Patent Grant
7879162
This application is related to the following co-pending applications that are filed on even date herewith and are assigned to the same assignee: L12 ALUMINUM ALLOYS WITH BIMODAL AND TRIMODAL DISTRIBUTION, Ser. No. 12/148,395; DISPERSION STRENGTHENED L12 ALUMINUM ALLOYS, Ser. No. 12/148,432; HEAT TREATABLE L12 ALUMINUM ALLOYS, Ser. No. 12/148,383; HIGH STRENGTH L12 ALUMINUM ALLOYS, Ser. No. 12/148,394; HIGH STRENGTH L12 ALUMINUM ALLOYS, Ser. No. 12/148,382; HEAT TREATABLE L12 ALUMINUM ALLOYS, Ser. No. 12/148,396; HIGH STRENGTH L12 ALUMINUM ALLOYS, Ser. No. 12/148,387; HIGH STRENGTH L12 ALUMINUM ALLOYS, Ser. No. 12/148,459; and L12 STRENGTHENED AMORPHOUS ALUMINUM ALLOYS, Ser. No. 12/148,458.
The present invention relates generally to aluminum alloys and more specifically to heat treatable aluminum alloys strengthened by L12 phase dispersions that are useful at temperatures from about −420° F. (−251° C.) up to about 650° F. (343° C.).
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.). U.S. Patent Application Publication No. 2006/0093512 by the current inventor discloses an aluminum magnesium alloy strengthened with a dispersion of Al3X dispersoids with the L12 structure where X comprises Sc, Gd, and Zr. The alloy provides excellent mechanical properties in the temperature range of about −420° F. (−250° C.) up to about 573° F. (300° C.).
An aluminum magnesium alloy strengthened by L12 precipitates with excellent mechanical properties in the temperature range of about 420° F. (−250° C.) to about 650° F. (343° C.) would be useful.
The present invention is an aluminum magnesium alloy that is strengthened with L12 dispersoids. The alloys have mechanical properties suitable for application at temperature ranges from about −420° F. (−251° C.) to about 650° F. (343° C.). The alloys comprise magnesium, coherent Al3Sc L12 dispersoids, and coherent Al3X L12 dispersoids where X is at least one element selected from scandium, erbium, thulium, ytterbium, and lutetium, and at least one element selected from gadolinium, yttrium, zirconium, titanium, hafnium, and niobium. The balance is substantially aluminum.
The alloys can also contain one or more elements selected from zinc, copper, lithium and silicon.
The alloys have less than about 1 weight percent total impurities.
The alloys can be formed by any rapid solidification technique that includes atomization, melt spinning, splat quenching, spray deposition, cold spray, plasma spray, laser melting, ball milling, and cryomilling. The alloys with smaller amounts of alloying elements can also be formed by casting and deformation processing.
The alloys can be heat treated at a temperature of about 800° F. (426° C.) to about 1100° F. (593° C.) for about 30 minutes to about four hours, followed by quenching in liquid and thereafter aged at a temperature of about 200° F. (93° C.) to about 600° F. (315° C.) for about two to about forty-eight hours.
The alloys of this invention are based on the aluminum magnesium system. The amount of magnesium in these alloys ranges from about 1 to about 8 weight percent, more preferably about 3 to about 7.5 weight percent, and even more preferably about 4 to about 6.5 weight percent.
The aluminum magnesium phase diagram is shown in
The alloys may also optionally contain at least one element selected from zinc, copper, lithium and silicon to produce additional strengthening. The amount of zinc in these alloys ranges from about 3 to about 12 weight percent, more preferably about 4 to about 10 weight percent, and even more preferably about 5 to about 9 weight percent. The amount of copper in these alloys ranges from about 0.2 to about 3 weight percent, more preferably about 0.5 to about 2.5 weight percent, and even more preferably about 1 to about 2.5 weight percent. The amount of lithium in these alloys ranges from about 0.5 to about 3 weight percent, more preferably about 1 to about 2.5 weight percent, and even more preferably about 1 to about 2 weight percent. The amount of silicon in these alloys ranges from about 4 to about 25 weight percent silicon, more preferably about 4 to about 18 weight percent, and even more preferably about 5 to about 11 weight percent.
