperature of about l205C. and an incipient melting temperature of about 1260C.
  • the annealing temperature is about l240C.
  • the annealing time is dependent on the size of the workpiece. I have found that -20 minutes at this temperature is preferred in a workpiece less than /2 inch thick.
  • the critical strain is defined as that amount of strain which is just sufficient to cause the growth of very large grains during subsequent recrystallization.
  • the crux of the critical strain concept is that a certain minimum strain is required to cause recrystallization during subsequent heating. If this strain is exceeded, the recrystallized grain diameter is essentially inversely related to the amount of tensile strain.
  • the critical strain in most of the nickel-base superalloys used in this invention are on the order of l-3percent at room temperature. This amount of plastic strain may be introduced in a tensile machine at a strain rate of 0.02 in./in./min.
  • the desired structure may also be achieved by rolling a test piece at room temperature to 2percent total reduction in thickness.
  • the state of critical strain can also be achieved by straining the workpiece at any temperature below the recrystallization temperature, although larger amounts of strain are required at higher temperature due to dynamic recovery during straining.
  • the critical strain at l200-l400F. is typically about 8-10percent.
  • the material is unidirectionally recrystallized to provide a body having an elongated parallel grain or monocrystalline structure.
  • This is performed by drawing the material through a gradient furnace. I have found the number of grains in the crosssection is essentially related to the efficiency of the gradient. In the preferred embodiment of the invention, the gradient is at least l000F./inch. The maximum temperature of the thermal gradient is above the 'y' solvus and below the incipient melting temperature of the alloy in question. The rate at which the workpiece is drawn through the gradient depends on the alloy, but it has been found that speeds of /z2inch/hour results in an elongated grain structure in most superalloys. I have found that slower speeds in this range result in monocrystalline structures, while higher speeds result in elongated parallel columnar grains.
  • EXAMPLE 1 A billet was prepared from Rene 120 nickel-base superalloy powders having the composition shown in the table above, except the carbon level was 0.05 percent rather than the typical 0.17 percent.
  • the loose powders, having mesh sizes +200, were encapsulated in a 3%. inch diameter stainless steel capsule having a 0.216 inch wall thickness.
  • the capsule cavity and powder were evacuated to 10* Torr, heated to 500C. under vacuum to remove volatile impurities, cooled to room temperature and sealed.
  • the entire capsule was then heated to 1175 C. for 2hours and extruded through a die aperture of 0.6 inch X 1.0 inch, approximately an 18/1 reduction. Two and one half inch lengths were cut from the billet. Four tabs, each 0.6 inch X 2- /2 inch X 0.072 inch were cut from the center of each length. Tabs were machined into tensile specimens having a gauge length of 0.150 inch X 0.072 inch X 1.0 inch.
  • the specimens were then subjected to various combinations of a prior anneal followed by being subjected to a strain at room temperature. Thereafter, the specimens were passed through a gradient fumace having a maximum temperature of 1260C., which is slightly below the alloy incipient melt temperature but above the 'y' solvus temperature. The temperature gradient was about l093C./in.
  • variable speed anneals were used for some specimens, in which about inch length of gauge section was passed through the hot zone of the furnace at a predetermined speed of about fiiinch/hn; then the drive motor speed was increased to about /2 inch/hr. for another of gauge length, and so on. This determined for a given set of processing conditions whether elongated grain growth would or would not occur.
  • EXAMPLE II A billet of Rene 120 was prepared using the same procedure as Example I, the only difference being the billet contained a normal carbon level, 0.17 percent. It was observed the combination of a heat treatment for 15 minutes at l240C. followed by a strain at room temperature of about 2 percent total strain resulted upon gradient annealing in an elongated grain structure in all attempts tried.
  • EXAMPLE III A billet of Rene 120 was prepared using the same procedure as Example I, except that the powder contained 2 vol.pct. Al O which had been milled into the metal powder by a mechanical milling process and the extrusion temperature was 1232C.
  • a billet of Rene 80 was prepared from superalloy powders having the composition shown in the table above.
