US20130240095A1 - Heat treatable l12 aluminum alloys - Google Patents

Heat treatable l12 aluminum alloys Download PDF

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US20130240095A1
US20130240095A1 US13/873,678 US201313873678A US2013240095A1 US 20130240095 A1 US20130240095 A1 US 20130240095A1 US 201313873678 A US201313873678 A US 201313873678A US 2013240095 A1 US2013240095 A1 US 2013240095A1
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alloys
aluminum
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silicon
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Awadh B. Pandey
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Raytheon Technologies Corp
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    • CCHEMISTRY; METALLURGY
    • C22METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22CALLOYS
    • C22C21/00Alloys based on aluminium
    • C22C21/02Alloys based on aluminium with silicon as the next major constituent
    • C22C21/04Modified aluminium-silicon alloys
    • CCHEMISTRY; METALLURGY
    • C22METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22CALLOYS
    • C22C1/00Making non-ferrous alloys
    • C22C1/02Making non-ferrous alloys by melting
    • C22C1/026Alloys based on aluminium
    • BPERFORMING OPERATIONS; TRANSPORTING
    • B21MECHANICAL METAL-WORKING WITHOUT ESSENTIALLY REMOVING MATERIAL; PUNCHING METAL
    • B21BROLLING OF METAL
    • B21B1/00Metal-rolling methods or mills for making semi-finished products of solid or profiled cross-section; Sequence of operations in milling trains; Layout of rolling-mill plant, e.g. grouping of stands; Succession of passes or of sectional pass alternations
    • B21B1/46Metal-rolling methods or mills for making semi-finished products of solid or profiled cross-section; Sequence of operations in milling trains; Layout of rolling-mill plant, e.g. grouping of stands; Succession of passes or of sectional pass alternations for rolling metal immediately subsequent to continuous casting
    • BPERFORMING OPERATIONS; TRANSPORTING
    • B21MECHANICAL METAL-WORKING WITHOUT ESSENTIALLY REMOVING MATERIAL; PUNCHING METAL
    • B21CMANUFACTURE OF METAL SHEETS, WIRE, RODS, TUBES OR PROFILES, OTHERWISE THAN BY ROLLING; AUXILIARY OPERATIONS USED IN CONNECTION WITH METAL-WORKING WITHOUT ESSENTIALLY REMOVING MATERIAL
    • B21C23/00Extruding metal; Impact extrusion
    • B21C23/002Extruding materials of special alloys so far as the composition of the alloy requires or permits special extruding methods of sequences
    • BPERFORMING OPERATIONS; TRANSPORTING
    • B21MECHANICAL METAL-WORKING WITHOUT ESSENTIALLY REMOVING MATERIAL; PUNCHING METAL
    • B21JFORGING; HAMMERING; PRESSING METAL; RIVETING; FORGE FURNACES
    • B21J5/00Methods for forging, hammering, or pressing; Special equipment or accessories therefor
    • CCHEMISTRY; METALLURGY
    • C22METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22CALLOYS
    • C22C21/00Alloys based on aluminium
    • C22C21/02Alloys based on aluminium with silicon as the next major constituent
    • CCHEMISTRY; METALLURGY
    • C22METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22FCHANGING THE PHYSICAL STRUCTURE OF NON-FERROUS METALS AND NON-FERROUS ALLOYS
    • C22F1/00Changing the physical structure of non-ferrous metals or alloys by heat treatment or by hot or cold working
    • C22F1/04Changing the physical structure of non-ferrous metals or alloys by heat treatment or by hot or cold working of aluminium or alloys based thereon
    • C22F1/043Changing the physical structure of non-ferrous metals or alloys by heat treatment or by hot or cold working of aluminium or alloys based thereon of alloys with silicon as the next major constituent