Exemplary aluminum alloys of this invention include, but are not limited to (in weight percent):
about Al-(1-8)Mg-(0.1-4)Sc-(0.1-20)Gd;
about Al-(1-8)Mg-(0.1-20)Er-(0.1-20)Gd;
about Al-(1-8)Mg-(0.1-15)Tm-(0.1-20)Gd;
about Al-(1-8)Mg-(0.1-25)Yb-(0.1-20)Gd;
about Al-(1-8)Mg-(0.1-25)Lu-(0.1-20)Gd;
about Al-(1-8)Mg-(0.1-4)Sc-(0.1-20)Y;
about Al-(1-8)Mg-(0.1-20)Er-(0.1-20)Y;
about Al-(1-8)Mg-(0.1-15)Tm-(0.1-20)Y;
about Al-(1-8)Mg-(0.1-25)Yb-(0.1-20)Y;
about Al-(1-8)Mg-(0.1-25)Lu-(0.1-20)Y;
about Al-(1-8Mg-(0.1-4)Sc-(0.05-3.0)Zr;
about Al-(1-8)Mg-(0.1-20)Er-(0.05-4.0)Zr;
about Al-(1-8)Mg-(0.1-15)Tm-(0.05-4.0)Zr;
about Al-(1-8)Mg-(0.1-25)Yb-(0.05-4.0)Zr;
about Al-(1-8)Mg-(0.1-25)Lu-(0.05-4.0)Zr;
about Al-(1-8)Mg-(0.1-4)Sc-(0.05-10)Ti;
about Al-(1-8)Mg-(0.1-20)Er-(0.05-10)Ti;
about Al-(1-8)Mg-(0.1-15)Tm-(0.05-10)Ti;
about Al-(1-8)Mg-(0.1-25)Yb-(0.05-10)Ti;
about Al-(1-8)Mg-(0.1-25)Lu-(0.05-10)Ti;
about Al-(1-8)Mg-(0.1-4)Sc-(0.05-10)Hf;
about Al-(1-8)Mg-(0.1-20)Er-(0.05-10)Hf;
about Al-(1-8)Mg-(0.1-15)Tm-(0.05-10)Hf;
about Al-(1-8)Mg-(0.1-25)Yb-(0.05-10)Hf;
about Al-(1-8)Mg-(0.1-25)Lu-(0.05-10)Hf;
about Al-(1-8)Mg-(0.1-4)Sc-(0.05-5)Nb;
about Al-(1-8)Mg-(0.1-20)Er-(0.05-5)Nb;
about Al-(1-8)Mg-(0.1-15)Tm-(0.05-5)Nb;
about Al-(1-8)Mg-(0.1-25)Yb-(0.05-5)Nb; and
about Al-(1-8)Mg-(0.1-25)Lu-(0.05-5)Nb.
In the inventive aluminum based alloys disclosed herein, scandium is a potent strengthener that has low diffusivity and low solubility in aluminum. Scandium forms equilibrium Al3Sc intermetallic dispersoids that have an L12 structure that is an ordered face centered cubic structure with Sc 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. In the alloys of this invention these Al3Sc dispersoids are made stronger and more resistant to coarsening at elevated temperatures by adding suitable alloying elements such as gadolinium, yttrium, zirconium, titanium, hafnium, niobium, or combinations thereof, that enter Al3Sc in solution.
