US20100254850A1 - Ceracon forging of l12 aluminum alloys - Google Patents

Ceracon forging of l12 aluminum alloys Download PDF

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US20100254850A1
US20100254850A1 US12419677 US41967709A US2010254850A1 US 20100254850 A1 US20100254850 A1 US 20100254850A1 US 12419677 US12419677 US 12419677 US 41967709 A US41967709 A US 41967709A US 2010254850 A1 US2010254850 A1 US 2010254850A1
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weight percent
powder
aluminum
aluminum alloy
l1
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Awadh B. Pandey
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United Technologies Corp
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    • CCHEMISTRY; METALLURGY
    • C22METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22CALLOYS
    • C22C1/00Making alloys
    • C22C1/04Making alloys by powder metallurgy
    • C22C1/0408Light metal alloys
    • C22C1/0416Aluminium-based alloys
    • BPERFORMING OPERATIONS; TRANSPORTING
    • B22CASTING; POWDER METALLURGY
    • B22FWORKING METALLIC POWDER; MANUFACTURE OF ARTICLES FROM METALLIC POWDER; MAKING METALLIC POWDER
    • B22F3/00Manufacture of workpieces or articles from metallic powder characterised by the manner of compacting or sintering; Apparatus specially adapted therefor ; Presses and furnaces
    • B22F3/17Manufacture of workpieces or articles from metallic powder characterised by the manner of compacting or sintering; Apparatus specially adapted therefor ; Presses and furnaces by forging
    • BPERFORMING OPERATIONS; TRANSPORTING
    • B22CASTING; POWDER METALLURGY
    • B22FWORKING METALLIC POWDER; MANUFACTURE OF ARTICLES FROM METALLIC POWDER; MAKING METALLIC POWDER
    • B22F2998/00Supplementary information concerning processes or compositions relating to powder metallurgy
    • B22F2998/10Processes characterised by the sequence of their steps
    • BPERFORMING OPERATIONS; TRANSPORTING
    • B22CASTING; POWDER METALLURGY
    • B22FWORKING METALLIC POWDER; MANUFACTURE OF ARTICLES FROM METALLIC POWDER; MAKING METALLIC POWDER
    • B22F2999/00Aspects linked to processes or compositions used in powder metallurgy

Abstract

A method for producing high strength aluminum alloy consolidated billets containing L12 dispersoids by Ceracon forging is disclosed. The method comprises forming an aluminum alloy powder compact preform containing L12 dispersoid forming elements therein and encompassing the preform in a flowable pressure transmitting medium in a die in a hydraulic press. The die, pressure transmitting medium and preform are then heated and the preform is forged by applying pressure to the pressure transmitting medium by the ram of the hydraulic press. The unequal axial and radial strain resulting from this type of forging results in improved mechanical properties of L12 aluminum alloys.

