EP2403967A2 - Alliages d'aluminium l1 2 haute résistance produits par cryobroyage - Google Patents

Alliages d'aluminium l1 2 haute résistance produits par cryobroyage

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Publication number
EP2403967A2
EP2403967A2 EP10749389A EP10749389A EP2403967A2 EP 2403967 A2 EP2403967 A2 EP 2403967A2 EP 10749389 A EP10749389 A EP 10749389A EP 10749389 A EP10749389 A EP 10749389A EP 2403967 A2 EP2403967 A2 EP 2403967A2
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EP
European Patent Office
Prior art keywords
weight percent
aluminum
dispersoids
billet
aluminum alloy
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Application number
EP10749389A
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German (de)
English (en)
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EP2403967A4 (fr
Inventor
Awadh B. Pandey
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Raytheon Technologies Corp
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United Technologies Corp
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Publication of EP2403967A2 publication Critical patent/EP2403967A2/fr
Publication of EP2403967A4 publication Critical patent/EP2403967A4/fr
Withdrawn legal-status Critical Current

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    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C21/00—Alloys based on aluminium
    • C22C21/003—Alloys based on aluminium containing at least 2.6% of one or more of the elements: tin, lead, antimony, bismuth, cadmium, and titanium
    • 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
    • B22F9/00—Making metallic powder or suspensions thereof
    • B22F9/02—Making metallic powder or suspensions thereof using physical processes
    • B22F9/04—Making metallic powder or suspensions thereof using physical processes starting from solid material, e.g. by crushing, grinding or milling
    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C1/00—Making non-ferrous alloys
    • C22C1/04—Making non-ferrous alloys by powder metallurgy
    • C22C1/0408—Light metal alloys
    • C22C1/0416—Aluminium-based alloys
    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C21/00—Alloys based on aluminium
    • C22C21/02—Alloys based on aluminium with silicon as the next major constituent
    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C21/00—Alloys based on aluminium
    • C22C21/06—Alloys based on aluminium with magnesium as the next major constituent
    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C21/00—Alloys based on aluminium
    • C22C21/10—Alloys based on aluminium with zinc as the next major constituent
    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C21/00—Alloys based on aluminium
    • C22C21/12—Alloys based on aluminium with copper as the next major constituent
    • 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
    • B22F3/12—Both compacting and sintering
    • B22F3/14—Both compacting and sintering simultaneously
    • B22F2003/145—Both compacting and sintering simultaneously by warm compacting, below debindering temperature
    • 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
    • B22F9/00—Making metallic powder or suspensions thereof
    • B22F9/02—Making metallic powder or suspensions thereof using physical processes
    • B22F9/04—Making metallic powder or suspensions thereof using physical processes starting from solid material, e.g. by crushing, grinding or milling
    • B22F2009/041—Making metallic powder or suspensions thereof using physical processes starting from solid material, e.g. by crushing, grinding or milling by mechanical alloying, e.g. blending, milling
    • 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
    • B22F2998/00—Supplementary information concerning processes or compositions relating to powder metallurgy
    • B22F2998/10—Processes characterised by the sequence of their steps
    • 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
    • B22F2999/00—Aspects linked to processes or compositions used in powder metallurgy

Definitions

  • the present invention relates generally to aluminum alloys and more specifically to a method for forming high strength aluminum alloy powder having Ll 2 dispersoids therein.
  • 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.
  • U.S. Patent No. 6,248,453 owned by the assignee of the present invention discloses aluminum alloys strengthened by dispersed Al 3 X Ll 2 intermetallic phases where X is selected from the group consisting of Sc, Er, Lu, Yb, Tm, and Lu.
  • the Al 3 X 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 Ll 2 aluminum alloys are stable up to 572°F (300°C).
  • Ll 2 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.
  • the present invention is a method for consolidating cryomilled aluminum alloy powders into useful components with high temperature strength and fracture toughness.
  • powders include an aluminum alloy having coherent Ll 2 Al 3 X 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 powders are classified by sieving and blended to improve homogeneity. Cryomilling is an essential step in the manufacturing process. The cryomilled powders are then vacuum degassed in a container that is then sealed.