Definitions

  • PA0006926U-U73.12-332KL HIGH STRENGTH ALUMINUM ALLOYS WITH L1 2 PRECIPITATES, Ser. No. 12/148,426, Attorney Docket No. PA0006924U-U73.12-334KL; HIGH STRENGTH L1 2 ALUMINUM ALLOYS, Ser. No. 12/148,459, Attorney Docket No. PA0006923U-U73.12-335KL; and L1 2 STRENGTHENED AMORPHOUS ALUMINUM ALLOYS, Ser. No. 12/148,458, Attorney Docket No. PA0001359U-U73.12-336KL.
  • the present invention relates generally to aluminum alloys and more specifically to heat treatable aluminum alloys produced by melt processing and strengthened by L1 2 phase dispersions.
  • 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. Pat. No. 6,248,453 discloses aluminum alloys strengthened by dispersed Al 3 X L1 2 intermetallic phases where X is selected from the group consisting of Sc, Er, Lu, Yb, Tm, and U.
  • 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 L1 2 aluminum alloys are stable up to 572° F. (300° C.).
  • the alloys need to be manufactured by expensive rapid solidification processes with cooling rates in excess of 1.8 ⁇ 10 3 F/sec (10 3 C/sec).
  • U.S. Patent Application Publication No. 2006/0269437 discloses an aluminum alloy that contains scandium and other elements. While the alloy is effective at high temperatures, it is not capable of being heat treated using a conventional age hardening mechanism.
  • the present invention is heat treatable aluminum alloys that can be cast, wrought, or formed by rapid solidification, and thereafter heat treated.
  • the alloys can achieve high temperature performance and can be used at temperatures up to about 650° F. (343° C.).
  • These alloys comprise silicon, and an Al 3 X L1 2 dispersoid 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.
  • X is at least one first element selected from scandium, erbium, thulium, ytterbium, and lutetium
  • X is at least one second element selected from gadolinium, yttrium, zirconium, titanium, hafnium, and niobium.
  • the balance is substantially aluminum.
  • the alloys may also contain magnesium and, optionally, copper, and have less than 1.0 weight percent total impurities.
  • the alloys are formed by a process selected from casting, deformation processing and rapid solidification.
  • the alloys are then heat treated at a temperature of from about 800° F. (426° C.) to about 1100° F. (593° C.) for between about 30 minutes and four hours, followed by quenching in water, and thereafter aged at a temperature from about 200° F. (93° C.) to about 600° F. (315° C.) for about two to forty eight hours.
  • FIG. 1 is an aluminum silicon phase diagram.
  • FIG. 2 is an aluminum scandium phase diagram.
  • FIG. 3 is an aluminum erbium phase diagram.
  • FIG. 4 is an aluminum thulium phase diagram.
  • FIG. 5 is an aluminum ytterbium phase diagram.
  • FIG. 6 is an aluminum lutetium phase diagram.
  • the alloys of this invention are based on the aluminum silicon system.
  • the aluminum silicon phase diagram is shown in FIG. 1 .
  • the binary 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.
  • Hypoeutectic alloys with less than 12.6 weight percent silicon solidify with a microstructure consisting of primary aluminum grains in a finely divided aluminum/silicon eutectic matrix phase.
  • Hypereutectic alloys with silicon contents greater than the eutectic composition solidify with a microstructure of primary silicon grains in a finely divided aluminum/silicon eutectic matrix phase. Alloys of this invention include alloys with the addition of about 4 to about 25 weight percent silicon, more preferably about 4 to about 18 weight percent silicon, and even more preferably about 5 to about 11 weight percent silicon.
  • Aluminum copper and aluminum magnesium alloys are heat treatable with Al 2 Cu ( ⁇ ′), Al 2 CuMg (S′) and Si crystals precipitating in aluminum-copper-silicon alloys; Mg 2 Si and Si crystals precipitating in aluminum-magnesium-silicon alloys following a solution heat treatment, quench, and age process.
  • strengthening phases are Al 2 Cu ( ⁇ ′), Al 2 CuMg (S′), Mg 2 Si and Si crystals following a solution heat treatment, quench, and age process.
  • Mg 2 Al 3 ( ⁇ ) phase precipitates as large intermetallic particle in high magnesium containing aluminum alloys which is not desired from strengthening point of view.
  • the presence of L1 2 phase prevents formation of ⁇ phase in material which improves ductility and toughness of material.
  • the alloys of this invention contain phases consisting of aluminum copper solid solutions, aluminum magnesium solid solutions, and aluminum copper magnesium solid solutions.
  • the solid solutions are dispersions of Al 3 X having an L1 2 structure where X is at least one first element selected from scandium, erbium, thulium, ytterbium, and lutetium. Also present is at least one second element selected from gadolinium, yttrium, zirconium, titanium, hafnium, and niobium.
  • Exemplary aluminum alloys of this invention include, but are not limited to (in weight percent):
  • alloys with the addition of about 0.2 to about 3 weight percent Mg are alloys with the addition of about 0.2 to about 3 weight percent Mg, more preferably alloys with the addition of about 0.3 to about 1.5 weight percent Mg, and even more preferably alloys with the addition of about 0.5 to about 1.5 weight percent Mg; and alloys with the addition of about 0.5 to about 5 weight percent Cu, more preferably alloys with the addition of about 1 to about 4 weight percent Cu, and even more preferably alloys with the addition of about 2 to about 4 weight percent Cu.
  • 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, ytterbium, lutetium, that have an L1 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.
  • Al 3 Sc dispersoids 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° F. (450° C.).
  • Addition of magnesium in solid solution in aluminum increases the lattice parameter of the aluminum matrix, and decreases the lattice parameter mismatch further increasing the resistance of the Al 3 Sc to coarsening.
  • Addition of copper increases the strength of alloys through precipitation of Al 2 Cu ( ⁇ ′) and Al 2 CuMg (S′) phases.
  • these 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° F. (450° C.).
  • Addition of magnesium in solid solution in aluminum increases the lattice parameter of the aluminum matrix, and decreases the lattice parameter mismatch further increasing the resistance of the Al 3 Er to coarsening.
  • 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.
  • 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° F. (450° C.).
  • Addition of magnesium in solid solution in aluminum increases the lattice parameter of the aluminum matrix and decreases the lattice parameter mismatch further increasing the resistance to coarsening of the dispersoid.
  • 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.
  • 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.
  • This low interfacial energy makes the Al 3 Yb dispersoids thermally stable and resistant to coarsening up to temperatures as high as about 842° F. (450° C.).
  • Addition of magnesium in solid solution in aluminum increases the lattice parameter of the aluminum matrix and decreases the lattice parameter mismatch further increasing the resistance to coarsening of the Al 3 Yb.
  • 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° F. (450° C.).
  • Addition of magnesium in solid solution in aluminum increases the lattice parameter of the aluminum matrix and decreases the lattice parameter mismatch further increasing the resistance to coarsening of Al 3 Lu.
  • 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.
  • Gadolinium forms metastable Al 3 Gd dispersoids in the aluminum matrix that have an L1 2 structure in the metastable condition.
  • the Al 3 Gd dispersoids are stable up to temperatures as high as about 842° F. (450° C.) due to their low diffusivity in aluminum.
  • the Al 3 Gd dispersoids have a D0 19 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 L1 2 phase which results in improved thermal and structural stability.
  • Yttrium forms metastable Al 3 Y dispersoids in the aluminum matrix that have an L1 2 structure in the metastable condition and a D0 19 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 L1 2 dispersoids which results in improved thermal and structural stability.
  • Zirconium forms Al 3 Zr dispersoids in the aluminum matrix that have an L1 2 structure in the metastable condition and D0 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 L1 2 structure in the metastable condition and D0 22 structure in the equilibrium condition.
  • the metastable Al 3 Ti dispersoids 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 results in improved thermal and structural stability.
  • Hafnium forms metastable Al 3 Hf dispersoids in the aluminum matrix that have an L1 2 structure in the metastable condition and a D0 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 L1 2 structure in the metastable condition and a D0 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 L1 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 anti-phase boundary on the glide plane. The energy to create the anti-phase boundary is the origin of the strengthening.
  • the cubic L1 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.
  • the amount of scandium present in the alloys of this invention if any may vary from about 0.1 to about 0.5 weight percent, more preferably from about 0.1 to about 0.35 weight percent, and even more preferably from about 0.1 to about 0.25 weight percent.
  • the Al—Sc phase diagram shown in FIG. 2 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 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 L1 2 intermetallic Al 3 Sc 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 10 3 ° C./second.
  • RSP rapid solidification processing
  • the amount of erbium present in the alloys of this invention may vary from about 0.1 to about 6.0 weight percent, more preferably from about 0.1 to about 4 weight percent, and even more preferably from about 0.2 to 2 weight percent.
  • the Al—Er phase diagram shown in FIG. 3 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 L1 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 3 ° C./second. Alloys with erbium in excess of the eutectic composition (hypereutectic alloys) cooled normally will have a microstructure consisting of relatively large Al 3 Er dispersoid in a finely divided aluminum-Al 3 Er eutectic phase matrix.
  • the amount of thulium present in the alloys of this invention may vary from about 0.1 to about 10 weight percent, more preferably from about 0.2 to about 6 weight percent, and even more preferably from about 0.2 to about 4 weight percent.
  • the Al—Tm phase diagram shown in FIG. 4 indicates a eutectic reaction at about 10 weight percent thulium at about 1193° F. (645° C.).
  • Thulium forms metastable Al 3 Tm dispersoids in the aluminum matrix that have an L1 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 L1 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 3 ° C./second.
  • RSP rapid solidification processing
  • the amount of ytterbium present in the alloys of this invention may vary from about 0.1 to about 15 weight percent more preferably from about 0.2 to about 8 weight percent, and even more preferably from about 0.2 to about 4 weight percent.
  • the Al—Yb phase diagram shown in FIG. 5 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 L1 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 3 ° C./second.
  • RSP rapid solidification processing
  • the amount of lutetium present in the alloys of this invention may vary from about 0.1 to about 12 weight percent, more preferably from 0.2 to about 8 weight percent, and even more preferably from about 0.2 to about 4 weight percent.
  • the Al—Lu phase diagram shown in FIG. 6 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 L1 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 3 ° C./second.
  • RSP rapid solidification processing
  • the amount of gadolinium present in the alloys of this invention may vary from about 0.1 to about 4 weight percent, more preferably from 0.2 to about 2 weight percent, and even more preferably from about 0.5 to about 2 weight percent.
  • the amount of yttrium present in the alloys of this invention may vary from about 0.1 to about 4 weight percent, more preferably from 0.2 to about 2 weight percent, and even more preferably from about 0.5 to about 2 weight percent.
  • the amount of zirconium present in the alloys of this invention may vary from about 0.05 to about 1 weight percent, more preferably from 0.1 to about 0.75 weight percent, and even more preferably from about 0.1 to about 0.5 weight percent.
  • the amount of titanium present in the alloys of this invention may vary from about 0.05 to 2 about weight percent, more preferably from 0.1 to about 1 weight percent, and even more preferably from about 0.1 to about 0.5 weight percent.
  • the amount of hafnium present in the alloys of this invention may vary from about 0.05 to about 2 weight percent, more preferably from 0.1 to about 1 weight percent, and even more preferably from about 0.1 to about 0.5 weight percent.
  • the amount of niobium present in the alloys of this invention may vary from about 0.05 to about 1 weight percent, more preferably from 0.1 to about 0.75 weight percent, and even more preferably from about 0.1 to about 0.5 weight percent.
  • inventive alloys include at least one of about 0.001 weight percent to about 0.10 weight percent sodium, about 0.001 weight percent to about 0.10 weight percent calcium, about 0.001 to about 0.10 weight percent strontium, about 0.001 to about 0.10 weight percent antimony, 0.001 to 0.10 weight percent barium and about 0.001 to about 0.10 weight percent phosphorus. These are added to refine the microstructure of the eutectic phase and the primary silicon particle morphology and size.
  • These aluminum alloys may be made by any and all consolidation and fabrication processes known to those in the art such as casting (without further deformation), deformation processing (wrought processing), rapid solidification processing, forging, extrusion, rolling, die forging, powder metallurgy and others.
  • the rapid solidification process should have a cooling rate greater that about 10 3 ° C./second including but not limited to powder processing, atomization, melt spinning, splat quenching, spray deposition, cold spray, plasma spray, laser melting and deposition, ball milling and cryomilling.
  • Preferred exemplary aluminum alloys of this invention include, but are not limited to (in weight percent):
  • alloys with the addition of about 0.2-3 weight percent Mg more preferably alloys with the addition of about 0.3-1.5 weight percent Mg, and even more preferably alloys with the addition of about 0.5-1.5 weight percent Mg; and alloys with the addition of about 0.5-5 weight percent Cu, more preferably alloys with the addition of about 1-4 weight percent Cu, and even more preferably alloys with the addition of about 2-4 weight percent Cu.
  • alloys with the addition of about 0.2-3 weight percent Mg more preferably alloys with the addition of about 0.3-1.5 weight percent Mg, and even more preferably alloys with the addition of about 0.5-1.5 weight percent Mg; and alloys with the addition of about 0.5-5 weight percent Cu, more preferably alloys with the addition of about 1-4 weight percent Cu, and even more preferably alloys with the addition of about 2-4 weight percent Cu.