Description

    CROSS-REFERENCE TO RELATED APPLICATION(S)
  • This application is related to the following co-pending applications that were filed on Dec. 9, 2008 herewith and are assigned to the same assignee: CONVERSION PROCESS FOR HEAT TREATABLE L12 ALUMINUM ALLOYS, Ser. No. 12/316,020; A METHOD FOR FORMING HIGH STRENGTH ALUMINUM ALLOYS CONTAINING L12 INTERMETALLIC DISPERSOIDS, Ser. No. 12/316,046; and A METHOD FOR PRODUCING HIGH STRENGTH ALUMINUM ALLOY POWDER CONTAINING L12 INTERMETALLIC DISPERSOIDS, Ser. No. 12/316,047.
  • This application is also related to the following co-pending applications that were filed on Apr. 18, 2008, and are assigned to the same assignee: L12 ALUMINUM ALLOYS WITH BIMODAL AND TRIMODAL DISTRIBUTION, Ser. No. 12/148,395; DISPERSION STRENGTHENED L12 ALUMINUM ALLOYS, Ser. No. 12/148,432; HEAT TREATABLE L12 ALUMINUM ALLOYS, Ser. No. 12/148,383; HIGH STRENGTH L12 ALUMINUM ALLOYS, Ser. No. 12/148,394; HIGH STRENGTH L12 ALUMINUM ALLOYS, Ser. No. 12/148,382; HEAT TREATABLE L12 ALUMINUM ALLOYS, Ser. No. 12/148,396; HIGH STRENGTH L12 ALUMINUM ALLOYS, Ser. No. 12/148,387; HIGH STRENGTH ALUMINUM ALLOYS WITH L12 PRECIPITATES, Ser. No. 12/148,426; HIGH STRENGTH L12 ALUMINUM ALLOYS, Ser. No. 12/148,459; and L12 STRENGTHENED AMORPHOUS ALUMINUM ALLOYS, Ser. No. 12/148,458.
  • BACKGROUND
  • The present invention relates generally to aluminum alloys and more specifically to a method for forming high strength aluminum alloy billets having L12 dispersoids therein.
  • The combination of high strength, ductility, and fracture toughness, as well as low density, make aluminum alloys natural candidates for aerospace and space applications. However, their use is typically limited to temperatures below about 300° F. (149° C.) since most aluminum alloys start to lose strength in that temperature range as a result of coarsening of strengthening precipitates.
  • The development of aluminum alloys with improved elevated temperature mechanical properties is a continuing process. Some attempts have included aluminum-iron and aluminum-chromium based alloys such as Al—Fe—Ce, Al—Fe—V—Si, Al—Fe—Ce—W, and Al—Cr—Zr—Mn that contain incoherent dispersoids. These alloys, however, also lose strength at elevated temperatures due to particle coarsening. In addition, these alloys exhibit ductility and fracture toughness values lower than other commercially available aluminum alloys.
  • Other attempts have included the development of mechanically alloyed Al—Mg and Al—Ti alloys containing ceramic dispersoids. These alloys exhibit improved high temperature strength due to the particle dispersion, but the ductility and fracture toughness are not improved.
  • U.S. Pat. No. 6,248,453 owned by the assignee of the present invention discloses aluminum alloys strengthened by dispersed Al3X L12 intermetallic phases where X is selected from the group consisting of Sc, Er, Lu, Yb, Tm, and Lu. The Al3X particles are coherent with the aluminum alloy matrix and are resistant to coarsening at elevated temperatures. The improved mechanical properties of the disclosed dispersion strengthened L12 aluminum alloys are stable up to 572° F. (300° C.). U.S. Patent Application Publication No. 2006/0269437 Al also owned commonly discloses a high strength aluminum alloy that contains scandium and other elements that is strengthened by L12 dispersoids.
  • L12 strengthened aluminum alloys have high strength and improved fatigue properties compared to commercially available aluminum alloys. Fine grain size results in improved mechanical properties of materials. Hall-Petch strengthening has been known for decades where strength increases as grain size decreases. An optimum grain size for optimum strength is in the nano range of about 30 to 100 nm. These alloys also have lower ductility.
  • SUMMARY
  • The present invention is a method for forming aluminum alloy billets with high strength and acceptable fracture toughness. In embodiments, the alloys have coherent L12 Al3X dispersoids where X is at least one first element selected from scandium, erbium, thulium, ytterbium, and lutetium, and at least one second element selected from gadolinium, yttrium, zirconium, titanium, hafnium, and niobium. The balance is substantially aluminum containing at least one alloying element selected from silicon, magnesium, lithium, copper, zinc, and nickel.
  • The alloys are formed by encompassing a powder preform of an aluminum alloy body containing L12 dispersoid forming elements in a heated, flowable pressure transmitting medium, and rapidly compressing the powder to consolidate the perform to form the billet. Use of graphite or a ceramic as the pressure transfer medium causes a non-isostatic pressure field to form in the chamber. During consolidation, the powder preform undergoes an axial compression that exceeds radial expansion. The resulting shear strain cleans the surface oxide from the particles and increases metal-to-metal contact improving forging density.
  • This method is known in the industry as Ceracon forging. It has not been attempted using aluminum alloy containing L12 dispersoids. Examples of the Ceracon forging process are shown in U.S. Pat. No. 4,667,497, U.S. Patent Application Publication No. 2005/0147520 and U.S. Pat. No. 7,097,807 and are included in their entirety by reference.
  • The compression takes place in a closed container and results in billets of aluminum alloy containing L12 dispersoids with a density of essentially 100%.
  • BRIEF DESCRIPTION OF THE DRAWINGS
  • FIG. 1 is an aluminum scandium phase diagram.
  • FIG. 2 is an aluminum erbium phase diagram.
  • FIG. 3 is an aluminum thulium phase diagram.
  • FIG. 4 is an aluminum ytterbium phase diagram.
  • FIG. 5 is an aluminum lutetium phase diagram.
  • FIG. 6 is a schematic diagram of a vertical gas atomizer.
  • FIG. 7 is a scanning electron micrograph of the gas atomized inventive L12 aluminum alloy powder.
  • FIG. 8 is a diagram showing the processing steps to consolidate L12 aluminum alloy powder.
  • FIG. 9 is a schematic diagram illustrating non-isostatic forging.
  • FIG. 10 is a diagram showing the processing steps to Ceracon forge an L12 aluminum alloy powder preform.
  • DETAILED DESCRIPTION 1. L12 Aluminum Alloys
  • The alloy products of this invention are formed from aluminum based alloys with high strength and fracture toughness for applications at temperatures from about −420° F. (−251 ° C.) up to about 650° F. (343° C.). The aluminum alloy comprises a solid solution of aluminum and at least one element selected from silicon, magnesium, lithium, copper, zinc, and nickel strengthened by L12 coherent precipitates where X is at least one first element selected from scandium, erbium, thulium, ytterbium, and lutetium, and at least one second element selected from gadolinium, yttrium, zirconium, titanium, hafnium, and niobium.
  • The aluminum silicon system is a simple eutectic alloy system with a eutectic reaction at 12.5 weight percent silicon and 1077° F. (577° C.). There is little solubility of silicon in aluminum at temperatures up to 930° F. (500° C.) and none of aluminum in silicon. However, the solubility can be extended significantly by utilizing rapid solidification techniques
  • The binary aluminum magnesium system is a simple eutectic at 36 weight percent magnesium and 842° F. (450° C.). There is complete solubility of magnesium and aluminum in the rapidly solidified inventive alloys discussed herein
  • The binary aluminum lithium system is a simple eutectic at 8 weight percent lithium and 1105° (596° C.). The equilibrium solubility of 4 weight percent lithium can be extended significantly by rapid solidification techniques. There can be complete solubility of lithium in the rapid solidified inventive alloys discussed herein.
  • The binary aluminum copper system is a simple eutectic at 32 weight percent copper and 1018° F. (548° C.). There can be complete solubility of copper in the rapidly solidified inventive alloys discussed herein.
  • The aluminum zinc binary system is a eutectic alloy system involving a monotectoid reaction and a miscibility gap in the solid state. There is a eutectic reaction at 94 weight percent zinc and 718° F. (381° C.). Zinc has maximum solid solubility of 83.1 weight percent in aluminum at 717.8° F. (381° C.) which can be extended by rapid solidification processes. Decomposition of the super saturated solid solution of zinc in aluminum gives rise to spherical and ellipsoidal GP zones which are coherent with the matrix and act to strengthen the alloy.
  • The aluminum nickel binary system is a simple eutectic at 5.7 weight percent nickel and 1183.8° F. (639.9° C.). There is little solubility of nickel in aluminum. However, the solubility can be extended significantly by utilizing rapid solidification processes. The equilibrium phase in the aluminum nickel eutectic system is L12 intermetallic Al3Ni.
  • In the aluminum based alloys disclosed herein, scandium, erbium, thulium, ytterbium, and lutetium are potent strengtheners that have low diffusivity and low solubility in aluminum. All these elements form equilibrium Al3X intermetallic dispersoids where X is at least one of scandium, erbium, thulium, ytterbium, and lutetium, that have an L12 structure that is an ordered face centered cubic structure with the X atoms located at the corners and aluminum atoms located on the cube faces of the unit cell.
  • Scandium forms Al3Sc dispersoids that are fine and coherent with the aluminum matrix. Lattice parameters of aluminum and Al3Sc are very close (0.405 nm and 0.410 nm respectively), indicating that there is minimal or no driving force for causing growth of the Al3Sc dispersoids. This low interfacial energy makes the Al3Sc dispersoids thermally stable and resistant to coarsening up to temperatures as high as about 842° F. (450° C.). Additions of magnesium in aluminum increase the lattice parameter of the aluminum matrix, and decrease the lattice parameter mismatch further increasing the resistance of the Al3Sc to coarsening. Additions of zinc, copper, lithium, silicon, and nickel provide solid solution and precipitation strengthening in the aluminum alloys. In the alloys of this invention these Al3Sc dispersoids are made stronger and more resistant to coarsening at elevated temperatures by adding suitable alloying elements such as gadolinium, yttrium, zirconium, titanium, hafnium, niobium, or combinations thereof, that enter Al3Sc in solution.
  • Erbium forms Al3Er dispersoids in the aluminum matrix that are fine and coherent with the aluminum matrix. The lattice parameters of aluminum and Al3Er are close (0.405 nm and 0.417 nm respectively), indicating there is minimal driving force for causing growth of the Al3Er dispersoids. This low interfacial energy makes the Al3Er dispersoids thermally stable and resistant to coarsening up to temperatures as high as about 842° F. (450° C.). Additions of magnesium in aluminum increase the lattice parameter of the aluminum matrix, and decrease the lattice parameter mismatch further increasing the resistance of the Al3Er to coarsening. Additions of zinc, copper, lithium, silicon, and nickel provide solid solution and precipitation strengthening in the aluminum alloys. In the alloys of this invention, these Al3Er dispersoids are made stronger and more resistant to coarsening at elevated temperatures by adding suitable alloying elements such as gadolinium, yttrium, zirconium, titanium, hafnium, niobium, or combinations thereof that enter Al3Er in solution.