  • the loose powders, having 325 mesh size, were encapsulated in a 2-1/16 inch carbon steel extrusion can, evacuated, sealed, and extruded through a k inch round die aperture (approximately an 18/ l reduction) at 1093C.
  • Lengths of the extruded bar were solution annealed for 2 hours at l2l0C. and cooled to room temperature in an air blast. Microstructural examination revealed a fine-grained structure having an average diameter of approximately 20 microns.
  • Tensile specimens having a A inch gauge diameter and 1 inch gauge length were machined from the /2 inch diameter stock.
  • Four specimens were deformed in tension at room temperature to respective strains of 2.3, 3.0, 3.5, and 10 percent at a strain rate of 0.02 inlinlmin.
  • the four specimens were surface ground to 0.1 inch thick flats to insure uniform heating during the gradient anneal.
  • Specimens were then subjected to a 50 percent l-lCl-SO percent HNO acid solution for & hour to remove potential grain nuclei at the specimen surface.
  • the two remaining specimens were machined.
  • the second reduced gauge sections were machined to dimensions of 0.50 inch long X 0.072 thick X 0.060 inch wide, so that the stress in the second gauge section was forty percent of that in the first gauge section. Test results were:
  • An article of manufacture having a chemically inert phase dispersed therein comprising a unidirectionally recrystallized nickel-base superalloy body free of eutectic module inhomogenieties having an aligned elongated grain or monocrystalline structure and containing up to 10 percent by volume of chemically chamically inert phase dispersed therein wherein said nickel-base superalloy body is prepared by a method comprising the steps of:
  • the strained material is unidirectionally recrystallized by drawing the material through a gradient furnace having a temperature gradient of at least 1000F. per inch and the material is drawn at a rate of about 0.5-2 inches per hour.

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Abstract

A method of making superalloy bodies characterized by having aligned elongated grains is provided by employing powder metallurgy procedures. The method involves hot compacting a nickel-base superalloy to a dense solid, heat treating the solid to form an essentially single phase gamma structure, imparting a critical strain to the single phase material and then subjecting the material to a unidirectional recrystallization step to form an elongated columnar grained structure having grain boundaries substantially parallel to the direction of recrystallization.

Description

United States Patent 1191 1111 3,9
Buchanan [4 Nov. 18, 1975 METHOD OF MAKING SUPERALLOY 3,671,223 6/1972 Thompson at al. 75/171 BODIES 3,783,032 1/1974 Walker et al. 75/171 75 l t Ed d B h B 1 nven or y R anan umt Hills Primary ExaminerW. Stallard Attorney, Agent, or Firm-F. Wesley Turner; Joseph Asslgnee! General Electric p y T. Cohen; Jerome C. Squillaro Schenectady, NY.
[21] Appl' 473407 A method of making superalloy bodies characterized Related US, Appli ti D t by having aligned elongated grains is provided by em- [62] Division of Ser. No. 402,306. Oct. 1, 1973, Pat. No. P' Powder melzlnurgy Procedures The method 3,850,701 involves hot compactmg a nickel-base superalloy to a I dense solid, heat treating the solid to form an essen- [52] US. Cl. 148/32 tiany Single Phase 7 Structure imparting a critical 51 Int. Cl. c221 1/10 Strain the Single P material and subjecting [58] Field f g 75/171; 143/115 R 115 F, the material to a unidirectional recrystallization step 148/32 to form an elongated columnar grained structure having grain boundaries substantially parallel to the direc- [56] References Cited tion of recrystallization.
Reichman 75/171 SUPERALLOY POWDER H07 COMPACT 016' CR/ T/CAL SUDJECT/NG STRAIN UNID/RCT/ONILLY RECRYS TALL IZ ING SUPERALLO) POWDER HOT CDMPAC T //V6 HEA T TREA TING SUBJECT/N6 T0 CR/T/CAL ST/PA/IV U/V/D/RECT/O/VALL Y RECR Y5 TALL lZ/N6 CONSOLIDATED SUPERALLOY BODY METHOD OF MAKING SUPERALLOY BODIES This is a division, of application Ser. No. 402,306, filed Oct. 1,1973, Now U.S. Pat. No. 3,850,702.