Abstract

High temperature heat treatable aluminum alloys that can be used at temperatures from about −420° F. (−251° C.) up to about 650° F. (343° C.) are described. The alloys are strengthened by dispersion of particles based on the L12 intermetallic compound Al3X. These alloys comprise aluminum; silicon; at least one of scandium, erbium, thulium, ytterbium, and lutetium; and at least one of gadolinium, yttrium, zirconium, titanium, hafnium, and niobium. Magnesium and copper are optional alloying elements.

Description

    CROSS-REFERENCE TO RELATED APPLICATIONS
  • This application a divisional of U.S. patent application Ser. No. 12/148,383, filed Apr. 18, 2008, which is related to the following co-pending applications that were filed on even date therewith and were assigned to the same assignee: L12 ALUMINUM ALLOYS WITH BIMODAL AND TRIMODAL DISTRIBUTION, Ser. No. 12/148,395, Attorney Docket No. PA0006933U-U73.12-325KL; DISPERSION STRENGTHENED L12 ALUMINUM ALLOYS, Ser. No. 12/148,432, Attorney Docket No. PA0006932U-U73.12-326KL; HIGH STRENGTH L12 ALUMINUM ALLOYS, Ser. No. 12/148,394, Attorney Docket No. PA0006929U-U73.12-329KL; HIGH STRENGTH L12 ALUMINUM ALLOYS, Ser. No. 12/148,382, Attorney Docket No. PA0006928U-U73.12-330KL; HEAT TREATABLE L12 ALUMINUM ALLOYS, Ser. No. 12/148,396, Attorney Docket No. PA0006927U-U73.12-331KL; HIGH STRENGTH L12 ALUMINUM ALLOYS, Ser. No. 12/148,387, Attorney Docket No. PA0006926U-U73.12-332KL; HIGH STRENGTH ALUMINUM ALLOYS WITH L12 PRECIPITATES, Ser. No. 12/148,426, Attorney Docket No. PA0006924U-U73.12-334KL; HIGH STRENGTH L12 ALUMINUM ALLOYS, Ser. No. 12/148,459, Attorney Docket No. PA0006923U-U73.12-335KL; and L12 STRENGTHENED AMORPHOUS ALUMINUM ALLOYS, Ser. No. 12/148,458, Attorney Docket No. PA0001359U-U73.12-336KL.
  • BACKGROUND
  • The present invention relates generally to aluminum alloys and more specifically to heat treatable aluminum alloys produced by melt processing and strengthened by L12 phase dispersions.
  • 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 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 U. 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.). In order to create aluminum alloys containing fine dispersions of Al3X L12 particles, the alloys need to be manufactured by expensive rapid solidification processes with cooling rates in excess of 1.8×103 F/sec (103 C/sec). U.S. Patent Application Publication No. 2006/0269437 discloses an aluminum alloy that contains scandium and other elements. While the alloy is effective at high temperatures, it is not capable of being heat treated using a conventional age hardening mechanism.
  • Heat treatable aluminum alloys strengthened by coherent L12 intermetallic phases produced by standard, inexpensive melt processing techniques would be useful.
  • SUMMARY
  • The present invention is heat treatable aluminum alloys that can be cast, wrought, or formed by rapid solidification, and thereafter heat treated. The alloys can achieve high temperature performance and can be used at temperatures up to about 650° F. (343° C.).
  • These alloys comprise silicon, and an Al3X L12 dispersoid 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.
  • The alloys may also contain magnesium and, optionally, copper, and have less than 1.0 weight percent total impurities.
  • The alloys are formed by a process selected from casting, deformation processing and rapid solidification. The alloys are then heat treated at a temperature of from about 800° F. (426° C.) to about 1100° F. (593° C.) for between about 30 minutes and four hours, followed by quenching in water, and thereafter aged at a temperature from about 200° F. (93° C.) to about 600° F. (315° C.) for about two to forty eight hours.
  • BRIEF DESCRIPTION OF THE DRAWINGS
  • FIG. 1 is an aluminum silicon phase diagram.
  • FIG. 2 is an aluminum scandium phase diagram.
  • FIG. 3 is an aluminum erbium phase diagram.
  • FIG. 4 is an aluminum thulium phase diagram.
  • FIG. 5 is an aluminum ytterbium phase diagram.
  • FIG. 6 is an aluminum lutetium phase diagram.
  • DETAILED DESCRIPTION
  • The alloys of this invention are based on the aluminum silicon system. The aluminum silicon phase diagram is shown in FIG. 1. The binary 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. Hypoeutectic alloys with less than 12.6 weight percent silicon solidify with a microstructure consisting of primary aluminum grains in a finely divided aluminum/silicon eutectic matrix phase. Hypereutectic alloys with silicon contents greater than the eutectic composition solidify with a microstructure of primary silicon grains in a finely divided aluminum/silicon eutectic matrix phase. Alloys of this invention include alloys with the addition of about 4 to about 25 weight percent silicon, more preferably about 4 to about 18 weight percent silicon, and even more preferably about 5 to about 11 weight percent silicon.
  • Other alloys of this invention include aluminum silicon containing copper or magnesium, or both copper and magnesium. Copper and magnesium are completely soluble in aluminum in the compositions of the inventive alloys discussed herein. Aluminum copper and aluminum magnesium alloys are heat treatable with Al2Cu (θ′), Al2CuMg (S′) and Si crystals precipitating in aluminum-copper-silicon alloys; Mg2Si and Si crystals precipitating in aluminum-magnesium-silicon alloys following a solution heat treatment, quench, and age process. In aluminum-copper-magnesium-silicon alloys, strengthening phases are Al2Cu (θ′), Al2CuMg (S′), Mg2Si and Si crystals following a solution heat treatment, quench, and age process. Mg2Al3 (β) phase precipitates as large intermetallic particle in high magnesium containing aluminum alloys which is not desired from strengthening point of view. The presence of L12 phase prevents formation of β phase in material which improves ductility and toughness of material. The alloys of this invention contain phases consisting of aluminum copper solid solutions, aluminum magnesium solid solutions, and aluminum copper magnesium solid solutions. In the solid solutions are dispersions of Al3X having an L12 structure where X is at least one first element selected from scandium, erbium, thulium, ytterbium, and lutetium. Also present is at least one second element selected from gadolinium, yttrium, zirconium, titanium, hafnium, and niobium.