  • Thulium forms metastable Al3Tm dispersoids in the aluminum matrix that are fine and coherent with the aluminum matrix. The lattice parameters of aluminum and Al3Tm are close (0.405 nm and 0.420 nm respectively), indicating there is minimal driving force for causing growth of the Al3Tm dispersoids. This low interfacial energy makes the Al3Tm dispersoids thermally stable and resistant to coarsening up to temperatures as high as about 842° F. (450° C.). Additions of magnesium in aluminum increase the lattice parameter of the aluminum matrix, and decrease the lattice parameter mismatch further increasing the resistance of the Al3Tm to coarsening. Additions of zinc, copper, lithium, silicon, and nickel provide solid solution and precipitation strengthening in the aluminum alloys. In the alloys of this invention these Al3Tm dispersoids are made stronger and more resistant to coarsening at elevated temperatures by adding suitable alloying elements such as gadolinium, yttrium, zirconium, titanium, hafnium, niobium, or combinations thereof that enter Al3Tm in solution.
  • Ytterbium forms Al3Yb dispersoids in the aluminum matrix that are fine and coherent with the aluminum matrix. The lattice parameters of Al and Al3Yb are close (0.405 nm and 0.420 nm respectively), indicating there is minimal driving force for causing growth of the Al3Yb dispersoids. This low interfacial energy makes the Al3Yb dispersoids thermally stable and resistant to coarsening up to temperatures as high as about 842° F. (450° C.). Additions of magnesium in aluminum increase the lattice parameter of the aluminum matrix, and decrease the lattice parameter mismatch further increasing the resistance of the Al3Yb to coarsening. Additions of zinc, copper, lithium, silicon, and nickel provide solid solution and precipitation strengthening in the aluminum alloys. In the alloys of this invention, these Al3Yb dispersoids are made stronger and more resistant to coarsening at elevated temperatures by adding suitable alloying elements such as gadolinium, yttrium, zirconium, titanium, hafnium, niobium, or combinations thereof that enter Al3Yb in solution.
  • Lutetium forms Al3Lu dispersoids in the aluminum matrix that are fine and coherent with the aluminum matrix. The lattice parameters of Al and Al3Lu are close (0.405 nm and 0.419 nm respectively), indicating there is minimal driving force for causing growth of the Al3Lu dispersoids. This low interfacial energy makes the Al3Lu dispersoids thermally stable and resistant to coarsening up to temperatures as high as about 842° F. (450° C.). Additions of magnesium in aluminum increase the lattice parameter of the aluminum matrix, and decrease the lattice parameter mismatch further increasing the resistance of the Al3Lu to coarsening. Additions of zinc, copper, lithium, silicon, and nickel provide solid solution and precipitation strengthening in the aluminum alloys. In the alloys of this invention, these Al3Lu dispersoids are made stronger and more resistant to coarsening at elevated temperatures by adding suitable alloying elements such as gadolinium, yttrium, zirconium, titanium, hafnium, niobium, or mixtures thereof that enter Al3Lu in solution.
  • Gadolinium forms metastable Al3Gd dispersoids in the aluminum matrix that are stable up to temperatures as high as about 842° F. (450° C.) due to their low diffusivity in aluminum. The Al3Gd dispersoids have a D019 structure in the equilibrium condition. Despite its large atomic size, gadolinium has fairly high solubility in the Al3X intermetallic dispersoids (where X is scandium, erbium, thulium, ytterbium or lutetium). Gadolinium can substitute for the X atoms in Al3X intermetallic, thereby forming an ordered L12 phase which results in improved thermal and structural stability.
  • Yttrium forms metastable Al3Y dispersoids in the aluminum matrix that have an L12 structure in the metastable condition and a D019 structure in the equilibrium condition. The metastable Al3Y dispersoids have a low diffusion coefficient which makes them thermally stable and highly resistant to coarsening. Yttrium has a high solubility in the Al3X intermetallic dispersoids allowing large amounts of yttrium to substitute for X in the Al3X L12 dispersoids which results in improved thermal and structural stability.
  • Zirconium forms Al3Zr dispersoids in the aluminum matrix that have an L12 structure in the metastable condition and D023 structure in the equilibrium condition. The metastable Al3Zr dispersoids have a low diffusion coefficient which makes them thermally stable and highly resistant to coarsening. Zirconium has a high solubility in the Al3X dispersoids allowing large amounts of zirconium to substitute for X in the Al3X dispersoids, which results in improved thermal and structural stability.
  • Titanium forms Al3Ti dispersoids in the aluminum matrix that have an L12 structure in the metastable condition and DO22 structure in the equilibrium condition. The metastable Al3Ti despersoids have a low diffusion coefficient which makes them thermally stable and highly resistant to coarsening. Titanium has a high solubility in the Al3X dispersoids allowing large amounts of titanium to substitute for X in the Al3X dispersoids, which result in improved thermal and structural stability.
  • Hafnium forms metastable Al3Hf dispersoids in the aluminum matrix that have an L12 structure in the metastable condition and a D023 structure in the equilibrium condition. The Al3Hf dispersoids have a low diffusion coefficient, which makes them thermally stable and highly resistant to coarsening. Hafnium has a high solubility in the Al3X dispersoids allowing large amounts of hafnium to substitute for scandium, erbium, thulium, ytterbium, and lutetium in the above mentioned Al3X dispersoids, which results in stronger and more thermally stable dispersoids.
  • Niobium forms metastable Al3Nb dispersoids in the aluminum matrix that have an L12 structure in the metastable condition and a D022 structure in the equilibrium condition. Niobium has a lower solubility in the Al3X dispersoids than hafnium or yttrium, allowing relatively lower amounts of niobium than hafnium or yttrium to substitute for X in the Al3X dispersoids. Nonetheless, niobium can be very effective in slowing down the coarsening kinetics of the Al3X dispersoids because the Al3Nb dispersoids are thermally stable. The substitution of niobium for X in the above mentioned Al3X dispersoids results in stronger and more thermally stable dispersoids.
  • Al3X L12 precipitates improve elevated temperature mechanical properties in aluminum alloys for two reasons. First, the precipitates are ordered intermetallic compounds. As a result, when the particles are sheared by glide dislocations during deformation, the dislocations separate into two partial dislocations separated by an anti-phase boundary on the glide plane. The energy to create the anti-phase boundary is the origin of the strengthening. Second, the cubic L12 crystal structure and lattice parameter of the precipitates are closely matched to the aluminum solid solution matrix. This results in a lattice coherency at the precipitate/matrix boundary that resists coarsening. The lack of an interphase boundary results in a low driving force for particle growth and resulting elevated temperature stability. Alloying elements in solid solution in the dispersed strengthening particles and in the aluminum matrix that tend to decrease the lattice mismatch between the matrix and particles will tend to increase the strengthening and elevated temperature stability of the alloy.
  • L12 phase strengthened aluminum alloys are important structural materials because of their excellent mechanical properties and the stability of these properties at elevated temperature due to the resistance of the coherent dispersoids in the microstructure to particle coarsening. The mechanical properties are optimized by maintaining a high volume fraction of L12 dispersoids in the microstructure. The L12 dispersoid concentration following aging scales as the amount of L12 phase forming elements in solid solution in the aluminum alloy following quenching. Examples of L12 phase forming elements include but are not limited to Sc, Er, Th, Yb, and Lu. The concentration of alloying elements in solid solution in alloys cooled from the melt is directly proportional to the cooling rate. Exemplary aluminum alloys for the bimodal system alloys of this invention include, but are not limited to (in weight percent unless otherwise specified):
  • about Al—M-(0.1-4)Sc-(0.1-20)Gd;
  • about Al—M-(0.1-20)Er-(0.1-20)Gd;
  • about Al—M-(0.1-15)Tm-(0.1-20)Gd;
  • about Al—M-(0.1-25)Yb-(0.1-20)Gd;
  • about Al—M-(0.1-25)Lu-(0.1-20)Gd;
  • about Al—M-(0.1-4)Sc-(0.1-20)Y;
  • about Al—M-(0.1-20)Er-(0.1-20)Y;
  • about Al—M-(0.1-15)Tm-(0.1-20)Y;
  • about Al—M-(0.1-25)Yb-(0.1-20)Y;
  • about Al—M-(0.1-25)Lu-(0.1-20)Y;
  • about Al—M-(0.1-4)Sc-(0.05-4)Zr;
  • about Al—M-(0.1-20)Er-(0.05-4)Zr;
  • about Al—M-(0.1-15)Tm-(0.05-4)Zr;
  • about Al—M-(0.1-25)Yb-(0.05-4)Zr;
  • about Al—M-(0.1-25)Lu-(0.05-4)Zr;
  • about Al—M-(0.1-4)Sc-(0.05-10)Ti;
  • about Al—M-(0.1-20)Er-(0.05-10)Ti;
  • about Al—M-(0.1-15)Tm-(0.05-10)Ti;
  • about Al—M-(0.1-25)Yb-(0.05-10)Ti;
  • about Al—M-(0.1-25)Lu-(0.05-10)Ti;
  • about Al—M-(0.1-4)Sc-(0.05-10)Hf;
  • about Al—M-(0.1-20)Er-(0.05-10)Hf;
  • about Al—M-(0.1-15)Tm-(0.05-10)Hf;
  • about Al—M-(0.1-25)Yb-(0.05-10)Hf;
  • about Al—M-(0.1-25)Lu-(0.05-10)Hf;
  • about Al—M-(0.1-4)Sc-(0.05-5)Nb;
  • about Al—M-(0.1-20)Er-(0.05-5)Nb;
  • about Al—M-(0.1-15)Tm-(0.05-5)Nb;
  • about Al—M-(0.1-25)Yb-(0.05-5)Nb; and
  • about Al—M-(0.1-25)Lu-(0.05-5)Nb.
  • M is at least one of about (4-25) weight percent silicon, (1-8) weight percent magnesium, (0.5-3) weight percent lithium, (0.2-6.5) weight percent copper, (3-12) weight percent zinc, and (1-12) weight percent nickel.
  • The amount of silicon present in the fine grain matrix of this invention if any may vary from about 4 to about 25 weight percent, more preferably from about 4 to about 18 weight percent, and even more preferably from about 5 to about 11 weight percent.
  • The amount of magnesium present in the fine grain matrix of this invention if any may vary from about 1 to about 8 weight percent, more preferably from about 3 to about 7.5 weight percent, and even more preferably from about 4 to about 6.5 weight percent.
  • The amount of lithium present in the fine grain matrix of this invention if any may vary from about 0.5 to about 3 weight percent, more preferably from about 1 to about 2.5 weight percent, and even more preferably from about 1 to about 2 weight percent.
  • The amount of copper present in the fine grain matrix of this invention if any may vary from about 0.2 to about 6.5 weight percent, more preferably from about 0.5 to about 5.0 weight percent, and even more preferably from about 2 to about 4.5 weight percent.