  • the sealed container i.e. can
  • the sealed container is compressed by vacuum hot pressing, hot isostatic pressing or blind die compaction to densify the powder charge.
  • the can is removed and the billet is extruded, forged and/or rolled into useful shapes with high temperature strength and fracture toughness.
  • 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. 6A and 6B are SEM photos of the gas atomized inventive Ll 2 aluminum alloy powder.
  • FIG. 7A and 7B are photomicrographs of cross-sections showing the cellular microstructure of the gas atomized inventive Ll 2 aluminum alloy powder.
  • FIG. 8A and 8B are photomicrographs of cryomilled powder of the inventive Ll 2 aluminum alloy powder.
  • FIG. 9A and 9B are photomicrographs of cross-sections of cryomilled powder of the inventive Ll 2 aluminum alloy powder.
  • FIG. 10 is a diagram showing the processing steps to consolidate Ll 2 aluminum alloy powder.
  • FIG. 11 is a photo of a 3-inch diameter copper jacketed Ll 2 aluminum alloy billet.
  • FIG. 12 is a photo of extrusion dies for 3-inch diameter billet.
  • FIG. 13 is a photo of extruded Ll 2 aluminum alloy rods from 3-inch diameter billets.
  • FIG. 14 is a photo of machined Ll 2 aluminum alloy billets.
  • FIG. 15 is a photo of a machined three-piece Ll 2 aluminum alloy billet assembly for 6-inch copper jacketed extrusion billet.
  • FIG. 16 is a photo of extruded Ll 2 aluminum alloy rods from 6-inch diameter billets. DETAILED DESCRIPTION
  • Alloy powders refined by this invention are formed from aluminum based alloys with high strength and fracture toughness for applications at temperatures from about - 42O 0 F (-251 0 C) up to about 65O 0 F (343 0 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 Ll 2 Al 3 X 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 0 F (577 0 C). There is little solubility of silicon in aluminum at temperatures up to 93O 0 F (500 0 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 0 F (45O 0 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 0 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 0 F (548 0 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 0 F (381 0 C). Zinc has maximum solid solubility of 83.1 weight percent in aluminum at 717.8 0 F (381 0 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 0 F (639.9 0 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 Ll 2 intermetallic Al 3 Ni.
  • scandium, erbium, thulium, ytterbium, and lutetium are potent strengtheners that have low diffusivity and low solubility in aluminum. All these elements form equilibrium Al 3 X intermetallic dispersoids where X is at least one of scandium, erbium, thulium, ytterbium, and lutetium, that have an Ll 2 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 Al 3 Sc dispersoids that are fine and coherent with the aluminum matrix.
  • Lattice parameters of aluminum and Al 3 Sc 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 Al 3 Sc dispersoids.
  • This low interfacial energy makes the Al 3 Sc dispersoids thermally stable and resistant to coarsening up to temperatures as high as about 842 0 F (45O 0 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 Al 3 Sc to coarsening.
  • Additions of zinc, copper, lithium, silicon, and nickel provide solid solution and precipitation strengthening in the aluminum alloys.
  • Al 3 Sc 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 Al 3 Sc in solution.
  • suitable alloying elements such as gadolinium, yttrium, zirconium, titanium, hafnium, niobium, or combinations thereof, that enter Al 3 Sc in solution.
  • Erbium forms Al 3 Er dispersoids in the aluminum matrix that are fine and coherent with the aluminum matrix.
  • the lattice parameters of aluminum and Al 3 Er are close (0.405 nm and 0.417 nm respectively), indicating there is minimal driving force for causing growth of the Al 3 Er dispersoids.
  • This low interfacial energy makes the Al 3 Er dispersoids thermally stable and resistant to coarsening up to temperatures as high as about 842 0 F (45O 0 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 Al 3 Er to coarsening.
  • Additions of zinc, copper, lithium, silicon, and nickel provide solid solution and precipitation strengthening in the aluminum alloys.
  • Al 3 Er 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 Al 3 Er in solution.