  • Exemplary aluminum alloys of this invention include, but are not limited to (in weight percent):
  • about Al-(4-25)Si-(0.1-0.5)Sc-(0.1-4)Gd;
  • about Al-(4-25)Si-(0.1-6)Er-(0.1-4)Gd;
  • about Al-(4-25)Si-(0.1-10)Tm-(0.1-4)Gd;
  • about Al-(4-25)Si-(0.1-15)Yb-(0.1-4)Gd;
  • about Al-(4-25)Si-(0.1-12)Lu-(0.1-4)Gd;
  • about Al-(4-25)Si-(0.1-0.5)Sc-(0.1-4)Y;
  • about Al-(4-25)Si-(0.1-6)Er-(0.1-4)Y;
  • about Al-(4-25)Si-(0.1-10)Tm-(0.1-4)Y;
  • about Al-(4-25)Si-(0.1-15)Yb-(0.1-4)Y;
  • about Al-(4-25)Si-(0.1-12)Lu-(0.1-4)Y;
  • about Al-(4-25)Si-(0.1-0.5)Sc-(0.05-1)Zr;
  • about Al-(4-25)Si-(0.1-6)Er-(0.05-1)Zr;
  • about Al-(4-25)Si-(0.1-10)Tm-(0.05-1)Zr;
  • about Al-(4-25)Si-(0.1-15)Yb-(0.05-1)Zr;
  • about Al-(4-25)Si-(0.1-12)Lu-(0.05-1)Zr;
  • about Al-(4-25)Si-(0.1-0.5)Sc-(0.05-2)Ti;
  • about Al-(4-25)Si-(0.1-6)Er-(0.05-2)Ti;
  • about Al-(4-25)Si-(0.1-10)Tm-(0.05-2)Ti;
  • about Al-(4-25)Si-(0.1-15)Yb-(0.05-2)Ti;
  • about Al-(4-25)Si-(0.1-12)Lu-(0.05-2)Ti;
  • about Al-(4-25)Si-(0.1-0.5)Sc-(0.05-2)Hf;
  • about Al-(4-25)Si-(0.1-6)Er-(0.05-2)Hf;
  • about Al-(4-25)Si-(0.1-10)Tm-(0.05-2)Hf;
  • about Al-(4-25)Si-(0.1-15)Yb-(0.05-2)Hf;
  • about Al-(4-25)Si-(0.1-12)Lu-(0.05-2)Hf;
  • about Al-(4-25)Si-(0.1-0.5)Sc-(0.05-1)Nb;
  • about Al-(4-25)Si-(0.1-6)Er-(0.05-1)Nb;
  • about Al-(4-25)Si-(0.1-10)Tm-(0.05-1)Nb;
  • about Al-(4-25)Si-(0.1-15)Yb-(0.05-1)Nb; and
  • about Al-(4-25)Si-(0.1-12)Lu-(0.05-1)Nb.
  • Examples of similar alloys to these are alloys with the addition of about 0.2 to about 3 weight percent Mg, more preferably alloys with the addition of about 0.3 to about 1.5 weight percent Mg, and even more preferably alloys with the addition of about 0.5 to about 1.5 weight percent Mg; and alloys with the addition of about 0.5 to about 5 weight percent Cu, more preferably alloys with the addition of about 1 to about 4 weight percent Cu, and even more preferably alloys with the addition of about 2 to about 4 weight percent Cu.
  • In the inventive 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, ytterbium, 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.). Addition of magnesium in solid solution in aluminum increases the lattice parameter of the aluminum matrix, and decreases the lattice parameter mismatch further increasing the resistance of the Al3Sc to coarsening. Addition of copper increases the strength of alloys through precipitation of Al2Cu (θ′) and Al2CuMg (S′) phases. 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.). Addition of magnesium in solid solution in aluminum increases the lattice parameter of the aluminum matrix, and decreases the lattice parameter mismatch further increasing the resistance of the Al3Er to coarsening. Addition of copper increases the strength of alloys through precipitation of Al2Cu (θ′) and Al2CuMg (S′) phases. 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.). Addition of magnesium in solid solution in aluminum increases the lattice parameter of the aluminum matrix and decreases the lattice parameter mismatch further increasing the resistance to coarsening of the dispersoid. Addition of copper increases the strength of alloys through precipitation of Al2Cu (θ′) and Al2CuMg (S′) phases. 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.). Addition of magnesium in solid solution in aluminum increases the lattice parameter of the aluminum matrix and decreases the lattice parameter mismatch further increasing the resistance to coarsening of the Al3Yb. Addition of copper increases the strength of alloys through precipitation of Al2Cu (θ′) and Al2CuMg (S′) phases. 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.). Addition of magnesium in solid solution in aluminum increases the lattice parameter of the aluminum matrix and decreases the lattice parameter mismatch further increasing the resistance to coarsening of Al3Lu. Addition of copper increases the strength of alloys through precipitation of Al2Cu (θ′) and Al2CuMg (S′) phases. 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 have an L12 structure in the metastable condition. The Al3Gd dispersoids 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 D022 structure in the equilibrium condition. The metastable Al3Ti dispersoids 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 results 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.
  • The amount of scandium present in the alloys of this invention if any may vary from about 0.1 to about 0.5 weight percent, more preferably from about 0.1 to about 0.35 weight percent, and even more preferably from about 0.1 to about 0.25 weight percent. The Al—Sc phase diagram shown in FIG. 2 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 alloys of this invention, if any, may vary from about 0.1 to about 6.0 weight percent, more preferably from about 0.1 to about 4 weight percent, and even more preferably from about 0.2 to 2 weight percent. The Al—Er phase diagram shown in FIG. 3 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. Alloys with erbium in excess of the eutectic composition (hypereutectic alloys) cooled normally will have a microstructure consisting of relatively large Al3Er dispersoid in a finely divided aluminum-Al3Er eutectic phase matrix.
  • The amount of thulium present in the alloys of this invention, if any, may vary from about 0.1 to about 10 weight percent, more preferably from about 0.2 to about 6 weight percent, and even more preferably from about 0.2 to about 4 weight percent. The Al—Tm phase diagram shown in FIG. 4 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 15 weight percent more preferably from about 0.2 to about 8 weight percent, and even more preferably from about 0.2 to about 4 weight percent. The Al—Yb phase diagram shown in FIG. 5 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 12 weight percent, more preferably from 0.2 to about 8 weight percent, and even more preferably from about 0.2 to about 4 weight percent. The Al—Lu phase diagram shown in FIG. 6 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 4 weight percent, more preferably from 0.2 to about 2 weight percent, and even more preferably from about 0.5 to about 2 weight percent.
  • The amount of yttrium present in the alloys of this invention, if any, may vary from about 0.1 to about 4 weight percent, more preferably from 0.2 to about 2 weight percent, and even more preferably from about 0.5 to about 2 weight percent.