Erbium forms Al3Er dispersoids in the aluminum matrix that are fine and coherent with the aluminum matrix. The lattice parameters of aluminum and Al3Er are close (0.405 nm and 0.417 nm respectively), indicating there is minimal driving force for causing growth of the Al3Er dispersoids. This low interfacial energy makes the Al3Er dispersoids thermally stable and resistant to coarsening up to temperatures as high as about 842° F. (450° C.). Additions of magnesium in solid solution in aluminum increase the lattice parameter of the aluminum matrix, and decrease the lattice parameter mismatch further increasing the resistance of the Al3Er to coarsening. In the alloys of this invention, these Al3Er dispersoids are made stronger and more resistant to coarsening at elevated temperatures by adding suitable alloying elements such as gadolinium, yttrium, zirconium, titanium, hafnium, niobium, or combinations thereof that enter Al3Er in solution.
Thulium forms metastable Al3Tm dispersoids in the aluminum matrix that are fine and coherent with the aluminum matrix. The lattice parameters of aluminum and Al3Tm are close (0.405 nm and 0.420 nm respectively), indicating there is minimal driving force for causing growth of the Al3Tm dispersoids. This low interfacial energy makes the Al3Tm dispersoids thermally stable and resistant to coarsening up to temperatures as high as about 842° F. (450° C.). Additions of magnesium in solid solution in aluminum increase the lattice parameter of the aluminum matrix, and decrease the lattice parameter mismatch further increasing the resistance of the Al3Tm to coarsening. In the alloys of this invention these Al3Tm dispersoids are made stronger and more resistant to coarsening at elevated temperatures by adding suitable alloying elements such as gadolinium, yttrium, zirconium, titanium, hafnium, niobium, or combinations thereof that enter Al3Tm in solution.
Ytterbium forms Al3Yb dispersoids in the aluminum matrix that are fine and coherent with the aluminum matrix. The lattice parameters of Al and Al3Yb are close (0.405 nm and 0.420 nm respectively), indicating there is minimal driving force for causing growth of the Al3Yb dispersoids. This low interfacial energy makes the Al3Yb dispersoids thermally stable and resistant to coarsening up to temperatures as high as about 842° F. (450° C.). Additions of magnesium in solid solution in aluminum increase the lattice parameter of the aluminum matrix, and decrease the lattice parameter mismatch further increasing the resistance of the Al3Yb to coarsening. In the alloys of this invention, these Al3Yb dispersoids are made stronger and more resistant to coarsening at elevated temperatures by adding suitable alloying elements such as gadolinium, yttrium, zirconium, titanium, hafnium, niobium, or combinations thereof that enter Al3Yb in solution.
Lutetium forms Al3Lu dispersoids in the aluminum matrix that are fine and coherent with the aluminum matrix. The lattice parameters of Al and Al3Lu are close (0.405 nm and 0.419 nm respectively), indicating there is minimal driving force for causing growth of the Al3Lu dispersoids. This low interfacial energy makes the Al3Lu dispersoids thermally stable and resistant to coarsening up to temperatures as high as about 842° F. (450° C.). Additions of magnesium in solid solution in aluminum increase the lattice parameter of the aluminum matrix, and decrease the lattice parameter mismatch further increasing the resistance of the Al3Lu to coarsening. 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 and a D019 structure in the equilibrium condition. The Al3Gd dispersoids are stable up to temperatures as high as about 842° F. (450° C.) due to their low diffusivity in aluminum. Despite its large atomic size, gadolinium has fairly high solubility in the Al3Sc intermetallic dispersoid. 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.
Additions of zinc, copper, lithium and silicon increase the strength of these alloys through additional solid solution hardening and precipitation hardening of Zn2Mg (η′), Al2Cu (θ′), Al2CuMg (S′), Al3Li (δ′), Al2LiMg, Mg2Si and Si phases, respectively. These phases precipitate as coherent fine particles which can provide considerable strengthening in the alloys. Precipitation of these phases can be controlled during heat treatment.