  • suitable alloying elements such as gadolinium, yttrium, zirconium, titanium, hafnium, niobium, or combinations thereof that enter Al 3 Er in solution.
  • Thulium forms metastable Al 3 Tm dispersoids in the aluminum matrix that are fine and coherent with the aluminum matrix.
  • the lattice parameters of aluminum and Al 3 Tm are close (0.405 nm and 0.420 nm respectively), indicating there is minimal driving force for causing growth of the Al 3 Tm dispersoids.
  • This low interfacial energy makes the Al 3 Tm dispersoids thermally stable and resistant to coarsening up to temperatures as high as about 842 0 F (45O 0 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 Al 3 Tm to coarsening.
  • Al 3 Tm 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 Al 3 Tm in solution.
  • suitable alloying elements such as gadolinium, yttrium, zirconium, titanium, hafnium, niobium, or combinations thereof that enter Al 3 Tm in solution.
  • Ytterbium forms Al 3 Yb dispersoids in the aluminum matrix that are fine and coherent with the aluminum matrix.
  • the lattice parameters of Al and Al 3 Yb are close (0.405 nm and 0.420 nm respectively), indicating there is minimal driving force for causing growth of the Al 3 Yb dispersoids.
  • Al 3 Yb dispersoids thermally stable and resistant to coarsening up to temperatures as high as about 842 0 F (45O 0 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 Al 3 Yb to coarsening.
  • Additions of zinc, copper, lithium, silicon, and nickel provide solid solution and precipitation strengthening in the aluminum alloys.
  • These Al 3 Yb 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 Al 3 Yb in solution.
  • Al 3 Lu dispersoids forms Al 3 Lu dispersoids in the aluminum matrix that are fine and coherent with the aluminum matrix.
  • the lattice parameters of Al and Al 3 Lu are close (0.405 nm and 0.419 nm respectively), indicating there is minimal driving force for causing growth of the Al 3 Lu dispersoids.
  • This low interfacial energy makes the Al 3 Lu dispersoids thermally stable and resistant to coarsening up to temperatures as high as about 842 0 F (45O 0 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 Al 3 Lu to coarsening.
  • Additions of zinc, copper, lithium, silicon, and nickel provide solid solution and precipitation strengthening in the aluminum alloys.
  • Al 3 Lu 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 Al 3 Lu in solution.
  • suitable alloying elements such as gadolinium, yttrium, zirconium, titanium, hafnium, niobium, or mixtures thereof that enter Al 3 Lu in solution.
  • Gadolinium forms metastable Al 3 Gd dispersoids in the aluminum matrix that are stable up to temperatures as high as about 842 0 F (45O 0 C) due to their low diffusivity in aluminum.
  • the Al 3 Gd dispersoids have a DO ⁇ structure in the equilibrium condition.
  • gadolinium has fairly high solubility in the Al 3 X intermetallic dispersoids (where X is scandium, erbium, thulium, ytterbium or lutetium).
  • Gadolinium can substitute for the X atoms in Al 3 X intermetallic, thereby forming an ordered Ll 2 phase which results in improved thermal and structural stability.
  • Yttrium forms metastable Al 3 Y dispersoids in the aluminum matrix that have an Ll 2 structure in the metastable condition and a DOig structure in the equilibrium condition.
  • the metastable Al 3 Y dispersoids have a low diffusion coefficient, which makes them thermally stable and highly resistant to coarsening.
  • Yttrium has a high solubility in the Al 3 X intermetallic dispersoids allowing large amounts of yttrium to substitute for X in the Al 3 X Ll 2 dispersoids, which results in improved thermal and structural stability.
  • Zirconium forms Al 3 Zr dispersoids in the aluminum matrix that have an Ll 2 structure in the metastable condition and DO 23 structure in the equilibrium condition.
  • the metastable Al 3 Zr dispersoids have a low diffusion coefficient, which makes them thermally stable and highly resistant to coarsening.
  • Zirconium has a high solubility in the Al 3 X dispersoids allowing large amounts of zirconium to substitute for X in the Al 3 X dispersoids, which results in improved thermal and structural stability.