  • The amount of zirconium present in the alloys of this invention, if any, may vary from about 0.05 to about 1 weight percent, more preferably from 0.1 to about 0.75 weight percent, and even more preferably from about 0.1 to about 0.5 weight percent.
  • The amount of titanium present in the alloys of this invention, if any, may vary from about 0.05 to 2 about weight percent, more preferably from 0.1 to about 1 weight percent, and even more preferably from about 0.1 to about 0.5 weight percent.
  • The amount of hafnium present in the alloys of this invention, if any, may vary from about 0.05 to about 2 weight percent, more preferably from 0.1 to about 1 weight percent, and even more preferably from about 0.1 to about 0.5 weight percent.
  • The amount of niobium present in the alloys of this invention, if any, may vary from about 0.05 to about 1 weight percent, more preferably from 0.1 to about 0.75 weight percent, and even more preferably from about 0.1 to about 0.5 weight percent.
  • In order to have the best properties for the alloys 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, 0.1 weight percent cobalt, and 0.1 weight percent nickel. The total quantity of additional elements should not exceed about 1% by weight, including the above listed elements.
  • Other additions in the inventive alloys include at least one of about 0.001 weight percent to about 0.10 weight percent sodium, about 0.001 weight percent to about 0.10 weight percent calcium, about 0.001 to about 0.10 weight percent strontium, about 0.001 to about 0.10 weight percent antimony, 0.001 to 0.10 weight percent barium and about 0.001 to about 0.10 weight percent phosphorus. These are added to refine the microstructure of the eutectic phase and the primary silicon particle morphology and size.
  • These aluminum alloys may be made by any and all consolidation and fabrication processes known to those in the art such as casting (without further deformation), deformation processing (wrought processing), rapid solidification processing, forging, extrusion, rolling, die forging, powder metallurgy and others. The rapid solidification process should have a cooling rate greater that about 103° C./second including but not limited to powder processing, atomization, melt spinning, splat quenching, spray deposition, cold spray, plasma spray, laser melting and deposition, ball milling and cryomilling.
  • Preferred exemplary aluminum alloys of this invention include, but are not limited to (in weight percent):
  • about Al-(4-18)Si-(0.1-0.35)Sc-(0.2-2)Gd;
  • about Al-(4-18)Si-(0.1-4)Er-(0.2-2)Gd;
  • about Al-(4-18)Si-(0.2-6)Tm-(0.2-2)Gd;
  • about Al-(4-18)Si-(0.2-8)Yb-(0.2-2)Gd;
  • about Al-(4-18)Si-(0.2-8)Lu-(0.2-2)Gd;
  • about Al-(4-18)Si-(0.1-0.35)Sc-(0.2-2)Y;
  • about Al-(4-18)Si-(0.1-4)Er-(0.2-2)Y;
  • about Al-(4-18)Si-(0.2-6)Tm-(0.2-2)Y;
  • about Al-(4-18)Si-(0.2-8)Yb-(0.2-2)Y;
  • about Al-(4-18)Si-(0.2-8)Lu-(0.2-2)Y;
  • about Al-(4-18)Si-(0.1-0.35)Sc-(0.1-0.75)Zr;
  • about Al-(4-18)Si-(0.1-4)Er-(0.1-0.75)Zr;
  • about Al-(4-18)Si-(0.2-6)Tm-(0.1-0.75)Zr;
  • about Al-(4-18)Si-(0.2-8)Yb-(0.1-0.75)Zr;
  • about Al-(4-18)Si-(0.2-8)Lu-(0.1-0.75)Zr;
  • about Al-(4-18)Si-(0.1-0.35)Sc-(0.1-1)Ti;
  • about Al-(4-18)Si-(0.1-4)Er-(0.1-1)Ti;
  • about Al-(4-18)Si-(0.2-6)Tm-(0.1-1)Ti;
  • about Al-(4-18)Si-(0.2-8)Yb-(0.1-1)Ti;
  • about Al-(4-18)Si-(0.2-8)Lu-(0.1-1)Ti;
  • about Al-(4-18)Si-(0.1-0.35)Sc-(0.1-1)Hf;
  • about Al-(4-18)Si-(0.1-4)Er-(0.1-1)Hf;
  • about Al-(4-18)Si-(0.2-6)Tm-(0.1-1)Hf;
  • about Al-(4-18)Si-(0.2-8)Yb-(0.1-1)Hf;
  • about Al-(4-18)Si-(0.2-8)Lu-(0.1-1)Hf;
  • about Al-(4-18)Si-(0.1-0.35)Sc-(0.1-0.75)Nb;
  • about Al-(4-18)Si-(0.2-2)Er-(0.1-0.75)Nb;
  • about Al-(4-18)Si-(0.2-6)Tm-(0.1-0.75)Nb;
  • about Al-(4-18)Si-(0.2-8)Yb-(0.1-0.75)Nb; and
  • about Al-(4-18)Si-(0.2-8)Lu-(0.1-0.75)Nb.
  • Examples of similar alloys to these are alloys with the addition of about 0.2-3 weight percent Mg, more preferably alloys with the addition of about 0.3-1.5 weight percent Mg, and even more preferably alloys with the addition of about 0.5-1.5 weight percent Mg; and alloys with the addition of about 0.5-5 weight percent Cu, more preferably alloys with the addition of about 1-4 weight percent Cu, and even more preferably alloys with the addition of about 2-4 weight percent Cu.
  • Even more preferred exemplary aluminum alloys of this invention include, but are not limited to (in weight percent):
  • about Al-(5-11)Si-(0.1-0.25)Sc-(0.5-2)Gd;
  • about Al-(5-11)Si-(0.2-2)Er-(0.5-2)Gd;
  • about Al-(5-11)Si-(0.2-4)Tm-(0.5-2)Gd;
  • about Al-(5-11)Si-(0.2-4)Yb-(0.5-2)Gd;
  • about Al-(5-11)Si-(0.2-4)Lu-(0.5-2)Gd;
  • about Al-(5-11)Si-(0.1-0.25)Sc-(0.5-2)Y;
  • about Al-(5-11)Si-(0.2-2)Er-(0.5-2)Y;
  • about Al-(5-11)Si-(0.2-4)Tm-(0.5-2)Y;
  • about Al-(5-11)Si-(0.2-4)Yb-(0.5-2)Y;
  • about Al-(5-11)Si-(0.2-4)Lu-(0.5-2)Y;
  • about Al-(5-11)Si-(0.1-0.25)Sc-(0.1-0.5)Zr;
  • about Al-(5-11)Si-(0.2-2)Er-(0.1-0.5)Zr;
  • about Al-(5-11)Si-(0.2-4)Tm-(0.1-0.5)Zr;
  • about Al-(5-11)Si-(0.2-4)Yb-(0.1-0.5)Zr;
  • about Al-(5-11)Si-(0.2-4)Lu-(0.1-0.5)Zr;
  • about Al-(5-11)Si-(0.1-0.25)Sc-(0.1-0.5)Ti;
  • about Al-(5-11)Si-(0.2-2)Er-(0.1-0.5)Ti;
  • about Al-(5-11)Si-(0.2-4)Tm-(0.1-0.5)Ti;
  • about Al-(5-11)Si-(0.2-4)Yb-(0.1-0.5)Ti;
  • about Al-(5-11)Si-(0.2-4)Lu-(0.1-0.5)Ti;
  • about Al-(5-11)Si-(0.1-0.25)Sc-(0.1-0.5)Hf;
  • about Al-(5-11)Si-(0.2-2)Er-(0.1-0.5)Hf;
  • about Al-(5-11)Si-(0.2-4)Tm-(0.1-0.5)Hf;
  • about Al-(5-11)Si-(0.2-4)Yb-(0.1-0.5)Hf;
  • about Al-(5-11)Si-(0.2-4)Lu-(0.1-0.5)Hf;
  • about Al-(5-11)Si-(0.1-0.25)Sc-(0.1-0.5)Nb;
  • about Al-(5-11)Si-(0.2-2)Er-(0.1-0.5)Nb;
  • about Al-(5-11)Si-(0.2-4)Tm-(0.1-0.5)Nb;
  • about Al-(5-11)Si-(0.2-4)Yb-(0.1-0.5)Nb; and
  • about Al-(5-11)Si-(0.2-4)Lu-(0.1-0.5)Nb.
  • Examples of similar alloys to these are alloys with the addition of about 0.2-3 weight percent Mg, more preferably alloys with the addition of about 0.3-1.5 weight percent Mg, and even more preferably alloys with the addition of about 0.5-1.5 weight percent Mg; and alloys with the addition of about 0.5-5 weight percent Cu, more preferably alloys with the addition of about 1-4 weight percent Cu, and even more preferably alloys with the addition of about 2-4 weight percent Cu.
  • 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 (5)