  • The amount of zinc present in the fine grain matrix of this invention if any may vary from about 3 to about 12 weight percent, more preferably from about 4 to about 10 weight percent, and even more preferably from about 5 to about 9 weight percent.
  • The amount of nickel present in the fine grain matrix of this invention if any may vary from about 1 to about 12 weight percent, more preferably from about 2 to about 10 weight percent, and even more preferably from about 4 to about 10 weight percent.
  • The amount of scandium present in the fine grain matrix of this invention if any may vary from 0.1 to about 4 weight percent, more preferably from about 0.1 to about 3 weight percent, and even more preferably from about 0.2 to about 2.5 weight percent. The Al—Sc phase diagram shown in FIG. 1 indicates a eutectic reaction at about 0.5 weight percent scandium at about 1219° F. (659° C.) resulting in a solid solution of scandium and aluminum and Al3Sc dispersoids. Aluminum alloys with less than 0.5 weight percent scandium can be quenched from the melt to retain scandium in solid solution that may precipitate as dispersed L12 intermetallic Al3Sc following an aging treatment. Alloys with scandium in excess of the eutectic composition (hypereutectic alloys) can only retain scandium in solid solution by rapid solidification processing (RSP) where cooling rates are in excess of about 103° C./second.
  • The amount of erbium present in the fine grain matrix of this invention, if any, may vary from about 0.1 to about 20 weight percent, more preferably from about 0.3 to about 15 weight percent, and even more preferably from about 0.5 to about 10 weight percent. The Al—Er phase diagram shown in FIG. 2 indicates a eutectic reaction at about 6 weight percent erbium at about 1211° F. (655° C.). Aluminum alloys with less than about 6 weight percent erbium can be quenched from the melt to retain erbium in solid solutions that may precipitate as dispersed L12 intermetallic Al3Er following an aging treatment. Alloys with erbium in excess of the eutectic composition can only retain erbium in solid solution by rapid solidification processing (RSP) where cooling rates are in excess of about 103° C./second.
  • The amount of thulium present in the alloys of this invention, if any, may vary from about 0.1 to about 15 weight percent, more preferably from about 0.2 to about 10 weight percent, and even more preferably from about 0.4 to about 6 weight percent. The Al—Tm phase diagram shown in FIG. 3 indicates a eutectic reaction at about 10 weight percent thulium at about 1193° F. (645° C.). Thulium forms metastable Al3Tm dispersoids in the aluminum matrix that have an L12 structure in the equilibrium condition. The Al3Tm dispersoids have a low diffusion coefficient which makes them thermally stable and highly resistant to coarsening. Aluminum alloys with less than 10 weight percent thulium can be quenched from the melt to retain thulium in solid solution that may precipitate as dispersed metastable L12 intermetallic Al3Tm following an aging treatment. Alloys with thulium in excess of the eutectic composition can only retain Tm in solid solution by rapid solidification processing (RSP) where cooling rates are in excess of about 103° C./second.
  • The amount of ytterbium present in the alloys of this invention, if any, may vary from about 0.1 to about 25 weight percent, more preferably from about 0.3 to about 20 weight percent, and even more preferably from about 0.4 to about 10 weight percent. The Al—Yb phase diagram shown in FIG. 4 indicates a eutectic reaction at about 21 weight percent ytterbium at about 1157° F. (625° C.). Aluminum alloys with less than about 21 weight percent ytterbium can be quenched from the melt to retain ytterbium in solid solution that may precipitate as dispersed L12 intermetallic Al3Yb following an aging treatment. Alloys with ytterbium in excess of the eutectic composition can only retain ytterbium in solid solution by rapid solidification processing (RSP) where cooling rates are in excess of about 103° C./second.
  • The amount of lutetium present in the alloys of this invention, if any, may vary from about 0.1 to about 25 weight percent, more preferably from about 0.3 to about 20 weight percent, and even more preferably from about 0.4 to about 10 weight percent. The Al—Lu phase diagram shown in FIG. 5 indicates a eutectic reaction at about 11.7 weight percent Lu at about 1202° F. (650° C.). Aluminum alloys with less than about 11.7 weight percent lutetium can be quenched from the melt to retain Lu in solid solution that may precipitate as dispersed L12 intermetallic Al3Lu following an aging treatment. Alloys with Lu in excess of the eutectic composition can only retain Lu in solid solution by rapid solidification processing (RSP) where cooling rates are in excess of about 103° C./second.
  • The amount of gadolinium present in the alloys of this invention, if any, may vary from about 0.1 to about 20 weight percent, more preferably from about 0.3 to about 15 weight percent, and even more preferably from about 0.5 to about 10 weight percent.
  • The amount of yttrium present in the alloys of this invention, if any, may vary from about 0.1 to about 20 weight percent, more preferably from about 0.3 to about 15 weight percent, and even more preferably from about 0.5 to about 10 weight percent.
  • The amount of zirconium present in the alloys of this invention, if any, may vary from about 0.05 to about 4 weight percent, more preferably from about 0.1 to about 3 weight percent, and even more preferably from about 0.3 to about 2 weight percent.
  • The amount of titanium present in the alloys of this invention, if any, may vary from about 0.05 to about 10 weight percent, more preferably from about 0.2 to about 8 weight percent, and even more preferably from about 0.4 to about 4 weight percent.
  • The amount of hafnium present in the alloys of this invention, if any, may vary from about 0.05 to about 10 weight percent, more preferably from about 0.2 to about 8 weight percent, and even more preferably from about 0.4 to about 5 weight percent.
  • The amount of niobium present in the alloys of this invention, if any, may vary from about 0.05 to about 5 weight percent, more preferably from about 0.1 to about 3 weight percent, and even more preferably from about 0.2 to about 2 weight percent.
  • In order to have the best properties for the fine grain matrix of this invention, it is desirable to limit the amount of other elements. Specific elements that should be reduced or eliminated include no more than about 0.1 weight percent iron, 0.1 weight percent chromium, 0.1 weight percent manganese, 0.1 weight percent vanadium, and 0.1 weight percent cobalt. The total quantity of additional elements should not exceed about 1% by weight, including the above listed impurities and other elements.
  • 2. Ceracon Forging of L12 Aluminum Alloys
  • It is advantageous to form L12 strengthened aluminum alloy product from powder. The major reason is that the rapid cooling rate experienced during powder formation from the melt results in high supersaturation of intermetallic L12 phase forming elements in the powder. The high supersaturation leads to a maximum amount of the strengthening phase dispersed throughout the structure in the final consolidated part.
  • The highest cooling rates observed in commercially viable processes are achieved by gas atomization of molten metals to produce powder. Gas atomization is a two fluid process wherein a stream of molten metal is disintegrated by a high velocity gas stream. The end result is that the particles of molten metal eventually become spherical due to surface tension and finely solidify in powder form. Heat from the liquid droplets is transferred to the atomization gas by convection. The solidification rates, depending on the gas and the surrounding environment, can be very high and can exceed 106° C./second. Cooling rates greater than 103° C./second are typically specified to ensure supersaturation of alloying elements in gas atomized L12 aluminum alloy powder in the inventive process described herein.
  • A schematic of typical vertical gas atomizer 100 is shown in FIG. 6A. FIG. 6A is taken from R. Germain, Powder Metallurgy Science Second Edition MPIF (1994) (chapter 3, p. 101) and is incorporated herein by reference. Vacuum or inert gas induction melter 102 is positioned at the top of free flight chamber 104. Vacuum induction melter 102 contains melt 106 which flows by gravity or gas overpressure through nozzle 108. A close up view of nozzle 108 is shown in FIG. 6B. Melt 106 enters nozzle 108 and flows downward till it meets high pressure gas stream from gas source 110 where it is transformed into a spray of droplets. The droplets eventually become spherical due to surface tension and rapidly solidify into spherical powder 112 which collects in collection chamber 114. The gas recirculates through cyclone collector 116 which collects fine powder 118 before returning to the input gas stream. As can be seen from FIG. 6A, the surroundings to which the melt and eventual powder are exposed are completely controlled.
  • There are many effective nozzle designs known in the art to produce spherical metal powder. Nozzle designs with short gas-to-melt separation distances produce finer powders. Confined nozzle designs where gas meets the molten stream at a short distance just after it leaves the atomization nozzle are preferred for the production of the inventive L12 aluminum alloy powders disclosed herein. Higher superheat temperatures cause lower melt viscosity and longer cooling times. Both result in smaller spherical particles.
  • A large number of processing parameters are associated with gas atomization that affect the final product. Examples include melt superheat, gas pressure, metal flow rate, gas type, and gas purity. In gas atomization, the particle size is related to the energy input to the metal. Higher gas pressures, higher superheat temperatures and lower metal flow rates result in smaller particle sizes. Higher gas pressures provide higher gas velocities for a given atomization nozzle design.
  • To maintain purity, inert gases are used, such as helium, argon, and nitrogen. Helium is preferred for rapid solidification because the high heat transfer coefficient of the gas leads to high quenching rates and high supersaturation of alloying elements.
  • Lower metal flow rates and higher gas flow rates favor production of finer powders. The particle size of gas atomized melts typically has a log normal distribution. An example of spherical L12 aluminum alloy powder is shown in the scanning electron micrograph (SEM) of FIG. 7.