  • Titanium forms Al 3 Ti dispersoids in the aluminum matrix that have an Ll 2 structure in the metastable condition and DO 22 structure in the equilibrium condition.
  • the metastable Al 3 Ti despersoids have a low diffusion coefficient, which makes them thermally stable and highly resistant to coarsening.
  • Titanium has a high solubility in the Al 3 X dispersoids allowing large amounts of titanium to substitute for X in the Al 3 X dispersoids, which result in improved thermal and structural stability.
  • Hafnium forms metastable Al 3 Hf dispersoids in the aluminum matrix that have an Ll 2 structure in the metastable condition and a DO 23 structure in the equilibrium condition.
  • the Al 3 Hf dispersoids have a low diffusion coefficient, which makes them thermally stable and highly resistant to coarsening.
  • Hafnium has a high solubility in the Al 3 X dispersoids allowing large amounts of hafnium to substitute for scandium, erbium, thulium, ytterbium, and lutetium in the above-mentioned Al 3 X dispersoids, which results in stronger and more thermally stable dispersoids.
  • Niobium forms metastable Al 3 Nb dispersoids in the aluminum matrix that have an Ll 2 structure in the metastable condition and a DO 22 structure in the equilibrium condition.
  • Niobium has a lower solubility in the Al 3 X dispersoids than hafnium or yttrium, allowing relatively lower amounts of niobium than hafnium or yttrium to substitute for X in the Al 3 X dispersoids. Nonetheless, niobium can be very effective in slowing down the coarsening kinetics of the Al 3 X dispersoids because the Al 3 Nb dispersoids are thermally stable. The substitution of niobium for X in the above mentioned Al 3 X dispersoids results in stronger and more thermally stable dispersoids.
  • Al 3 X Ll 2 precipitates improve elevated temperature mechanical properties in aluminum alloys for two reasons.
  • 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 antiphase boundary on the glide plane. The energy to create the anti-phase boundary is the origin of the strengthening.
  • the cubic Ll 2 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.
  • Ll 2 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 Ll 2 dispersoids in the microstructure.
  • 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-(O. l-4)Sc-(0.1-2O)Gd; about Al-M-(0.1-2O)Er-(0.1-2O)Gd; about Al-M-(O. l-15)Tm-(0.1-2O)Gd; about Al-M-(O. l-25)Yb-(0.1-2O)Gd; about Al-M-(O. l-25)Lu-(0.1-2O)Gd; about Al-M-(O. l-4)Sc-(0. l-20)Y; about Al-M-(0.1-2O)Er-(O.
  • M is at least one of about (4-25) weight percent silicon, (1-8) weight percent magnesium, (0.5-3) weight percent lithium, (0.2-3) weight percent copper, (3-12) weight percent zinc, and (1-12) weight percent nickel.
  • the amount of silicon present in the fine grain matrix 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 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 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 may vary from about 0.2 to about 3 weight percent, more preferably from about 0.5 to about 2.5 weight percent, and even more preferably from about 1 to about 2.5 weight percent.
  • the amount of zinc present in the fine grain matrix may vary from about
  • the amount of nickel present in the fine grain matrix 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 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 0 F (659 0 C) resulting in a solid solution of scandium and aluminum and Al 3 Sc 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 Ll 2 intermetallic Al 3 Sc following an aging treatment.
  • Alloys with scandium in excess of the eutectic composition can only retain scandium in solid solution by rapid solidification processing (RSP) where cooling rates are in excess of about 10 3o C/second.
  • RSP rapid solidification processing
  • the amount of erbium present in the fine grain matrix 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 0 F (655 0 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 Ll 2 intermetallic Al 3 Er 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 10 3o C/second.
  • RSP rapid solidification processing
  • the amount of thulium present in the alloys, 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.
  • Thulium forms metastable Al 3 Tm dispersoids in the aluminum matrix that have an Ll 2 structure in the equilibrium condition.
  • the Al 3 Tm 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 Ll 2 intermetallic Al 3 Tm 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 10 3o C/second.
  • RSP rapid solidification processing
  • the amount of ytterbium present in the alloys 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 0 F (625 0 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 Ll 2 intermetallic Al 3 Yb 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 10 3o C/second.