1. A method of forming a heat treatable aluminum alloy, the method comprising:
(a) forming a melt comprising:
about 4.0 to about 25.0 weight percent silicon;
about 0.2 to about 3.0 weight percent magnesium;
about 0.5 to about 5.0 weight percent copper;
at least one first element selected from the group comprising about 0.1 to about 0.5 weight percent scandium, about 0.1 to about 6.0 weight percent erbium, about 0.1 to about 10 weight percent thulium, about 0.1 to about 15.0 weight percent ytterbium, and about 0.1 to about 12 weight percent lutetium;
at least one second element selected from the group comprising about 0.1 to about 4.0 weight percent gadolinium, about 0.1 to about 4.0 weight percent yttrium, about 0.05 to about 1.0 weight percent zirconium, about 0.05 to about 2.0 weight percent titanium, about 0.05 to about 2.0 weight percent hafnium, and about 0.05 to about 1.0 weight percent niobium; and
the balance substantially aluminum;
(b) solidifying the melt to form a solid body; and
(c) heat treating the solid body.
2. The method of claim 1 further comprising:
refining the structure of the solid body by deformation processing including at least one of: extrusion, forging and rolling.
3. The method of claim 1, wherein solidifying comprises a casting process.
4. The method of claim 1, wherein solidifying comprises a rapid solidification process in which the cooling rate is greater than about 103° C./second including at least one of: powder processing, atomization, melt spinning, splat quenching, spray deposition, cold spray, plasma spray, laser melting and deposition, ball milling, and cryomilling.
5. The method of claim 1, wherein the heat treating comprises:
solution heat treatment at about 800° F. (426° C.) to about 1100° F. (593° C.) for about thirty minutes to four hours; and
quenching; and
aging at a temperature of about 200° F. (93° C.) to about 600° F. (315° C.) for about two to forty eight hours.
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Cited By (1)