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 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
The amount of erbium present in the alloys of this invention, if any, may vary from about 0.1 to about 20 weight percent, more preferably from about 0.3 to about 15 weight percent, and even more preferably from about 0.5 to about 10 weight percent. The Al—Er phase diagram shown in
The amount of thulium present in the alloys of this invention, if any, may vary from about 0.1 to about 15.0 weight percent, more preferably from about 0.2 to about 10 weight percent, and even more preferably from about 0.4 to about 6 weight percent. The Al—Tm phase diagram shown in
The amount of ytterbium present in the alloys of this invention, if any, may vary from about 0.1 to about 25 weight percent, more preferably from about 0.3 to about 20 weight percent, and even more preferably from about 0.4 to about 10 weight percent. The Al—Yb phase diagram shown in
The amount of lutetium present in the alloys of this invention, if any, may vary from about 0.1 to about 25 weight percent, more preferably from about 0.3 to about 20 weight percent, and even more preferably from about 0.4 to about 10 weight percent. The Al—Lu phase diagram shown in
The amount of gadolinium present in the alloys of this invention, if any, may vary from about 0.1 to about 20 weight percent, more preferably from about 0.3 to about 15 weight percent, and even more preferably from about 0.5 to about 10 weight percent.
The amount of yttrium present in the alloys of this invention, if any, may vary from about 0.1 to about 20 weight percent, more preferably from about 0.3 to about 15 weight percent, and even more preferably from about 0.5 to about 10 weight percent.
The amount of zirconium present in the alloys of this invention, if any, may vary from about 0.05 to about 4 weight percent, more preferably from about 0.1 to about 3 weight percent, and even more preferably from about 0.3 to about 2 weight percent.
The amount of titanium present in the alloys of this invention, if any, may vary from about 0.05 to about 10 weight percent, more preferably from about 0.2 to about 8 weight percent, and even more preferably from about 0.4 to about 4 weight percent.
The amount of hafnium present in the alloys of this invention, if any, may vary from about 0.05 to about 10 weight percent, more preferably from about 0.2 to about 8 weight percent, and even more preferably from about 0.4 to about 5 weight percent.
The amount of niobium present in the alloys of this invention, if any, may vary from about 0.05 to about 5 weight percent, more preferably from about 0.1 to about 3 weight percent, and even more preferably from about 0.2 to about 2 weight percent.
In order to have the best properties for the 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.
The alloy of the present invention can be processed by any rapid solidification technique utilizing cooling rates in excess of 103° C./second. The rapid solidification process includes melt spinning, splat quenching, atomization, spray deposition, cold spray, vacuum plasma spray, and laser melting. The particular processing technique is not important. The most important aspect is the cooling rate of the process. A higher cooling rate is required for the alloys with larger amount of solute additions. These processes produce different forms of the product such as ribbon, flake or powder. Atomization is the most commonly used rapid solidification technique to produce a large volume of powder. Cooling rate experienced during atomization depends on the powder size and usually varies from about 103 to 105° C./second. Finer size (−325 mesh) of powder is preferred to have maximum supersaturation of alloying elements that can precipitate out during extrusion of powder. For higher supersaturation of alloying elements, helium gas atomization is preferred. Helium gas provides higher heat transfer coefficient leading to higher cooling rate in the powder. The ribbon or powder of alloy can be compacted using vacuum hot pressing, hot isostatic pressing or blind die compaction after suitable vacuum degassing. Compaction takes place by shear deformation in vacuum hot pressing and blind die compaction, whereas diffusional creep is key for compaction in hot isostatic pressing.
The alloy powder of the present invention can also be produced using mechanical alloying or cryomilling where powder is milled using high energy ball milling at room temperature or at cryogenic temperature in liquid nitrogen environment. While both mechanical alloying and cryomilling processes can provide supersaturation of alloying elements, cryomilling is preferred because it has less oxygen content. Cryomilling introduces oxynitride particles in the grains that can provide additional strengthening to the alloy at high temperature by increasing threshold stress for dislocation climb. In addition, the nitride particles when located on grain boundaries can reduce the grain boundary sliding in the alloy by pinning the dislocation resulting in reduced dislocation mobility in the grain boundaries.