  • Oxygen and hydrogen in the powder can degrade the mechanical properties of the final part. It is preferred to limit the oxygen in the L12 alloy powder to about 1 ppm to 2000 ppm. Oxygen is intentionally introduced as a component of the helium gas during atomization. A thin oxide coating on the L12 aluminum powder is beneficial for two reasons. First, the coating prevents agglomeration by contact sintering and secondly, the coating inhibits the chance of explosion of the powder. A controlled amount of oxygen is important in order to provide good ductility and fracture toughness in the final consolidated material. Hydrogen content in the powder is controlled by ensuring the dew point of the helium gas is low. A dew point of about minus 50° F. (minus 45.5° C.) to minus 100° F. (minus 73.3° C.) is preferred.
  • In preparation for final processing, the powder is classified according to size by sieving. To prepare the powder for sieving, if the powder has zero percent oxygen content, the powder may be exposed to nitrogen gas which passivates the powder surface and prevents agglomeration. Finer powder sizes result in improved mechanical properties of the end product. While minus 325 mesh (about 45 microns) powder can be used, minus 450 mesh (about 30 microns) powder is a preferred size in order to provide good mechanical properties in the end product. During the atomization process, powder is collected in collection chambers in order to prevent oxidation of the powder. Collection chambers are used at the bottom of atomization chamber 104 as well as at the bottom of cyclone collector 116. The powder is transported and stored in the collection chambers also. Collection chambers are maintained under positive pressure with nitrogen gas which prevents oxidation of the powder.
  • The process of consolidating the inventive alloy powders into useful forms is schematically illustrated in FIG. 8. L12 aluminum alloy powders 210 are first classified according to size by sieving (step 220). Fine particle sizes are required for optimum mechanical properties in the final part. Next, the classified powders are blended (step 230) in order to maintain microstructural homogeneity in the final part. Blending is necessary because different atomization batches produce powders with varying particle size distributions. Other benefits of blending will be discussed later. Powders may be optionally cryomilled (step 240) to minimize grain size and improve strength. Cryomilling is carried out in a high-energy ball mill under liquid nitrogen, and offers several benefits that will be discussed later.
  • The sieved, blended and (optionally) cryomilled powders are then put in a can (step 250) and vacuum degassed (step 260). Following vacuum degassing, the can is sealed (step 270) under vacuum and forged (step 280) to produce a densified preform. Finally, the preform is Ceracon forged (step 290) to produce a product with improved mechanical properties useful for subsequent service as a high temperature L12 strengthened aluminum alloy. Non-isostatic Ceracon forging will be described later.
  • Sieving (step 220) is a critical step in consolidation because the final mechanical properties relate directly to the particle size. Finer particle size results in finer L12 particle dispersion and finer grain size. Optimum mechanical properties have been observed with −450 mesh (30 micron) powder. Sieving (step 220) also limits the defect size in the powder. Before sieving, the powder is passivated with nitrogen gas in order to prevent agglomeration. Ultrasonic sieving is preferred for its efficiency.
  • Blending (step 230) is another critical step in the consolidation process because it results in improved uniformity of particle size distribution. Gas atomized L12 aluminum alloy powder generally exhibits a bimodal particle size distribution and cross blending of separate powder batches tends to homogenize the particle size distribution. Blending (step 230) is also necessary when separate metal and/or ceramic powders are added to the L12 base powder to form bimodal and trimodal consolidated alloy microstructures.
  • Cryomilling (step 240) can be used to refine the grain size of gas atomized L12 aluminum alloy powder as well as the final consolidated alloy microstructure. Cryomilling is described in U.S. Pat. No. 6,902,699, Fritzemeier et al. and in U.S. Pat. No. 7,344,675, Van Daam et al. and are incorporated herein in their entirety by reference. Cryomilling involves high-energy ball milling under liquid nitrogen. The liquid nitrogen environment prevents oxidation and prevents frictional heating of the powder and the resulting grain coarsening. During the process, the powder particles are repeatedly sheared, fractured and cold welded which results in a severely deformed microstructure containing a high dislocation density that, with continued deformation, evolves into a cellular structure consisting of extremely small dislocation free grains separated by high angle grain boundaries with high dislocation density. The grain size of the cellular microstructure is typically less than 100 nm and the microstructure is considered a nanostructure.
  • In addition, the nitrogen environment results in the formation of nitride particles that reside at the grain boundaries and in the grain interiors and resist coarsening at higher temperatures. Stearic acid is preferably added to the powder charge to prevent excessive agglomeration and to promote fracturing and rewelding of the L12 aluminum alloy particles during milling.
  • Following sieving (step 220), blending (step 230) and (optionally) cryomilling (Step 240), the powders are transferred to a can (step 250) where the powder is vacuum degassed (step 260) for about 12 hours to over 8 days at elevated temperatures. A temperature range of about 500° F. (260° C.) to about 900° F. (482° C.) is preferred and about 750° F. (399° C.) is more preferred. Dynamic degassing large amounts of powder are preferred to static degassing to expose all of the powder to a uniform temperature. Degassing removes the stearic acid lubricant as well as oxygen and hydrogen from the charge.
  • Following vacuum degassing (step 260), the vacuum line is crimped and welded shut. The powder is then consolidated into a dense preform by closed die forging or by quasi-isostatic forging (step 280).
  • An exemplary embodiment of this invention is to consolidate a canned L12 aluminum alloy powder preform into a substantially 100% dense billet by a quasi-isostatic Ceracon-type forging process. The forging process consists of uniaxially pressing the canned powder preform or solid part perform in a cylindrical press, wherein the preform is surrounded by pressure transmitting medium during the forging. A schematic of quasi-isostatic forging equipment 300 is shown in FIG. 9. The equipment consists of cylindrical die 310 on base 320 with ram 325 inserted in dye cavity 330. Quasi-isostatic forging is carried out at elevated temperatures schematically illustrated by heating coils 340.
  • The steps to Ceracon forge an L12 aluminum alloy powder preform are schematically illustrated in FIG. 10. First die 310 is (optionally) preheated (step 410). The preheat temperature is preferably about 40° F. (4.5° C.) to about 70° F. (21° C.) higher than the forging temperature to compensate for cooling that occurs during the loading process. Next, die 310 is partially filled with pressure transmitting medium (PTM) 360 (step 420). To minimize cooling during loading, the PTM can be (optionally) preheated, preferably to the same temperature as the die (step 450). Consolidated powder preform 350 is then inserted in die 310. Powder preform 350 can be (optionally) preheated, preferably to the same temperature as die 310 and PTM 360 (step 440). The remainder of die cavity 330 is then filled with PTM 360 (step 460). PTM 360 can be (optionally) heated (step 450). Ram 325 is then inserted in die 310 (step 470). Pressure 370 is applied to Ram 325 to accomplish quasi-isostatic Ceracon type forging.
  • One advantage of this process is that, if the preheating steps are followed, the run time is short. As a result, deleterious microstructural changes, such as grain and particle coarsening are minimized. Run times can be as short as a few minutes if automated loading and charging equipment is used. In addition, ductility and fracture toughness of L12 based aluminum alloys can be improved significantly due to dynamic bimodal pressure generated during quasi-isostatic forging that breaks apart continuous powder surface oxide that prevent metal-metal bonding and randomly distributes them in the material.
  • During quasi-isostatic Ceracon type forging, the billet axially deforms about 30% and radially deforms about 10%. Strain rates from about 0.1 min−1 to about 6 min−1 at forging temperatures from about 400° F. (204° C.) to about 900° F. (482° C.) are preferred. Forging pressures from 50 ksi (345 MPa) to 150 ksi (1034 MPa) are preferred.
  • The pressure transmitting medium can consist of graphite or other carbon containing powders or ceramic powders or both. By proper fixturing and preheating the pressure transmitting medium and the work piece, densification during quasi-isostatic forging can be extremely rapid on the order of minutes.
  • The unequal actual and radial deformation of the powder is schematically illustrated by the dotted line outline of deformed can 352 in FIG. 9.
  • Quasi-isostatic forging of a 60 to 80 percent dense aluminum-7.5 weight percent magnesium powder preform by this technique resulted in about a 30 percent axial compression and about a 10 percent radial expansion as taught by Meeks III et al. U.S. Pat. No. 7,097,807 and included herein by reference. The quasi-isostatic forging deformation of L12 aluminum alloy powder of this embodiment has beneficial effects on the microstructure and resulting properties of the consolidated billet. The nonuniform stress and resulting strain field in the powder during forging results in extensive shear deformation. The shear deformation deforms the powder, strips off the surface oxide from the L12 alloy particles and redistributes it throughout the consolidated L12 aluminum alloy powder forging as finely divided dispersoids. As a result, there is increased metal to metal contact during forging and resulting increased mechanical integrity. The redistributed surface oxide particles act as additional strengthening agents by resisting dislocation motion by Orowan strengthening.
  • In other embodiments, Ceracon type forging can be followed by hot isostatic pressing (HIP), forging, rolling and other deformation processing techniques.
  • To summarize, Ceracon type quasi-isostatic forging can be low cost and efficient. The pressure transmitting medium, as well as the vacuum sealed preform can be preheated prior to introduction of the perform into the forging chamber to minimize runtime. In addition, forging can be accomplished at high strain rates resulting in forging runs lasting less than a few minutes. The short run time results in an economical process that also further inhibits grain and L12 particle growth in the billet during forging due to the limited time at temperature. The pressure transmitting medium can be reused and loading and unloading the preforms and resulting forgings can be an automated procedure.
  • Although the present invention has been described with reference to preferred embodiments, workers skilled in the art will recognize that changes may be made in form and detail without departing from the spirit and scope of the invention.