* Cited by examiner, † Cited by third party
Publication number Priority date Publication date Assignee Title
WO2020150056A1 (en) * 2019-01-18 2020-07-23 Divergent Technologies, Inc. Aluminum alloys

Families Citing this family (34)

* Cited by examiner, † Cited by third party
Publication number Priority date Publication date Assignee Title
US9611522B2 (en) * 2009-05-06 2017-04-04 United Technologies Corporation Spray deposition of L12 aluminum alloys
JP5860873B2 (en) * 2010-06-16 2016-02-16 ノルスク・ヒドロ・アーエスアーNorsk Hydro Asa Castable heat resistant aluminum alloy
CN102021444B (en) * 2010-12-09 2012-08-22 北京科技大学 High-conductive heat-resistant aluminium alloy conductor and preparation method thereof
GB201102849D0 (en) * 2011-02-18 2011-04-06 Univ Brunel Method of refining metal alloys
ITTO20110257A1 (en) * 2011-03-24 2012-09-25 Avio Spa METHOD FOR REPAIRING AN ALUMINUM ALLOY COMPONENT
WO2012171552A1 (en) * 2011-06-14 2012-12-20 Federal-Mogul Nürnberg GmbH Piston for an internal combustion engine, method for producing a piston and use of an alloy for casting a piston
CN102296218A (en) * 2011-08-24 2011-12-28 吴江市精工铝字制造厂 High-strength heat-resistant magnalium alloy
US10174409B2 (en) * 2011-10-28 2019-01-08 Alcoa Usa Corp. High performance AlSiMgCu casting alloy
CN103361524B (en) * 2013-07-05 2015-05-20 苏州有色金属研究院有限公司 Composite modification method for hypereutectic aluminum-silicon alloy
WO2015006447A1 (en) 2013-07-10 2015-01-15 Alcoa Inc. Methods for producing forged products and other worked products
CN103627935A (en) * 2013-12-09 2014-03-12 国家电网公司 Non-heat-treated heat-resistant aluminium alloy monofilament and preparation method thereof
MX2016008166A (en) * 2013-12-20 2016-09-29 Alcoa Inc HIGH PERFORMANCE AlSiMgCu CASTING ALLOY.
FR3020291B1 (en) * 2014-04-29 2017-04-21 Saint Jean Ind METHOD FOR MANUFACTURING METAL OR METAL MATRIX COMPOSITE ARTICLES MADE OF ADDITIVE MANUFACTURING FOLLOWED BY A FORGING OPERATION OF SAID PARTS
CN104651673A (en) * 2015-03-09 2015-05-27 苏州圣谱拉新材料科技有限公司 Stretch-proof nickel-aluminum alloy material and preparation method thereof
CN108472712A (en) 2016-01-14 2018-08-31 奥科宁克公司 Method for producing forging product and other converted products
CN106756305B (en) * 2017-01-03 2018-07-13 江苏理工学院 A kind of Aluminum alloy modification processing method
CN106834815B (en) * 2017-02-27 2018-04-10 广东省材料与加工研究所 A kind of aluminium zirconium carbon rare earth fining agent and preparation method thereof
CN106884129A (en) * 2017-03-14 2017-06-23 广州金邦液态模锻技术有限公司 A kind of Technology for Heating Processing for extrusion casint aluminium alloy knuckle
CN107326228B (en) * 2017-06-23 2019-09-10 兰州理工大学 A kind of composite inoculating transcocrystallized Al-Si alloy and preparation method thereof
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Citations (6)

* Cited by examiner, † Cited by third party
Publication number Priority date Publication date Assignee Title
US5039476A (en) * 1989-07-28 1991-08-13 Ube Industries, Ltd. Method for production of powder metallurgy alloy
US6517954B1 (en) * 1998-07-29 2003-02-11 Miba Gleitlager Aktiengesellschaft Aluminium alloy, notably for a layer
US20030192627A1 (en) * 2002-04-10 2003-10-16 Lee Jonathan A. High strength aluminum alloy for high temperature applications
US7883590B1 (en) * 2008-04-18 2011-02-08 United Technologies Corporation Heat treatable L12 aluminum alloys
US7909947B2 (en) * 2008-04-18 2011-03-22 United Technologies Corporation High strength L12 aluminum alloys
US8409497B2 (en) * 2009-10-16 2013-04-02 United Technologies Corporation Hot and cold rolling high strength L12 aluminum alloys