The alloy may also be produced using casting processes such as squeeze casting, die casting, sand casting and permanent mold casting provided the alloy contains small amounts of Sc, Er, Tm, Yb, Lu, Gd, Y, Ti, Hf, or Nb.
Following consolidation and deformation processing, the alloys can be heat treated at a temperature of from about 800° F. (426° C.) to about 1100° F. (593° C.) for about thirty minutes to four hours followed by quenching in liquid and thereafter aged at a temperature from about 200° F. (93° C.) to about 600° F. (316° C.) for about two to forty-eight hours.
More preferred exemplary aluminum alloys of this invention include, but are not limited to (in weight percent):
about Al-(3-7.5)Mg-(0.1-3)Sc-(0.3-15)Gd;
about Al-(3-7.5)Mg-(0.3-15)Er-(0.3-15)Gd;
about Al-(3-7.5)Mg-(0.2-10)Tm-(0.3-15)Gd;
about Al-(3-7.5)Mg-(0.3-20)Yb-(0.3-15)Gd;
about Al-(3-7.5)Mg-(0.3-20)Lu-(0.3-15)Gd;
about Al-(3-7.5)Mg-(0.1-3)Sc-(0.3-15)Y;
about Al-(3-7.5)Mg-(0.3-15)Er-(0.3-15)Y;
about Al-(3-7.5)Mg-(0.2-10)Tm-(0.3-15)Y;
about Al-(3-7.5)Mg-(0.3-20)Yb-(0.3-15)Y;
about Al-(3-7.5)Mg-(0.3-20)Lu-(0.3-15)Y;
about Al-(3-7.5)Mg-(0.1-3)Sc-(0.1-3)Zr;
about Al-(3-7.5)Mg-(0.3-15)Er-(0.1-3)Zr;
about Al-(3-7.5)Mg-(0.2-10)Tm-(0.1-3)Zr;
about Al-(3-7.5)Mg-(0.3-20)Yb-(0.1-3)Zr;
about Al-(3-7.5)Mg-(0.3-20)Lu-(0.1-3)Zr;
about Al-(3-7.5)Mg-(0.1-3)Sc-(0.2-8)Ti;
about Al-(3-7.5)Mg-(0.3-15)Er-(0.2-8)Ti;
about Al-(3-7.5)Mg-(0.2-10)Tm-(0.2-8)Ti;
about Al-(3-7.5)Mg-(0.3-20)Yb-(0.2-8)Ti;
about Al-(3-7.5)Mg-(0.3-20)Lu-(0.2-8)Ti;
about Al-(3-7.5)Mg-(0.1-3)Sc-(0.2-8)Hf;
about Al-(3-7.5)Mg-(0.3-15)Er-(0.2-8)Hf;
about Al-(3-7.5)Mg-(0.2-10)Tm-(0.2-8)Hf;
about Al-(3-7.5)Mg-(0.3-20)Yb-(0.2-8)Hf;
about Al-(3-7.5)Mg-(0.3-20)Lu-(0.2-8)Hf;
about Al-(3-7.5)Mg-(0.1-3)Sc-(0.1-3)Nb;
about Al-(3-7.5)Mg-(0.3-15)Er-(0.1-3)Nb;
about Al-(3-7.5)Mg-(0.2-10)Tm-(0.1-3)Nb;
about Al-(3-7.5)Mg-(0.3-20)Yb-(0.1-3)Nb; and
about Al-(3-7.5)Mg-(0.3-20)Lu-(0.1-3)Nb.