Superalloys are heat resistant materials having superior strength and oxidation resistance at high tempera tures. Many of these alloys contain iron, nickel or C- balt alone or in combination as the principal element together with chromium to impart surface stability, and usually containing one or more minor constituents, such as molybdenum, tungsten, columbium, titanium and aluminum for the purpose of effecting strengthening. The physical properties of the superalloys make them particularly useful in the manufacture of gas turbine components.
The strength of superalloys is determined in part by their grain size. At low temperatures, fine grained equiaxed structures are preferred. At high temperatures (generally above 1600F.), large grain size structures are usually found to be stronger than fine-grained. This is believed related to the fact that failure generally originates at grain boundaries oriented perpendicular to the direction of the induced stress. An improved technique for cast superalloys used in gas turbine engines was developed by Ver Snyder, U.S. Pat. No. 3,260,505 which discloses the preparation of a blade having an elongated columnar structure with unidirectional crys tals aligned substantially parallel to the long axis of the blade. This procedure involves directional solidification whereby almost a complete elimination of grain boundaries normal to the primary stress axis occurs. A further advance was made by Piearcey, U.S. Pat. No. 3,494,709 wherein grain boundaries in superalloys were eliminated by making single crystal castings. These directionally solidified materials are not suitable for all applications. One disadvantage is that there is a substantial increase in cost over conventional castings. Further, I have found that variations in mechanical properties of a directionally solidified part occur at different distances from the chill due to differences in dendrite spacing and microstructure along the length of the casting.
An alternative processing technique for superalloys that has aroused recent interest is powder metallurgy. This has advantages for some applications from the point of view of cost reduction and improved properties and permits introduction of an inert refractory oxide for additional strengthening. Thus, U.S. Pat. No. 3,639,179 issued to Reichman et al. describes a process for making nickel-base superalloys having superior high temperature properties which employs powder metallurgy techniques. The process involves confining and densifying the powder into a billet, cold working the billet below the recrystallization temperature, recrystallizing the cold worked billet for a time sufficient to nucleate new grains and thereafter heat treating the recrystallized billet to effect growth of the grains to a desired equiaxed size. This process yielded a nickelbase superalloy characterized as being of large grain size and possessing superior tensile strength and stress rupture life at elevated temperatures.
In accordance with the present invention, I have discovered a method of making superalloy bodies characterized by having aligned elongated grains or a monocrystalline grain structure by employing powder metallurgy techniques. The method involves hot compacting a nickel-base superalloy powder to form a relatively dense solid, solution heat treating the solid at a temperature below the incipient melting point, subjecting the material to a sufficient strain below the recrystallization temperature to permit subsequent recrystallization, and unidirectionally recrystallizing the strained material in a temperature gradient at a maximum temperature below the incipient melting temperature and above the 'y solvus temperature, to form a body having an elongated columanar grain or monocrystalline grain structure. The product produced by my process is substantially similar in macrostructure to the cast superalloy articles prepared by directional solidification. In addition, inhomogeneities such as eutectic modules are avoided and it is possible to prepare a material having a uniform structure throughout without this segregation characteristic of cast bodies.
The accompanying drawing, which is a flow sheet of the novel process, while not intended as a definition essentially illustrates the invention. A full discussion is set forth herein below.