  • RSP rapid solidification processing
  • the amount of lutetium present in the alloys 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 0 F (65O 0 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 Ll 2 intermetallic Al 3 Lu 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 10 3o C/second.
  • RSP rapid solidification processing
  • the amount of gadolinium present in the alloys 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 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 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 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, 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 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.
  • FIGS. 8 A and 8B show photomicrographs of cryomilled powders at different magnifications indicating that the shape of the spherical powder has changed from spherical to irregular due to the milling operation.
  • the micro structure of cryomilled powder shown in FIGS. 9 A and 9B indicates that the cellular structure is refined and there is a more uniform distribution of fine particles present in the powder.
  • the process of consolidating the alloy powders into useful forms is schematically illustrated in FIG. 10.
  • Ll 2 aluminum alloy powders 10 are first classified according to size by sieving
  • step 20 Fine particle sizes are required for optimum mechanical properties in the final part.
  • the starting stock should be at least -325 mesh powder (step 30). Other benefits of blending will be discussed later. Powders are then cryomilled to decrease grain size and improve strength (step 40). 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 cryomilled powders are then put in a can (step 50) and vacuum degassed (step 60). Following vacuum degassing the can is sealed (step 70) under vacuum and hot pressed (step 80) to densify the powder compact. The compact is then hot worked (step 90) to refine the microstructure. Finally, the densified compact is machined into an extrusion billet and extruded (step 100) to produce a product with improved mechanical properties useful for subsequent service as a high temperature Ll 2 strengthened aluminum alloy. Extrusion results in optimum mechanical properties in the extrusion direction. If more uniform directional properties are required, forging and/or rolling (step 110) is necessary.
  • Sieving (step 20) is a preferred step in consolidation because the final mechanical properties relate directly to the particle size. Finer particle size results in finer Ll 2 particle dispersion. Optimum mechanical properties have been observed with -450 mesh (30 micron) powder. Sieving (step 20) also limits the defect size in the powder. Before sieving, the powder is passivated with nitrogen gas in order to improve the efficiency of sieving. If the as-atomized powder is oxygen deficient, the powder can have a tendancy to stick together which will lower the efficiency for sieving. Ultrasonic sieving is preferred for its efficiency. Blending (step 30) is a preferred step in the consolidation process because it results in improved uniformity of particle size distribution.
  • Gas atomized Ll 2 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 30 is also preferred when separate metal and/or ceramic powders are added to the Ll 2 base powder to form bimodal or trimodal consolidated alloy microstructures.
  • Cryomilling (step 40) is a preferred step and is used to refine the grain size of gas atomized Ll 2 aluminum alloy powder as well as the final consolidated alloy microstructure.
  • Cryomilling is described in U.S. Patent No. 6,902,699, Fritzemeier et al. and in U.S. Patent No. 7,344,675, Van Daam et al., both owned by the assignee of the present invention, and are incorporated herein in their entirety by reference.
  • Cryomilling involves high-energy ball milling under liquid nitrogen. The liquid nitrogen facilitates efficient breaking up of powder particles. The liquid nitrogen environment prevents oxidation and prevents frictional heating of the powder and the resulting grain coarsening.
  • the powder particles are repeatedly sheared, fractured and cold welded which results in a severely deformed structure 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 (0.04 microinch) and the structure is considered a nano structure.
  • the nitrogen environment results in the formation of nitride particles that reside at the grain boundaries and inside grains themselves 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 Ll 2 aluminum alloy particles during milling.
  • the powders are transferred to a can (step 50) where the powder is vacuum degassed (step 60) at elevated temperatures.
  • the can (step 50) is an aluminum container having a cylindrical, rectangular or other configuration with a central axis. Vacuum degassing times can range from 12 hours to over 8 days. A temperature range of about 500 0 F (26O 0 C) to about 900 0 F (482 0 C) is preferred and about 75O 0 F (399 0 C) is more preferred.
  • Dynamic degassing of large amounts of powder are preferred to static degassing. In dynamic degassing, the can is preferably rotated during 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 powder.