Family Cites Families (91)

* 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
FR2529909B1 (en) * 1982-07-06 1986-12-12 Centre Nat Rech Scient AMORPHOUS OR MICROCRYSTALLINE ALLOYS BASED ON ALUMINUM
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
FR2584095A1 (en) * 1985-06-28 1987-01-02 Cegedur AL ALLOYS WITH HIGH LI AND SI CONTENT AND METHOD OF MANUFACTURE
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
US4689090A (en) * 1986-03-20 1987-08-25 Aluminum Company Of America Superplastic aluminum alloys containing scandium
US4874440A (en) * 1986-03-20 1989-10-17 Aluminum Company Of America Superplastic aluminum products and alloys
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
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
US4927470A (en) * 1988-10-12 1990-05-22 Aluminum Company Of America Thin gauge aluminum plate product by isothermal treatment and ramp anneal
US4946517A (en) * 1988-10-12 1990-08-07 Aluminum Company Of America Unrecrystallized aluminum plate product by ramp annealing
AU620155B2 (en) * 1988-10-15 1992-02-13 Koji Hashimoto Amorphous aluminum alloys
US4933140A (en) * 1988-11-17 1990-06-12 Ceracon, Inc. Electrical heating of graphite grain employed in consolidation of objects
US4853178A (en) * 1988-11-17 1989-08-01 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
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 (en) * 1989-12-29 1998-03-09 本田技研工業株式会社 High-strength aluminum-based amorphous alloy
US5211910A (en) * 1990-01-26 1993-05-18 Martin Marietta Corporation Ultra high strength aluminum-base alloys
JP2619118B2 (en) * 1990-06-08 1997-06-11 健 増本 Particle-dispersed high-strength amorphous aluminum alloy
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 (en) * 1990-10-16 1999-03-03 本田技研工業株式会社 Method for producing high strength and high toughness aluminum alloy and alloy material
JPH04218637A (en) * 1990-12-18 1992-08-10 Honda Motor Co Ltd Manufacture of 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
JP2911673B2 (en) * 1992-03-18 1999-06-23 健 増本 High strength aluminum alloy
JPH0673479A (en) * 1992-05-06 1994-03-15 Honda Motor Co Ltd High strength and high toughness al alloy
CA2107421A1 (en) * 1992-10-16 1994-04-17 Steven Alfred Miller Atomization with low atomizing gas pressure
US5597529A (en) * 1994-05-25 1997-01-28 Ashurst Technology Corporation (Ireland Limited) Aluminum-scandium alloys
CA2190951A1 (en) * 1994-05-25 1995-11-30 William Troy Tack Aluminum-scandium alloys and uses thereof
WO1996010099A1 (en) * 1994-09-26 1996-04-04 Ashurst Technology Corporation (Ireland) Limited High strength aluminum casting alloys for structural applications
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 (en) * 1996-09-09 2008-04-23 住友電気工業株式会社 High strength and high toughness aluminum alloy and method for producing the same
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
JP3592052B2 (en) * 1997-12-01 2004-11-24 株式会社神戸製鋼所 Filler for welding aluminum alloy and method for welding aluminum alloy using the same
US6071324A (en) * 1998-05-28 2000-06-06 Sulzer Metco (Us) Inc. Powder of chromium carbide and nickel chromium
AT407404B (en) * 1998-07-29 2001-03-26 Miba Gleitlager Ag INTERMEDIATE LAYER, IN PARTICULAR BOND LAYER, FROM AN ALUMINUM-BASED ALLOY
DE19838015C2 (en) * 1998-08-21 2002-10-17 Eads Deutschland Gmbh Rolled, extruded, welded or forged component made of a weldable, corrosion-resistant, high-magnesium aluminum-magnesium alloy
DE19838018C2 (en) * 1998-08-21 2002-07-25 Eads Deutschland Gmbh Welded component made of a weldable, corrosion-resistant, high-magnesium aluminum-magnesium alloy
DE19838017C2 (en) * 1998-08-21 2003-06-18 Eads Deutschland Gmbh Weldable, corrosion resistant AIMg alloys, especially for traffic engineering
US6309594B1 (en) * 1999-06-24 2001-10-30 Ceracon, Inc. Metal consolidation process employing microwave heated pressure transmitting particulate
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
US6557289B2 (en) * 2000-05-18 2003-05-06 Smith & 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
EP1499753A2 (en) * 2002-04-24 2005-01-26 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 (en) * 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 Aluminum base alloy
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
DE10352932B4 (en) * 2003-11-11 2007-05-24 Eads Deutschland Gmbh Cast aluminum alloy
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
US7393559B2 (en) * 2005-02-01 2008-07-01 The Regents Of The University Of California Methods for production of FGM net shaped body for various applications
US7875132B2 (en) * 2005-05-31 2011-01-25 United Technologies Corporation High temperature aluminum alloys
JP5079225B2 (en) * 2005-08-25 2012-11-21 富士重工業株式会社 Method for producing metal powder comprising magnesium-based metal particles containing dispersed magnesium silicide grains
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

Patent Citations (6)

* Cited by examiner, † Cited by third party
Publication number Priority date Publication date Assignee Title
US5039476A (en) * 1989-07-28 1991-08-13 Ube Industries, Ltd. Method for production of powder metallurgy alloy
US6517954B1 (en) * 1998-07-29 2003-02-11 Miba Gleitlager Aktiengesellschaft Aluminium alloy, notably for a layer
US20030192627A1 (en) * 2002-04-10 2003-10-16 Lee Jonathan A. High strength aluminum alloy for high temperature applications
US7883590B1 (en) * 2008-04-18 2011-02-08 United Technologies Corporation Heat treatable L12 aluminum alloys
US7909947B2 (en) * 2008-04-18 2011-03-22 United Technologies Corporation High strength L12 aluminum alloys
US8409497B2 (en) * 2009-10-16 2013-04-02 United Technologies Corporation Hot and cold rolling high strength L12 aluminum alloys

Cited By (2)

* Cited by examiner, † Cited by third party
Publication number Priority date Publication date Assignee Title
WO2020150056A1 (en) * 2019-01-18 2020-07-23 Divergent Technologies, Inc. Aluminum alloys
CN113508184A (en) * 2019-01-18 2021-10-15 戴弗根特技术有限公司 Aluminium alloy

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