Nickel-base superalloys are strong, high temperature materials which are particularly useful in gas turbine engines. A substantial listing of these materials is set positions and Rupture Strengths 0f Superalloys, ASTM Data Series Publication No. DS9E, and may be represented by the nominal compositions in weight percent of the following superalloys:
TABLE I Rene Rene Rene 1N- Udimet Ingredient 738 500 Si 0.30 0.75 Cr 14.0 9.5 9.25 16.0 19.0 Ni Bal. Bal. Bal. Bal. Bal. Co 9.5 15.0 10.0 8.5 18.0 M0 4.0 3 .0 2.0 1.75 4.0 W 4.0 7.0 2.6
Cb 0.9 Ti 5.0 4.20 4.0 3.4 2.9 A1 3.0 5.50 4.25 3.4 2.9 B 0015 0.015 0.015 0.01 0.005 Zr 0.03 0.06 0.05 0.10 Fe 0.2 1.0 max. 0.50 4.0 Other 1.0 V 3.75 Ta 1.75 Ta Initially the nickel-base superalloy is in the form of a fine metal powder which is prepared in such a way that each powder particle is substantially of the same nominal composition as the final alloy composition. A conventional technique for preparing such powders is by atomization of a melt of the alloy. The use of a powder ensures alloy homogeneity and overcomes the problems resulting from alloy segregation which occurs in large ingot sections and causes variations in physical properties within a single large part or from separate parts made from the same ingot. In addition, it is possible with powder materials to disperse a chemically inert phase, e.g., alumina or yttria, uniformly through the alloy by various milling techniques toachieve additional high temperature strength. These inert phases tend to agglomerate when added to liquid cast metal, thus preventing their utilization in cast metals.
The next step involves hot compacting the metal powder into a dense solid. The hot compaction is performed either by extrusion or by hot pressing. It is preferred that during hot compacting, a protective atmosphere or vacuum be used to prevent oxidation of some of the reactive elements in the alloy. The alloy powder 3 may be extruded by canning it in a steel jacket and then hot extruding the billet to finished size or to stock which is machined to the desired final dimensions. At this point, in the absence of an addition of an inert phase, the dense solid consists essentially of a 'y' precipitate phase with a 'y matrix.
Thereafter, the dense solid is annealed to dissolve a substantial portion of the 'y phase. The reason for the anneal is related to the fact that the 7 phase appears to impede elongated grain growth. For each alloy, the annealing temperature should be above the y solvus and below the incipient melting temperature of the alloy. For example, Rene 120 has a 'y solvus temperature of about l205C. and an incipient melting temperature of about 1260C. In the preferred embodiment involving Rene 120, the annealing temperature is about l240C. The annealing time is dependent on the size of the workpiece. I have found that -20 minutes at this temperature is preferred in a workpiece less than /2 inch thick.
The critical strain is defined as that amount of strain which is just sufficient to cause the growth of very large grains during subsequent recrystallization. The crux of the critical strain concept is that a certain minimum strain is required to cause recrystallization during subsequent heating. If this strain is exceeded, the recrystallized grain diameter is essentially inversely related to the amount of tensile strain. By just exceeding the critical strain in the workpiece and drawing it through a thermal gradient, a monocrystalline or elongated grain structure results. The critical strain in most of the nickel-base superalloys used in this invention are on the order of l-3percent at room temperature. This amount of plastic strain may be introduced in a tensile machine at a strain rate of 0.02 in./in./min. However, the desired structure may also be achieved by rolling a test piece at room temperature to 2percent total reduction in thickness. In addition, the state of critical strain can also be achieved by straining the workpiece at any temperature below the recrystallization temperature, although larger amounts of strain are required at higher temperature due to dynamic recovery during straining. The critical strain at l200-l400F. is typically about 8-10percent.
After straining, the material is unidirectionally recrystallized to provide a body having an elongated parallel grain or monocrystalline structure. This is performed by drawing the material through a gradient furnace. I have found the number of grains in the crosssection is essentially related to the efficiency of the gradient. In the preferred embodiment of the invention, the gradient is at least l000F./inch. The maximum temperature of the thermal gradient is above the 'y' solvus and below the incipient melting temperature of the alloy in question. The rate at which the workpiece is drawn through the gradient depends on the alloy, but it has been found that speeds of /z2inch/hour results in an elongated grain structure in most superalloys. I have found that slower speeds in this range result in monocrystalline structures, while higher speeds result in elongated parallel columnar grains.