  • the vacuum line is crimped and welded shut.
  • the powder is then consolidated further by uniaxially hot pressing the evacuated can (along its central axis while radial movement is restrained) in a die or by hot isostatic pressing (HIP) the can in an isostatic press. At this point the powder charge is nearly 100 percent dense.
  • the billet can be compressed by blind die compaction (step 90) to further densify the structure. Blind die compaction is preferred to further densify the billet as prior consolidation processes may not provide 100% density due to insufficient load available with the press. If uniaxial hot pressing or hot isostatic pressing are used 100% density can be achieved and blind die compaction may not be required. Following densification, the can may be removed by machining.
  • the billet is machined into an extrusion billet, copper jacketed and extruded (step 100).
  • the extrusion process preferably increases the hardness and improves the tensile ductility.
  • Extrusion imparts directional mechanical properties to the material. Forging and/or rolling (step 110) can improve the uniformity of the short transverse mechanical properties.
  • FIG. 11 shows a 3-inch (7.62 cm) diameter copper jacketed Ll 2 aluminum alloy billet ready for extrusion.
  • FIG. 12 is a photo of three 3-inch (7.62 cm) diameter extrusion dies. Representative extrusions using the 3-inch (7.62 cm) diameter dies are shown in FIG. 13. A 12-inch (30.48 cm) ruler is included in the photo for size comparison. Larger 6-inch (15.24 cm) diameter billets were also extruded. Machined 6- inch (15.24 cm) diameter Ll 2 aluminum alloy extrusion billets are shown in FIG. 14.
  • FIG. 15 is a photo of a machined three-piece copper jacketed 6-inch (15.24 cm) diameter billet assembly. A 12-inch (30.48 cm) ruler is included in the photo for size comparison.
  • the upright cylinder behind the three-piece assembly is another machined, copper jacketed Ll 2 aluminum alloy extrusion billet.
  • the top rod is 46 inches (116.8 cm) long.
  • Representative processing parameters for 3-inch diameter Ll 2 aluminum alloy billets are listed in Table 1.
  • Table 1 shows processing details of degassing, vacuum hot pressing (VHP), and extrusion parameters used for fabrication of this material.
  • Degassing and consolidation of billets were performed in the range of 55O 0 F (288 0 C) to 700 0 F (371 0 C).
  • Extrusion was performed in the 35O 0 F (177 0 C) to 65O 0 F (343 0 C) temperature range where billet, die and liner temperatures were maintained equal.
  • Extrusion speed was maintained at 0.5 inches per minute (1.27 cm per minute). Lower speed is desired for higher strength due to less adiabatic heat generation during extrusion.
  • Table 2 includes tensile properties of extrusions that resulted from the above processing parmaters.
  • the measured tensile strength ranges from 101 ksi (696 MPa) to 120 ksi (827 MPa) and yield strength ranges from 97 ksi (667 MPa) to 108 ksi (745 MPa). These strength values are significantly higher than commercially available existing aluminum alloys. It should be noted that higher strength was obtained for lower extrusion temperature conditions.
  • billet number 3 showed a yield strength of 97 ksi (669 MPa) and tensile strength of 101 ksi (696 MPa) for an extrusion temperature of 65O 0 F (343 0 C), whereas billet number 6 showed a yield strength of 108 ksi (745 MPa) and a tensile strength of 120 ksi (827 MPa) for an extrusion temperature of 35O 0 F (177 0 C).
  • Billet number 2 showed a yield strength of 112 ksi (772 MPa) and a tensile strength of 118 (813 MPa) for an extrusion temperature of 400 0 F (204 0 C) which is very close to the strength obtained for an extrusion temperature of 35O 0 F (177 0 C).
  • Ductility which is measured by elongation and reduction in area also showed variations with extrusion temperature. Higher extrusion temperatures resulted in higher ductility. Higher strength values obtained for these extrusions made from cryomilled billets suggest that cryomilling has worked very effectively for Ll 2 strengthened aluminum alloys.