Claims (20)

  1. 1. A method for producing high strength aluminum alloy consolidated billets containing L12 dispersoids, comprising the steps of:
    forming an aluminum alloy powder containing L12 dispersoid forming elements,
    wherein the L12 dispersoid forming elements form Al3X dispersoids
    wherein X is at least one first element selected from the group comprising:
    about 0.1 to about 4.0 weight percent scandium, about 0.1 to about 20.0 weight percent erbium, about 0.1 to about 15.0 weight percent thulium, about 0.1 to about 25.0 weight percent ytterbium, and about 0.1 to about 25.0 weight percent lutetium;
    at least one second element selected from the group comprising:
    about 0.1 to about 20.0 weight percent gadolinium, about 0.1 to about 20.0 weight percent yttrium, about 0.05 to about 4.0 weight percent zirconium, about 0.05 to about 10.0 weight percent titanium, about 0.05 to about 10.0 weight percent hafnium, and about 0.05 to about 5.0 weight percent niobium; and
    the balance substantially aluminum;
    placing the powder in a container;
    vacuum degassing the container at about 500° F. (260° C.) to about 900° F. (482° C.) for about 12 hours to about 8 days;
    sealing the container;
    creating a preform by compressing the container by closed die forging or by quasi-isostatic forging to consolidate the powder;
    encompassing the aluminum alloy powder preform with a flowable pressure transmitting medium and heating the encompassed alloy powder; and
    uniaxially compressing the medium to thereby consolidate the aluminum powder; and
    removing the consolidated powder billet.
  2. 2. The method of claim 1, wherein the forging temperature is from about 400° F. (204° C.) to about 900° F. (482° C.).
  3. 3. The method of claim 1, wherein the axial strain rate during compressing the medium is from about 0.1 min−1 to 6 min−1.
  4. 4. The method of claim 1, wherein compressing the medium is at a pressure from about 50 Ksi (345 MPa) to about 150 Ksi (1034 MPa).
  5. 5. The method of claim 1, wherein the preform is vacuum sealed in an aluminum jacket.
  6. 6. The method of claim 1 wherein the pressure transmitting medium comprises graphite or other carbon containing powders or ceramic powders or both.
  7. 7. The method of claim 1 wherein the aluminum alloy powder contains at least one third element selected from the group consisting of silicon, magnesium, lithium, copper, zinc and nickel.
  8. 8. The method of claim 1 wherein the third element comprises at least one of about 4 to about 25 weight percent silicon, about 1 to about 8 weight percent magnesium, about 0.5 to about 3 weight percent lithium, about 0.2 to about 6.5 weight percent copper, about 3 to about 12 weight percent zinc, and about 1 to about 12 weight percent nickel.
  9. 9. A high strength aluminum alloy consolidated billet containing L12 dispersoids in an aluminum alloy matrix wherein:
    the L12 dispersoid forming elements form Al3X dispersoids wherein X is at least
    one first element selected from the group comprising:
    about 0.1 to about 4.0 weight percent scandium, about 0.1 to about 20.0 weight percent erbium, about 0.1 to about 15.0 weight percent thulium, about 0.1 to about 25.0 weight percent ytterbium, and about 0.1 to about 25.0 weight percent lutetium;
    at least one second element selected from the group comprising:
    about 0.1 to about 20.0 weight percent gadolinium, about 0.1 to about 20.0 weight percent yttrium, about 0.05 to about 4.0 weight percent zirconium, about 0.05 to about 10.0 weight percent titanium, about 0.05 to about 10.0 weight percent hafnium, and about 0.05 to about 5.0 weight percent niobium; and
    the balance substantially aluminum, wherein:
    the aluminum alloy consolidated billet containing L12 dispersoids is formed by the steps comprising:
    placing the powder in a container;
    vacuum degassing the container at about 500° F. (260° C.) to about 900° F. (482° C.) for about 12 hours to about 8 days;
    sealing the container;
    creating a preform by compressing the container by closed die forging or by quasi-isostatic forging to consolidate the powder;
    encompassing the aluminum alloy powder preform with a flowable pressure transmitting medium and heating the encompassed alloy powder; and
    encompassing the aluminum alloy powder preform with a flowable pressure transmitting medium and heating the encompassed alloy powder and medium; and
    uniaxially compressing the medium to thereby consolidate the aluminum powder; and
    removing the consolidated powder billet.
  10. 10. The alloy of claim 9, wherein the aluminum alloy powder contains at least one third element selected from the group consisting of silicon, magnesium, lithium, copper, zinc, and nickel.
  11. 11. The alloy of claim 10, wherein the third element comprises at least one of about 4 to about 25 weight percent silicon, about 1 to about 8 weight percent magnesium, about 0.5 to about 3 weight percent lithium, about 0.2 to about 6.5 weight percent copper, about 3 to about 12 weight percent zinc, and about 1 to about 12 weight percent nickel.
  12. 12. The alloy of claim 9, wherein the forging temperature is from about 400° F. (204° C.) to about 900° F. (482° C.).
  13. 13. The alloy of claim 9, wherein the axial strain rate during compressing the medium is from about 0.1 min−1 to 6 min−1.
  14. 14. The alloy of claim 9, wherein compressing the medium is at a pressure from about 50 Ksi (345 MPa) to about 150 Ksi (1034 MPa).
  15. 15. The alloy of claim 9, wherein the pressure transmitting medium comprises graphite or other carbon containing powders or ceramic powders or both.
  16. 16. A high strength aluminum alloy consolidated billet containing L12 dispersoids in an aluminum alloy matrix wherein:
    the L12 dispersoid forming elements form Al3X dispersoids wherein X is at least
    one first element selected from the group comprising:
    about 0.1 to about; 4.0 weight percent scandium, about 0.1 to about 20.0 weight percent erbium, about 0.1 to about 15.0 weight percent thulium, about 0.1 to about 25.0 weight percent ytterbium, and about 0.1 to about 25.0 weight percent lutetium;
    at least one second element selected from the group comprising:
    about 0.1 to about 20.0 weight percent gadolinium, about 0.1 to about 20.0 weight percent yttrium, about 0.05 to about 4.0 weight percent zirconium, about 0.05 to about 10.0 weight percent titanium, about 0.05 to about 10.0 weight percent hafnium, and about 0.05 to about 5.0 weight percent niobium; and
    the balance substantially aluminum, wherein:
    the aluminum alloy consolidated billet containing L12 dispersoids is formed by the steps comprising:
    placing the powder in a container;
    vacuum degassing the container at about 500° F. (260° C.) to about 900° F. (482° C.) for about 12 hours to about 8 days;
    sealing the container;
    creating a preform by compressing the container by closed die forging or by quasi-isostatic forging to consolidate the powder;
    encompassing the aluminum alloy powder preform with a flowable pressure transmitting medium and heating the encompassed alloy powder and medium; and
    uniaxially compressing the medium at an axial strain rate of from about 0.1 min−1 to about 6 min−1 at a pressure from about 50 KSi (345 MPa) to about 150 KSi (1034 MPa) to thereby consolidate the aluminum powder; and
    removing the consolidated powder in a billet.
  17. 17. The high strength aluminum alloy consolidated billet of claim 16 wherein the aluminum alloy powder contains at least one third element selected from the group consisting of silicon, magnesium, lithium, copper, zinc, and nickel.
  18. 18. The high strength aluminum alloy consolidated billet of claim 16 wherein the third element comprises at least one of about 4 to about 25 weight percent silicon, about 1 to about 8 weight percent magnesium, about 0.5 to about 3 weight percent lithium, about 0.2 to about 6.5 weight percent copper, about 3 to about 12 weight percent zinc, and about 1 to about 12 weight percent nickel.
  19. 19. The high strength aluminum alloy consolidated billet of claim 16 wherein the forging temperature is from about 400° F. (204° C.) to about 900° F. (482° C.).
  20. 20. The high strength aluminum alloy consolidated billet of claim 16 wherein the pressure transmitting medium comprises graphite or other carbon containing powders or ceramic powders or both.
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Cited By (2)