My invention is further illustrated by the following examples:
EXAMPLE 1 A billet was prepared from Rene 120 nickel-base superalloy powders having the composition shown in the table above, except the carbon level was 0.05 percent rather than the typical 0.17 percent. The loose powders, having mesh sizes +200, were encapsulated in a 3%. inch diameter stainless steel capsule having a 0.216 inch wall thickness. The capsule cavity and powder were evacuated to 10* Torr, heated to 500C. under vacuum to remove volatile impurities, cooled to room temperature and sealed.
The entire capsule was then heated to 1175 C. for 2hours and extruded through a die aperture of 0.6 inch X 1.0 inch, approximately an 18/1 reduction. Two and one half inch lengths were cut from the billet. Four tabs, each 0.6 inch X 2- /2 inch X 0.072 inch were cut from the center of each length. Tabs were machined into tensile specimens having a gauge length of 0.150 inch X 0.072 inch X 1.0 inch.
The specimens were then subjected to various combinations of a prior anneal followed by being subjected to a strain at room temperature. Thereafter, the specimens were passed through a gradient fumace having a maximum temperature of 1260C., which is slightly below the alloy incipient melt temperature but above the 'y' solvus temperature. The temperature gradient was about l093C./in.
To determine the speed range, if any, at which elongated grain growth would occur, variable speed anneals were used for some specimens, in which about inch length of gauge section was passed through the hot zone of the furnace at a predetermined speed of about fiiinch/hn; then the drive motor speed was increased to about /2 inch/hr. for another of gauge length, and so on. This determined for a given set of processing conditions whether elongated grain growth would or would not occur.
The results of the critical strain experiments on Rene of speciment gradient annealed at a maximum temperature of 1260C. are shown in Table II:
It may be concluded from the results that the necessary conditions for achieving elongated grain growth in this alloy are a high temperature heat treatment followed by a critical strain at room temperature greater than 1% total elongation, but less than 3 percent.
EXAMPLE II A billet of Rene 120 was prepared using the same procedure as Example I, the only difference being the billet contained a normal carbon level, 0.17 percent. It was observed the combination of a heat treatment for 15 minutes at l240C. followed by a strain at room temperature of about 2 percent total strain resulted upon gradient annealing in an elongated grain structure in all attempts tried.
EXAMPLE III A billet of Rene 120 was prepared using the same procedure as Example I, except that the powder contained 2 vol.pct. Al O which had been milled into the metal powder by a mechanical milling process and the extrusion temperature was 1232C.
It was found the same combination of heat treatment and strain resulted in elongated grain structure after recrystallizing in a temperature gradient. However, it was found that the room temperature strain resulting in op- .timum elongated grain structure was about'2.75-3.0 percent, somewhat higher than the 2 percent strain in the nondispersoid-containing alloy.
EXAMPLE IV A billet of Rene 80 was prepared from superalloy powders having the composition shown in the table above. The loose powders, having 325 mesh size, were encapsulated in a 2-1/16 inch carbon steel extrusion can, evacuated, sealed, and extruded through a k inch round die aperture (approximately an 18/ l reduction) at 1093C. Lengths of the extruded bar were solution annealed for 2 hours at l2l0C. and cooled to room temperature in an air blast. Microstructural examination revealed a fine-grained structure having an average diameter of approximately 20 microns.
Tensile specimens having a A inch gauge diameter and 1 inch gauge length were machined from the /2 inch diameter stock. Four specimens were deformed in tension at room temperature to respective strains of 2.3, 3.0, 3.5, and 10 percent at a strain rate of 0.02 inlinlmin. Following tensile deformation, the four specimens were surface ground to 0.1 inch thick flats to insure uniform heating during the gradient anneal. Specimens were then subjected to a 50 percent l-lCl-SO percent HNO acid solution for & hour to remove potential grain nuclei at the specimen surface.