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Abstract

Le procédé et l'appareil ci-décrits produisent des alliages d'aluminium haute résistance à partir d'une poudre cryobroyée contenant des dispersoïdes intermétalliques L12. La poudre cryobroyée est dégazée, scellée sous vide dans un récipient, chauffée, consolidée par pressage à chaud sous vide, et extrudée.
EP10749389.2A 2009-03-05 2010-03-05 Alliages d'aluminium l1 2 haute résistance produits par cryobroyage Withdrawn EP2403967A4 (fr)

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US12/398,712 US20100226817A1 (en) 2009-03-05 2009-03-05 High strength l12 aluminum alloys produced by cryomilling
PCT/US2010/026364 WO2010102206A2 (fr) 2009-03-05 2010-03-05 Alliages d'aluminium l12 haute résistance produits par cryobroyage

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Families Citing this family (7)

* Cited by examiner, † Cited by third party
Publication number Priority date Publication date Assignee Title
US20110064599A1 (en) * 2009-09-15 2011-03-17 United Technologies Corporation Direct extrusion of shapes with l12 aluminum alloys
US20180237893A1 (en) * 2017-02-22 2018-08-23 Orlando RIOS Rapidly solidified aluminum-rare earth element alloy and method of making the same
US11761061B2 (en) 2017-09-15 2023-09-19 Ut-Battelle, Llc Aluminum alloys with improved intergranular corrosion resistance properties and methods of making and using the same
JP7461292B2 (ja) * 2017-11-28 2024-04-03 クエステック イノベーションズ リミテッド ライアビリティ カンパニー 付加製造等の用途のための多成分アルミニウム合金
US11986904B2 (en) 2019-10-30 2024-05-21 Ut-Battelle, Llc Aluminum-cerium-nickel alloys for additive manufacturing
US11608546B2 (en) 2020-01-10 2023-03-21 Ut-Battelle Llc Aluminum-cerium-manganese alloy embodiments for metal additive manufacturing
US11781201B2 (en) * 2022-01-14 2023-10-10 Uif (University Industry Foundation), Yonsei University Aluminium alloy material and method of manufacturing the same

Family Cites Families (93)

* Cited by examiner, † Cited by third party
Publication number Priority date Publication date Assignee Title
US2006127A (en) * 1935-01-16 1935-06-25 Victor B Buck Navigation instrument
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
FR2529909B1 (fr) * 1982-07-06 1986-12-12 Centre Nat Rech Scient Alliages amorphes ou microcristallins a base d'aluminium
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
CH673240A5 (fr) * 1986-08-12 1990-02-28 Bbc Brown Boveri & Cie
JPS6447831A (en) * 1987-08-12 1989-02-22 Takeshi Masumoto High strength and heat resistant aluminum-based alloy and its production
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
AU620155B2 (en) * 1988-10-15 1992-02-13 Koji Hashimoto 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
US5059390A (en) * 1989-06-14 1991-10-22 Aluminum Company Of America Dual-phase, magnesium-based alloy having improved properties
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
US5039476A (en) * 1989-07-28 1991-08-13 Ube Industries, Ltd. Method for production of powder metallurgy alloy
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
JP2724762B2 (ja) * 1989-12-29 1998-03-09 本田技研工業株式会社 高強度アルミニウム基非晶質合金
US5211910A (en) * 1990-01-26 1993-05-18 Martin Marietta Corporation Ultra high strength aluminum-base alloys
JP2619118B2 (ja) * 1990-06-08 1997-06-11 健 増本 粒子分散型高強度非晶質アルミニウム合金
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
JP2864287B2 (ja) * 1990-10-16 1999-03-03 本田技研工業株式会社 高強度高靭性アルミニウム合金の製造方法および合金素材
JPH04218637A (ja) * 1990-12-18 1992-08-10 Honda Motor Co Ltd 高強度高靱性アルミニウム合金の製造方法