* Cited by examiner, † Cited by third party
Publication number Priority date Publication date Assignee Title
US20140373592A1 (en) * 2011-12-29 2014-12-25 Saint Jean Industries Method of dressing a forge die in the implementation of parts obtained by two successive operations of foundry casting followed by forging
GB2536483A (en) * 2015-03-19 2016-09-21 Avic Beijing Inst Of Aeronautical Mat (Avic Biam) A method of forming a metal component

Citations (90)

* Cited by examiner, † Cited by third party
Publication number Priority date Publication date Assignee Title
US3619181A (en) * 1968-10-29 1971-11-09 Aluminum Co Of America Aluminum scandium alloy
US3816080A (en) * 1971-07-06 1974-06-11 Int Nickel Co Mechanically-alloyed aluminum-aluminum oxide
US4041123A (en) * 1971-04-20 1977-08-09 Westinghouse Electric Corporation Method of compacting shaped powdered objects
US4259112A (en) * 1979-04-05 1981-03-31 Dwa Composite Specialties, Inc. Process for manufacture of reinforced composites
US4463058A (en) * 1981-06-16 1984-07-31 Atlantic Richfield Company Silicon carbide whisker composites
US4469537A (en) * 1983-06-27 1984-09-04 Reynolds Metals Company Aluminum armor plate system
US4499048A (en) * 1983-02-23 1985-02-12 Metal Alloys, Inc. Method of consolidating a metallic body
US4597792A (en) * 1985-06-10 1986-07-01 Kaiser Aluminum & Chemical Corporation Aluminum-based composite product of high strength and toughness
US4626294A (en) * 1985-05-28 1986-12-02 Aluminum Company Of America Lightweight armor plate and method
US4647321A (en) * 1980-11-24 1987-03-03 United Technologies Corporation Dispersion strengthened aluminum alloys
US4661172A (en) * 1984-02-29 1987-04-28 Allied Corporation Low density aluminum alloys and method
US4667497A (en) * 1985-10-08 1987-05-26 Metals, Ltd. Forming of workpiece using flowable particulate
US4689090A (en) * 1986-03-20 1987-08-25 Aluminum Company Of America Superplastic aluminum alloys containing scandium
US4710246A (en) * 1982-07-06 1987-12-01 Centre National De La Recherche Scientifique "Cnrs" Amorphous aluminum-based alloys
US4713216A (en) * 1985-04-27 1987-12-15 Showa Aluminum Kabushiki Kaisha Aluminum alloys having high strength and resistance to stress and corrosion
US4755221A (en) * 1986-03-24 1988-07-05 Gte Products Corporation Aluminum based composite powders and process for producing same
US4832741A (en) * 1986-08-12 1989-05-23 Bbc Brown Boveri Ag Powder-metallurgical process for the production of a green pressed article of high strength and of low relative density from a heat-resistant aluminum alloy
US4834810A (en) * 1988-05-06 1989-05-30 Inco Alloys International, Inc. High modulus A1 alloys
US4834942A (en) * 1988-01-29 1989-05-30 The United States Of America As Represented By The Secretary Of The Navy Elevated temperature aluminum-titanium alloy by powder metallurgy process
US4853178A (en) * 1988-11-17 1989-08-01 Ceracon, Inc. Electrical heating of graphite grain employed in consolidation of objects
US4865806A (en) * 1986-05-01 1989-09-12 Dural Aluminum Composites Corp. Process for preparation of composite materials containing nonmetallic particles in a metallic matrix
US4874440A (en) * 1986-03-20 1989-10-17 Aluminum Company Of America Superplastic aluminum products and alloys
US4915605A (en) * 1989-05-11 1990-04-10 Ceracon, Inc. Method of consolidation of powder aluminum and aluminum alloys
US4923532A (en) * 1988-09-12 1990-05-08 Allied-Signal Inc. Heat treatment for aluminum-lithium based metal matrix composites
US4927470A (en) * 1988-10-12 1990-05-22 Aluminum Company Of America Thin gauge aluminum plate product by isothermal treatment and ramp anneal
US4933140A (en) * 1988-11-17 1990-06-12 Ceracon, Inc. Electrical heating of graphite grain employed in consolidation of objects
US4946517A (en) * 1988-10-12 1990-08-07 Aluminum Company Of America Unrecrystallized aluminum plate product by ramp annealing
US4964927A (en) * 1989-03-31 1990-10-23 University Of Virginia Alumini Patents Aluminum-based metallic glass alloys
US4988464A (en) * 1989-06-01 1991-01-29 Union Carbide Corporation Method for producing powder by gas atomization
US5032352A (en) * 1990-09-21 1991-07-16 Ceracon, Inc. Composite body formation of consolidated powder metal part
US5053084A (en) * 1987-08-12 1991-10-01 Yoshida Kogyo K.K. High strength, heat resistant aluminum alloys and method of preparing wrought article therefrom
US5055257A (en) * 1986-03-20 1991-10-08 Aluminum Company Of America Superplastic aluminum products and alloys
US5059390A (en) * 1989-06-14 1991-10-22 Aluminum Company Of America Dual-phase, magnesium-based alloy having improved properties
US5066342A (en) * 1988-01-28 1991-11-19 Aluminum Company Of America Aluminum-lithium alloys and method of making the same
US5076865A (en) * 1988-10-15 1991-12-31 Yoshida Kogyo K. K. Amorphous aluminum alloys
US5076340A (en) * 1989-08-07 1991-12-31 Dural Aluminum Composites Corp. Cast composite material having a matrix containing a stable oxide-forming element
US5130209A (en) * 1989-11-09 1992-07-14 Allied-Signal Inc. Arc sprayed continuously reinforced aluminum base composites and method
US5133931A (en) * 1990-08-28 1992-07-28 Reynolds Metals Company Lithium aluminum alloy system
US5198045A (en) * 1991-05-14 1993-03-30 Reynolds Metals Company Low density high strength al-li alloy
US5211910A (en) * 1990-01-26 1993-05-18 Martin Marietta Corporation Ultra high strength aluminum-base alloys
US5226983A (en) * 1985-07-08 1993-07-13 Allied-Signal Inc. High strength, ductile, low density aluminum alloys and process for making same
US5256215A (en) * 1990-10-16 1993-10-26 Honda Giken Kogyo Kabushiki Kaisha Process for producing high strength and high toughness aluminum alloy, and alloy material
US5308410A (en) * 1990-12-18 1994-05-03 Honda Giken Kogyo Kabushiki Kaisha Process for producing high strength and high toughness aluminum alloy
US5312494A (en) * 1992-05-06 1994-05-17 Honda Giken Kogyo Kabushiki Kaisha High strength and high toughness aluminum alloy
US5318641A (en) * 1990-06-08 1994-06-07 Tsuyoshi Masumoto Particle-dispersion type amorphous aluminum-alloy having high strength
US5397403A (en) * 1989-12-29 1995-03-14 Honda Giken Kogyo Kabushiki Kaisha High strength amorphous aluminum-based alloy member
US5458700A (en) * 1992-03-18 1995-10-17 Tsuyoshi Masumoto High-strength aluminum alloy
US5462712A (en) * 1988-08-18 1995-10-31 Martin Marietta Corporation High strength Al-Cu-Li-Zn-Mg alloys
US5480470A (en) * 1992-10-16 1996-01-02 General Electric Company Atomization with low atomizing gas pressure
US5532069A (en) * 1993-12-24 1996-07-02 Tsuyoshi Masumoto Aluminum alloy and method of preparing the same
US5597529A (en) * 1994-05-25 1997-01-28 Ashurst Technology Corporation (Ireland Limited) Aluminum-scandium alloys
US5624632A (en) * 1995-01-31 1997-04-29 Aluminum Company Of America Aluminum magnesium alloy product containing dispersoids
US5882449A (en) * 1997-07-11 1999-03-16 Mcdonnell Douglas Corporation Process for preparing aluminum/lithium/scandium rolled sheet products
US6139563A (en) * 1997-09-25 2000-10-31 Allegiance Corporation Surgical device with malleable shaft
US6149737A (en) * 1996-09-09 2000-11-21 Sumitomo Electric Industries Ltd. High strength high-toughness aluminum alloy and method of preparing the same
US6248453B1 (en) * 1999-12-22 2001-06-19 United Technologies Corporation High strength aluminum alloy
US6254704B1 (en) * 1998-05-28 2001-07-03 Sulzer Metco (Us) Inc. Method for preparing a thermal spray powder of chromium carbide and nickel chromium
US6258318B1 (en) * 1998-08-21 2001-07-10 Eads Deutschland Gmbh Weldable, corrosion-resistant AIMG alloys, especially for manufacturing means of transportation
US6309594B1 (en) * 1999-06-24 2001-10-30 Ceracon, Inc. Metal consolidation process employing microwave heated pressure transmitting particulate
US6312643B1 (en) * 1997-10-24 2001-11-06 The United States Of America As Represented By The Secretary Of The Air Force Synthesis of nanoscale aluminum alloy powders and devices therefrom
US6315948B1 (en) * 1998-08-21 2001-11-13 Daimler Chrysler Ag Weldable anti-corrosive aluminum-magnesium alloy containing a high amount of magnesium, especially for use in automobiles
US6331218B1 (en) * 1994-11-02 2001-12-18 Tsuyoshi Masumoto High strength and high rigidity aluminum-based alloy and production method therefor
US20010054247A1 (en) * 2000-05-18 2001-12-27 Stall Thomas C. Scandium containing aluminum alloy firearm
US6355209B1 (en) * 1999-11-16 2002-03-12 Ceracon, Inc. Metal consolidation process applicable to functionally gradient material (FGM) compositons of tungsten, nickel, iron, and cobalt
US6368427B1 (en) * 1999-09-10 2002-04-09 Geoffrey K. Sigworth Method for grain refinement of high strength aluminum casting alloys
US6506503B1 (en) * 1998-07-29 2003-01-14 Miba Gleitlager Aktiengesellschaft Friction bearing having an intermediate layer, notably binding layer, made of an alloy on aluminium basis
US6517954B1 (en) * 1998-07-29 2003-02-11 Miba Gleitlager Aktiengesellschaft Aluminium alloy, notably for a layer
US6524410B1 (en) * 2001-08-10 2003-02-25 Tri-Kor Alloys, Llc Method for producing high strength aluminum alloy welded structures
US6531004B1 (en) * 1998-08-21 2003-03-11 Eads Deutschland Gmbh Weldable anti-corrosive aluminium-magnesium alloy containing a high amount of magnesium, especially for use in aviation
US6562154B1 (en) * 2000-06-12 2003-05-13 Aloca Inc. Aluminum sheet products having improved fatigue crack growth resistance and methods of making same
US6630008B1 (en) * 2000-09-18 2003-10-07 Ceracon, Inc. Nanocrystalline aluminum metal matrix composites, and production methods
US20030192627A1 (en) * 2002-04-10 2003-10-16 Lee Jonathan A. High strength aluminum alloy for high temperature applications
US6702982B1 (en) * 1995-02-28 2004-03-09 The United States Of America As Represented By The Secretary Of The Army Aluminum-lithium alloy
US20040046402A1 (en) * 2002-09-05 2004-03-11 Michael Winardi Drive-in latch with rotational adjustment
US20040055671A1 (en) * 2002-04-24 2004-03-25 Questek Innovations Llc Nanophase precipitation strengthened Al alloys processed through the amorphous state
US20040089382A1 (en) * 2002-11-08 2004-05-13 Senkov Oleg N. Method of making a high strength aluminum alloy composition
US20040170522A1 (en) * 2003-02-28 2004-09-02 Watson Thomas J. Aluminum base alloys
US20040191111A1 (en) * 2003-03-14 2004-09-30 Beijing University Of Technology Er strengthening aluminum alloy
US20050013725A1 (en) * 2003-07-14 2005-01-20 Chung-Chih Hsiao Aluminum based material having high conductivity
US6902699B2 (en) * 2002-10-02 2005-06-07 The Boeing Company Method for preparing cryomilled aluminum alloys and components extruded and forged therefrom
US20050147520A1 (en) * 2003-12-31 2005-07-07 Guido Canzona Method for improving the ductility of high-strength nanophase alloys
US20060011272A1 (en) * 2004-07-15 2006-01-19 Lin Jen C 2000 Series alloys with enhanced damage tolerance performance for aerospace applications
US20060093512A1 (en) * 2003-01-15 2006-05-04 Pandey Awadh B Aluminum based alloy
US20060172073A1 (en) * 2005-02-01 2006-08-03 Groza Joanna R Methods for production of FGM net shaped body for various applications
US20060269437A1 (en) * 2005-05-31 2006-11-30 Pandey Awadh B High temperature aluminum alloys
US20070048167A1 (en) * 2005-08-25 2007-03-01 Yutaka Yano Metal particles, process for manufacturing the same, and process for manufacturing vehicle components therefrom
US20070062669A1 (en) * 2005-09-21 2007-03-22 Song Shihong G Method of producing a castable high temperature aluminum alloy by controlled solidification
US7241328B2 (en) * 2003-11-25 2007-07-10 The Boeing Company Method for preparing ultra-fine, submicron grain titanium and titanium-alloy articles and articles prepared thereby
US7344675B2 (en) * 2003-03-12 2008-03-18 The Boeing Company Method for preparing nanostructured metal alloys having increased nitride content
US20080066833A1 (en) * 2006-09-19 2008-03-20 Lin Jen C HIGH STRENGTH, HIGH STRESS CORROSION CRACKING RESISTANT AND CASTABLE Al-Zn-Mg-Cu-Zr ALLOY FOR SHAPE CAST PRODUCTS