The specimens were then recrystallized in a temperature gradient at 1205C. at inch/hour. Microstructural examination revealed no evidence of elongated grain structure in the specimen deformed to 2.3 percent tensile strain, indicating the critical strain had not been reached. In the specimen strained to 10 percent strain, grains in the gauge length consisted of equiaxed grains about 20 microns in diameter. At both ends of the gauge length, where the plastic strain decreased from the nominal 10 percent to zero, a continuous region of columnar grains were observed to have been .nucleated, indicating the validity of the strain anneal concept for the subject material. On the specimen strained to 3.5 percent, large (-40 microns) equiaxed grains were observed in the gauge length, again ;with columnar grains at the ends of the gauge length indicating the critical strain in the gauge length had been exceeded. In the specimen strained to 3.0 percent, the grain structure in the gauge length consisted of long columnar grains believed typical of specimens strained to the critical strain.
EXAMPLE V following the procedure of Example 1, four test specimens were prepared from Rene 120 (0. l 7 percent C.)
@extrudedbar having an 0.6 inch X 1.0 inch cross-secspeed). Thereafter the 'strained samples were unidirectionally recrystallized in a gradient furnace at 1254C. at l inch/hr. to produce elongated grains in the specimen gauge section, with fine, equi-axed grains at the gauge length ends, in the specimen shoulders, and in the specimen grip sections. The specimens were then given a heat treatment according to the following schedule: 1236C. for 1 hour; lO93C. for 4 hours; and 900C. for 16 hours. Subsequently, a second gauge length having dimensions 0.50 inch long X 0.100 inch wide X 0.072 inch thick in the elongated grain portion was machined into the first gauge section of two of the four specimens.
Stress-rupture tests were performed in air at the fol- Sample V-A failed in 0.6 hours and sample V-B failed on loading. Both samples failed through the 0.150 inch width in the first gauge section in the fine-grained portion of the specimen. The stress in the 0.150 inch portion of the specimens is two-thirds that in the 0.100. inch portion. This indicates that the stress rupture strength of fine-grained Rene is poor at 1650F. /40 KS1 and l800F./20 KSI.
To evaluate the stress rupture strength of the elongated grain structure portion of the specimen, the two remaining specimens were machined. To ensure that failure occurred in the elongated portion of the specimen, the second reduced gauge sections were machined to dimensions of 0.50 inch long X 0.072 thick X 0.060 inch wide, so that the stress in the second gauge section was forty percent of that in the first gauge section. Test results were:
The above results indicate the degree of mechanical property improvement in Rene 120 occurring from production of an elongated grain structure.
It will be appreciated that the invention is not limited to the specific details shown in the examples and illustrations and that various modifications may be made within the ordinary skill in the art without departing from the spirit and scope of the invention.
I claim:
1. An article of manufacture having a chemically inert phase dispersed therein comprising a unidirectionally recrystallized nickel-base superalloy body free of eutectic module inhomogenieties having an aligned elongated grain or monocrystalline structure and containing up to 10 percent by volume of chemically chamically inert phase dispersed therein wherein said nickel-base superalloy body is prepared by a method comprising the steps of:
a. hot compacting a nickel-base superalloy powder to form .a relatively dense solid consisting essentially of a 'y precipitate phase within a 'y matrix,
b. solution heat treating the dense solid at a temperature below the incipient melting point and high perature to form a body having an aligned elongated grain or monocrystalline structure. 2. The article of claim 1, wherein the strain is equivalent to an elongation at room temperature of about 1-3 percent.
3. The article of claim 1, wherein the strain is equivalent to an elongation at a temperature of 1200-1400F. of 8-10 percent.
4. The article of claim 1, wherein the nickel-base superalloy powder is compacted by canning in a steel jacket under vacuum to form a billet and the billet is hot extruded.
5. The article of claim 1, wherein said chemically inert phase ia alumina.
8 6. The article of claim 1, wherein said inert phase is 7. The article of claim 1, wherein said alloy consists essentially in weight percent of about:
Composition Weight Aluminum 4.25 Chromium 9.25 Titanium 4.0 10 Cobalt 10.0 Molybdenum 2.0 Tungsten 7.0 Tantalum 3.75 Carbon 0. l 7 Boron 0.0 l 5 l5 Zirconium 0.05
8. The article of claim 4, wherein the strained material is unidirectionally recrystallized by drawing the material through a gradient furnace having a temperature gradient of at least 1000F. per inch and the material is drawn at a rate of about 0.5-2 inches per hour.