US5198045A (en) * 1991-05-14 1993-03-30 Reynolds Metals Company Low density high strength al-li alloy
JP2911673B2 (ja) * 1992-03-18 1999-06-23 健 増本 高強度アルミニウム合金
JPH0673479A (ja) * 1992-05-06 1994-03-15 Honda Motor Co Ltd 高強度高靱性Ａｌ合金
CA2107421A1 (fr) * 1992-10-16 1994-04-17 Steven Alfred Miller Methode de pulverisation a faible pression de gaz
JPH07179974A (ja) * 1993-12-24 1995-07-18 Takeshi Masumoto アルミニウム合金およびその製造方法
US5597529A (en) * 1994-05-25 1997-01-28 Ashurst Technology Corporation (Ireland Limited) Aluminum-scandium alloys
US5858131A (en) * 1994-11-02 1999-01-12 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
JP4080013B2 (ja) * 1996-09-09 2008-04-23 住友電気工業株式会社 高強度高靱性アルミニウム合金およびその製造方法
US5882449A (en) * 1997-07-11 1999-03-16 Mcdonnell Douglas Corporation Process for preparing aluminum/lithium/scandium rolled sheet products
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
US6071324A (en) * 1998-05-28 2000-06-06 Sulzer Metco (Us) Inc. Powder of chromium carbide and nickel chromium
AT407532B (de) * 1998-07-29 2001-04-25 Miba Gleitlager Ag Verbundwerkstoff aus zumindest zwei schichten
AT407404B (de) * 1998-07-29 2001-03-26 Miba Gleitlager Ag Zwischenschicht, insbesondere bindungsschicht, aus einer legierung auf aluminiumbasis
DE19838018C2 (de) * 1998-08-21 2002-07-25 Eads Deutschland Gmbh Geschweißtes Bauteil aus einer schweißbaren, korrosionsbeständigen hochmagnesiumhaltigen Aluminium-Magnesium-Legierung
DE19838015C2 (de) * 1998-08-21 2002-10-17 Eads Deutschland Gmbh Gewalztes, stranggepreßtes, geschweißtes oder geschmiedetes Bauteil aus einer schweißbaren, korrosionsbeständigen hochmagnesiumhaltigen Aluminium-Magnesium-Legierung
DE19838017C2 (de) * 1998-08-21 2003-06-18 Eads Deutschland Gmbh Schweißbare, korrosionsbeständige AIMg-Legierungen, insbesondere für die Verkehrstechnik
US6309594B1 (en) * 1999-06-24 2001-10-30 Ceracon, Inc. Metal consolidation process employing microwave heated pressure transmitting particulate
JP4080111B2 (ja) * 1999-07-26 2008-04-23 ヤマハ発動機株式会社 鍛造用アルミニウム合金製ビレットの製造方法
US6139653A (en) * 1999-08-12 2000-10-31 Kaiser Aluminum & Chemical Corporation Aluminum-magnesium-scandium alloys with zinc and copper
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
AU2001264646A1 (en) * 2000-05-18 2001-11-26 Smith And Wesson Corp. 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
US6524410B1 (en) * 2001-08-10 2003-02-25 Tri-Kor Alloys, Llc Method for producing high strength aluminum alloy welded structures
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
US6880871B2 (en) * 2002-09-05 2005-04-19 Newfrey Llc 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
JP3929978B2 (ja) * 2003-01-15 2007-06-13 ユナイテッド テクノロジーズ コーポレイション アルミニウム基合金
US7648593B2 (en) * 2003-01-15 2010-01-19 United Technologies Corporation Aluminum based alloy
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
US6866817B2 (en) * 2003-07-14 2005-03-15 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
US7547366B2 (en) * 2004-07-15 2009-06-16 Alcoa Inc. 2000 Series alloys with enhanced damage tolerance performance for aerospace applications
US7875132B2 (en) * 2005-05-31 2011-01-25 United Technologies Corporation High temperature aluminum alloys
JP5079225B2 (ja) * 2005-08-25 2012-11-21 富士重工業株式会社 マグネシウムシリサイド粒を分散した状態で含むマグネシウム系金属粒子からなる金属粉末を製造する方法
US7584778B2 (en) * 2005-09-21 2009-09-08 United Technologies Corporation 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 (1)

* Cited by examiner, † Cited by third party
Title
See references of WO2010102206A2 *

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