Family Cites Families (3)

* Cited by examiner, † Cited by third party
Publication number Priority date Publication date Assignee Title
US5061323A (en) * 1990-10-15 1991-10-29 The United States Of America As Represented By The Secretary Of The Navy Composition and method for producing an aluminum alloy resistant to environmentally-assisted cracking
US5330704A (en) * 1991-02-04 1994-07-19 Alliedsignal Inc. Method for producing aluminum powder alloy products having lower gas contents
US7691214B2 (en) * 2005-05-26 2010-04-06 Honeywell International, Inc. High strength aluminum alloys for aircraft wheel and brake components

Patent Citations (95)

* Cited by examiner, † Cited by third party
Publication number Priority date Publication date Assignee Title
US3619181A (en) * 1968-10-29 1971-11-09 Aluminum Co Of America Aluminum scandium alloy
US4041123A (en) * 1971-04-20 1977-08-09 Westinghouse Electric Corporation Method of compacting shaped powdered objects
US3816080A (en) * 1971-07-06 1974-06-11 Int Nickel Co Mechanically-alloyed aluminum-aluminum oxide
US4259112A (en) * 1979-04-05 1981-03-31 Dwa Composite Specialties, Inc. Process for manufacture of reinforced composites
US4647321A (en) * 1980-11-24 1987-03-03 United Technologies Corporation Dispersion strengthened aluminum alloys
US4463058A (en) * 1981-06-16 1984-07-31 Atlantic Richfield Company Silicon carbide whisker composites
US4710246A (en) * 1982-07-06 1987-12-01 Centre National De La Recherche Scientifique "Cnrs" Amorphous aluminum-based alloys
US4499048A (en) * 1983-02-23 1985-02-12 Metal Alloys, Inc. Method of consolidating a metallic body
US4469537A (en) * 1983-06-27 1984-09-04 Reynolds Metals Company Aluminum armor plate system
US4661172A (en) * 1984-02-29 1987-04-28 Allied Corporation Low density aluminum alloys and method
US4713216A (en) * 1985-04-27 1987-12-15 Showa Aluminum Kabushiki Kaisha Aluminum alloys having high strength and resistance to stress and corrosion
US4626294A (en) * 1985-05-28 1986-12-02 Aluminum Company Of America Lightweight armor plate and method
US4597792A (en) * 1985-06-10 1986-07-01 Kaiser Aluminum & Chemical Corporation Aluminum-based composite product of high strength and toughness
US5226983A (en) * 1985-07-08 1993-07-13 Allied-Signal Inc. High strength, ductile, low density aluminum alloys and process for making same
US4667497A (en) * 1985-10-08 1987-05-26 Metals, Ltd. Forming of workpiece using flowable particulate
US4874440A (en) * 1986-03-20 1989-10-17 Aluminum Company Of America Superplastic aluminum products and alloys
US4689090A (en) * 1986-03-20 1987-08-25 Aluminum Company Of America Superplastic aluminum alloys containing scandium
US5055257A (en) * 1986-03-20 1991-10-08 Aluminum Company Of America Superplastic aluminum products and alloys
US4755221A (en) * 1986-03-24 1988-07-05 Gte Products Corporation Aluminum based composite powders and process for producing same
US4865806A (en) * 1986-05-01 1989-09-12 Dural Aluminum Composites Corp. Process for preparation of composite materials containing nonmetallic particles in a metallic matrix
US4832741A (en) * 1986-08-12 1989-05-23 Bbc Brown Boveri Ag Powder-metallurgical process for the production of a green pressed article of high strength and of low relative density from a heat-resistant aluminum alloy
US5053084A (en) * 1987-08-12 1991-10-01 Yoshida Kogyo K.K. High strength, heat resistant aluminum alloys and method of preparing wrought article therefrom
US5066342A (en) * 1988-01-28 1991-11-19 Aluminum Company Of America Aluminum-lithium alloys and method of making the same
US4834942A (en) * 1988-01-29 1989-05-30 The United States Of America As Represented By The Secretary Of The Navy Elevated temperature aluminum-titanium alloy by powder metallurgy process
US4834810A (en) * 1988-05-06 1989-05-30 Inco Alloys International, Inc. High modulus A1 alloys
US5462712A (en) * 1988-08-18 1995-10-31 Martin Marietta Corporation High strength Al-Cu-Li-Zn-Mg alloys
US4923532A (en) * 1988-09-12 1990-05-08 Allied-Signal Inc. Heat treatment for aluminum-lithium based metal matrix composites
US4946517A (en) * 1988-10-12 1990-08-07 Aluminum Company Of America Unrecrystallized aluminum plate product by ramp annealing
US4927470A (en) * 1988-10-12 1990-05-22 Aluminum Company Of America Thin gauge aluminum plate product by isothermal treatment and ramp anneal
US5076865A (en) * 1988-10-15 1991-12-31 Yoshida Kogyo K. K. Amorphous aluminum alloys
US4853178A (en) * 1988-11-17 1989-08-01 Ceracon, Inc. Electrical heating of graphite grain employed in consolidation of objects
US4933140A (en) * 1988-11-17 1990-06-12 Ceracon, Inc. Electrical heating of graphite grain employed in consolidation of objects
US4964927A (en) * 1989-03-31 1990-10-23 University Of Virginia Alumini Patents Aluminum-based metallic glass alloys
US4915605A (en) * 1989-05-11 1990-04-10 Ceracon, Inc. Method of consolidation of powder aluminum and aluminum alloys
US4988464A (en) * 1989-06-01 1991-01-29 Union Carbide Corporation Method for producing powder by gas atomization
US5059390A (en) * 1989-06-14 1991-10-22 Aluminum Company Of America Dual-phase, magnesium-based alloy having improved properties
US5076340A (en) * 1989-08-07 1991-12-31 Dural Aluminum Composites Corp. Cast composite material having a matrix containing a stable oxide-forming element
US5130209A (en) * 1989-11-09 1992-07-14 Allied-Signal Inc. Arc sprayed continuously reinforced aluminum base composites and method
US5397403A (en) * 1989-12-29 1995-03-14 Honda Giken Kogyo Kabushiki Kaisha High strength amorphous aluminum-based alloy member
US5211910A (en) * 1990-01-26 1993-05-18 Martin Marietta Corporation Ultra high strength aluminum-base alloys
US5318641A (en) * 1990-06-08 1994-06-07 Tsuyoshi Masumoto Particle-dispersion type amorphous aluminum-alloy having high strength
US5133931A (en) * 1990-08-28 1992-07-28 Reynolds Metals Company Lithium aluminum alloy system
US5032352A (en) * 1990-09-21 1991-07-16 Ceracon, Inc. Composite body formation of consolidated powder metal part
US5256215A (en) * 1990-10-16 1993-10-26 Honda Giken Kogyo Kabushiki Kaisha Process for producing high strength and high toughness aluminum alloy, and alloy material
US5308410A (en) * 1990-12-18 1994-05-03 Honda Giken Kogyo Kabushiki Kaisha Process for producing high strength and high toughness aluminum alloy
US5198045A (en) * 1991-05-14 1993-03-30 Reynolds Metals Company Low density high strength al-li alloy
US5458700A (en) * 1992-03-18 1995-10-17 Tsuyoshi Masumoto High-strength aluminum alloy
US5312494A (en) * 1992-05-06 1994-05-17 Honda Giken Kogyo Kabushiki Kaisha High strength and high toughness aluminum alloy
US5480470A (en) * 1992-10-16 1996-01-02 General Electric Company Atomization with low atomizing gas pressure
US5532069A (en) * 1993-12-24 1996-07-02 Tsuyoshi Masumoto Aluminum alloy and method of preparing the same
US5597529A (en) * 1994-05-25 1997-01-28 Ashurst Technology Corporation (Ireland Limited) Aluminum-scandium alloys
US5620652A (en) * 1994-05-25 1997-04-15 Ashurst Technology Corporation (Ireland) Limited Aluminum alloys containing scandium with zirconium additions
US6331218B1 (en) * 1994-11-02 2001-12-18 Tsuyoshi Masumoto High strength and high rigidity aluminum-based alloy and production method therefor
US5624632A (en) * 1995-01-31 1997-04-29 Aluminum Company Of America Aluminum magnesium alloy product containing dispersoids
US6702982B1 (en) * 1995-02-28 2004-03-09 The United States Of America As Represented By The Secretary Of The Army Aluminum-lithium alloy
US6149737A (en) * 1996-09-09 2000-11-21 Sumitomo Electric Industries Ltd. High strength high-toughness aluminum alloy and method of preparing the same
US5882449A (en) * 1997-07-11 1999-03-16 Mcdonnell Douglas Corporation Process for preparing aluminum/lithium/scandium rolled sheet products
US6139563A (en) * 1997-09-25 2000-10-31 Allegiance Corporation Surgical device with malleable shaft
US6312643B1 (en) * 1997-10-24 2001-11-06 The United States Of America As Represented By The Secretary Of The Air Force Synthesis of nanoscale aluminum alloy powders and devices therefrom
US6254704B1 (en) * 1998-05-28 2001-07-03 Sulzer Metco (Us) Inc. Method for preparing a thermal spray powder of chromium carbide and nickel chromium
US6517954B1 (en) * 1998-07-29 2003-02-11 Miba Gleitlager Aktiengesellschaft Aluminium alloy, notably for a layer
US6506503B1 (en) * 1998-07-29 2003-01-14 Miba Gleitlager Aktiengesellschaft Friction bearing having an intermediate layer, notably binding layer, made of an alloy on aluminium basis
US6315948B1 (en) * 1998-08-21 2001-11-13 Daimler Chrysler Ag Weldable anti-corrosive aluminum-magnesium alloy containing a high amount of magnesium, especially for use in automobiles
US6531004B1 (en) * 1998-08-21 2003-03-11 Eads Deutschland Gmbh Weldable anti-corrosive aluminium-magnesium alloy containing a high amount of magnesium, especially for use in aviation
US6258318B1 (en) * 1998-08-21 2001-07-10 Eads Deutschland Gmbh Weldable, corrosion-resistant AIMG alloys, especially for manufacturing means of transportation
US6309594B1 (en) * 1999-06-24 2001-10-30 Ceracon, Inc. Metal consolidation process employing microwave heated pressure transmitting particulate
US6368427B1 (en) * 1999-09-10 2002-04-09 Geoffrey K. Sigworth Method for grain refinement of high strength aluminum casting alloys
US6355209B1 (en) * 1999-11-16 2002-03-12 Ceracon, Inc. Metal consolidation process applicable to functionally gradient material (FGM) compositons of tungsten, nickel, iron, and cobalt
US6248453B1 (en) * 1999-12-22 2001-06-19 United Technologies Corporation High strength aluminum alloy
US20010054247A1 (en) * 2000-05-18 2001-12-27 Stall Thomas C. Scandium containing aluminum alloy firearm
US6562154B1 (en) * 2000-06-12 2003-05-13 Aloca Inc. Aluminum sheet products having improved fatigue crack growth resistance and methods of making same
US6630008B1 (en) * 2000-09-18 2003-10-07 Ceracon, Inc. Nanocrystalline aluminum metal matrix composites, and production methods
US7097807B1 (en) * 2000-09-18 2006-08-29 Ceracon, Inc. Nanocrystalline aluminum alloy metal matrix composites, and production methods
US6524410B1 (en) * 2001-08-10 2003-02-25 Tri-Kor Alloys, Llc Method for producing high strength aluminum alloy welded structures
US20030192627A1 (en) * 2002-04-10 2003-10-16 Lee Jonathan A. High strength aluminum alloy for high temperature applications
US6918970B2 (en) * 2002-04-10 2005-07-19 The United States Of America As Represented By The Administrator Of The National Aeronautics And Space Administration High strength aluminum alloy for high temperature applications
US20040055671A1 (en) * 2002-04-24 2004-03-25 Questek Innovations Llc Nanophase precipitation strengthened Al alloys processed through the amorphous state
US20040046402A1 (en) * 2002-09-05 2004-03-11 Michael Winardi Drive-in latch with rotational adjustment
US6902699B2 (en) * 2002-10-02 2005-06-07 The Boeing Company Method for preparing cryomilled aluminum alloys and components extruded and forged therefrom
US7048815B2 (en) * 2002-11-08 2006-05-23 Ues, Inc. Method of making a high strength aluminum alloy composition
US20040089382A1 (en) * 2002-11-08 2004-05-13 Senkov Oleg N. Method of making a high strength aluminum alloy composition
US20060093512A1 (en) * 2003-01-15 2006-05-04 Pandey Awadh B Aluminum based alloy
US20040170522A1 (en) * 2003-02-28 2004-09-02 Watson Thomas J. Aluminum base alloys
US6974510B2 (en) * 2003-02-28 2005-12-13 United Technologies Corporation Aluminum base alloys
US7344675B2 (en) * 2003-03-12 2008-03-18 The Boeing Company Method for preparing nanostructured metal alloys having increased nitride content
US20040191111A1 (en) * 2003-03-14 2004-09-30 Beijing University Of Technology Er strengthening aluminum alloy
US20050013725A1 (en) * 2003-07-14 2005-01-20 Chung-Chih Hsiao Aluminum based material having high conductivity
US7241328B2 (en) * 2003-11-25 2007-07-10 The Boeing Company Method for preparing ultra-fine, submicron grain titanium and titanium-alloy articles and articles prepared thereby
US20050147520A1 (en) * 2003-12-31 2005-07-07 Guido Canzona Method for improving the ductility of high-strength nanophase alloys
US20060011272A1 (en) * 2004-07-15 2006-01-19 Lin Jen C 2000 Series alloys with enhanced damage tolerance performance for aerospace applications
US20060172073A1 (en) * 2005-02-01 2006-08-03 Groza Joanna R Methods for production of FGM net shaped body for various applications
US20060269437A1 (en) * 2005-05-31 2006-11-30 Pandey Awadh B High temperature aluminum alloys
US20070048167A1 (en) * 2005-08-25 2007-03-01 Yutaka Yano Metal particles, process for manufacturing the same, and process for manufacturing vehicle components therefrom
US20070062669A1 (en) * 2005-09-21 2007-03-22 Song Shihong G Method of producing a castable high temperature aluminum alloy by controlled solidification
US20080066833A1 (en) * 2006-09-19 2008-03-20 Lin Jen C HIGH STRENGTH, HIGH STRESS CORROSION CRACKING RESISTANT AND CASTABLE Al-Zn-Mg-Cu-Zr ALLOY FOR SHAPE CAST PRODUCTS

Non-Patent Citations (2)

* Cited by examiner, † Cited by third party
Title
James, R.S. "Aluminum-Lithium Alloys," Properties and Selection: Nonferrous Alloys and Special-Purpose Materials, Vol 2, ASM Handbook, ASM International, 1990. p 178-199. *
Lyle et al. "Aluminum Alloys." Ullmann's Encyclopedia of Industrial Chemistry. Available online 15 June 2000. *

Cited By (2)

* Cited by examiner, † Cited by third party
Publication number Priority date Publication date Assignee Title
US20140373592A1 (en) * 2011-12-29 2014-12-25 Saint Jean Industries Method of dressing a forge die in the implementation of parts obtained by two successive operations of foundry casting followed by forging
GB2536483A (en) * 2015-03-19 2016-09-21 Avic Beijing Inst Of Aeronautical Mat (Avic Biam) A method of forming a metal component

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