Claims (8)

1. AN ARTICLE OF MANUFACTURE HAING A CHEMICALLY INERT PHASE DISPERSED THEREIN COMPRISING A UNIDIRECTIONALLY RECRYSTALLIZED NICKEL-BASE SUPERALLOY BODY FREE OF EUTECTIC MODULE INHOMOGENIETIES HAVING AN ALIGNED ELONGATED GRAIN OR MONOCRYSTALLINE STRUCTURE AND CONTAINING UP TO 10 PERCENT BY VOLUME OF CHEMICALLY CHAMICALLY INERT PHASE DISPERSED THEREIN WHERIN SAID NICKEL-BASE SUPERALLOY BODY IS PREPAREED BY A METHOD COMPRISING THE STEPS OF: A. HOT COMPACTING A NICKEL-BASE SUPERALLOY POWDER TO FROM A RELATIVELY DENSE SOLID CONSISTING ESSENTAILLY OF A Y'' PRECIPITATE PHASE WITH A Y'' MATRIX. B. SOLUTION HEAT TREATING THE DENSE SOLID AT A TEMPERATURE BELOW THE INCIPIENT MELTING POINT AND HIGH ENOUGH TO DISSOLVE A SUBSTANTIAL PORTION OF THE Y'' PHASE. C. SUBJECTING THE MATERIAL TO A SUFFICIENT STRAIN BELOW THE RECRYSTALLIZATION TEMPERATURE TO PERMIT SUBSEQUENT RECRYSTALLIZATION, AND D. UNIDIRECTIONALLY RECRYSTALLIZATION THE STRAINED MATERIAL IN A TEMPERATURE GRADIENT BELOW THE INCIPIENT MELTING TEMPER-
2. The article of claim 1, wherein the strain is equivalent to an elongation at room temperature of about 1-3 percent.
3. The article of claim 1, wherein the strain is equivalent to an elongation at a temperature of 1200-1400*F. of 8-10 percent.
4. The article of claim 1, wherein the nickel-base superalloy powder is compacted by canning in a steel jacket under vacuum to form a billet and the billet is hot extruded.
5. The article of claim 1, wherein said chemically inert phase ia alumina.
6. The article of claim 1, wherein said inert phase is yttria.
7. The article of claim 1, wherein said alloy consists essentially in weight percent of about:
8. The article of claim 4, wherein the strained material is unidirectionally recrystallized by drawing the material through a gradient furnace having a temperature gradient of at least 1000*F. per inch and the material is drawn at a rate of about 0.5-2 inches per hour.
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NLAANVRAGE7102685,A NL171309C (en) 1970-03-02 1971-03-01 METHOD FOR THE MANUFACTURE OF A SEMICONDUCTOR BODY FORMING A SILICONE DIOXIDE LAYER ON A SURFACE OF A SILICONE MONOCRYSTALLINE BODY
FR7107147A FR2084089A5 (en) 1970-03-02 1971-03-02
US00120289A US3821783A (en) 1970-03-02 1971-03-02 Semiconductor device with a silicon monocrystalline body having a specific crystal plane
DE2109874A DE2109874C3 (en) 1970-03-02 1971-03-02 Semiconductor component with a monocrystalline silicon body and method for manufacturing
GB2288671A GB1318832A (en) 1970-03-02 1971-04-19 Semiconductor devices
US00402306A US3850702A (en) 1970-03-02 1973-10-01 Method of making superalloy bodies
US473407A US3920489A (en) 1970-03-02 1974-05-28 Method of making superalloy bodies
US483837A US3920492A (en) 1970-03-02 1974-06-27 Process for manufacturing a semiconductor device with a silicon monocrystalline body having a specific crystal plane

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US00402306A US3850702A (en) 1970-03-02 1973-10-01 Method of making superalloy bodies
US473407A US3920489A (en) 1970-03-02 1974-05-28 Method of making superalloy bodies
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