Even more preferred examples of similar alloys to these are alloys with about 4 to about 6.5 weight percent Mg. These exemplary alloys may also optionally contain at least one of the elements from zinc, copper, lithium and silicon to produce additional strengthening, and include, but are not limited to (in weight percent):
about Al-(4-6.5)Mg-(0.2-2.5)Sc-(0.5-10)Gd;
about Al-(4-6.5)Mg-(0.5-10)Er-(0.5-10)Gd;
about Al-(4-6.5)Mg-(0.4-6)Tm-(0.5-10)Gd;
about Al-(4-6.5)Mg-(0.4-10)Yb-(0.5-10)Gd;
about Al-(4-6.5)Mg-(0.4-10)Lu-(0.5-10)Gd;
about Al-(4-6.5)Mg-(0.2-2.5)Sc-(0.5-10)Y;
about Al-(4-6.5)Mg-(0.5-10)Er-(0.5-10)Y;
about Al-(4-6.5)Mg-(0.4-6)Tm-(0.5-10)Y;
about Al-(4-6.5)Mg-(0.4-10)Yb-(0.5-10)Y;
about Al-(4-6.5)Mg-(0.4-10)Lu-(0.5-10)Y;
about Al-(4-6.5)Mg-(0.2-2.5)Sc-(0.3-2)Zr;
about Al-(4-6.5)Mg-(0.5-10O)Er-(0.3-2)Zr;
about Al-(4-6.5)Mg-(0.4-6)Tm-(0.3-2)Zr;
about Al-(4-6.5)Mg-(0.4-10)Yb-(0.3-2)Zr;
about Al-(4-6.5)Mg-(0.4-10)Lu-(0.3-2)Zr;
about Al-(4-6.5)Mg-(0.2-2.5)Sc-(0.4-4)Ti;
about Al-(4-6.5)Mg-(0.5-10)Er-(0.4-4)Ti;
about Al-(4-6.5)Mg-(0.4-6)Tm-(0.4-4)Ti;
about Al-(4-6.5)Mg-(0.4-10)Yb-(0.4-4)Ti;
about Al-(4-6.5)Mg-(0.4-10)Lu-(0.4-4)Ti;
about Al-(4-6.5)Mg-(0.2-2.5)Sc-(0.4-5)Hf;
about Al-(4-6.5)Mg-(0.5-10)Er-(0.4-5)Hf;
about Al-(4-6.5)Mg-(0.4-6)Tm-(0.4-5)Hf;
about Al-(4-6.5)Mg-(0.4-10)Yb-(0.4-5)Hf;
about Al-(4-6.5)Mg-(0.4-10)Lu-(0.4-5)Hf;
about Al-(4-6.5)Mg-(0.2-2.5)Sc-(0.2-2)Nb;
about Al-(4-6.5)Mg-(0.5-10)Er-(0.2-2)Nb;
about Al-(4-6.5)Mg-(0.4-6)Tm-(0.2-2)Nb;
about Al-(4-6.5)Mg-(0.4-10)Yb-(0.2-2)Nb; and
about Al-(4-6.5)Mg-(0.4-10)Lu-(0.2-2)Nb.
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.
| Number | Name | Date | Kind |
|---|---|---|---|
| 3619181 | Willey et al. | Nov 1971 | A |
| 3816080 | Bomford et al. | Jun 1974 | A |
| 4041123 | Lange et al. | Aug 1977 | A |
| 4259112 | Dolowy, Jr. et al. | Mar 1981 | A |
| 4463058 | Hood et al. | Jul 1984 | A |
| 4469537 | Ashton et al. | Sep 1984 | A |
| 4499048 | Hanejko | Feb 1985 | A |
| 4597792 | Webster | Jul 1986 | A |
| 4626294 | Sanders, Jr. | Dec 1986 | A |
| 4647321 | Adam | Mar 1987 | A |
| 4661172 | Skinner et al. | Apr 1987 | A |
| 4667497 | Oslin et al. | May 1987 | A |
| 4689090 | Sawtell et al. | Aug 1987 | A |
| 4710246 | Le Caer et al. | Dec 1987 | A |
| 4713216 | Higashi et al. | Dec 1987 | A |
| 4755221 | Paliwal et al. | Jul 1988 | A |
| 4853178 | Oslin | Aug 1989 | A |
| 4865806 | Skibo et al. | Sep 1989 | A |
| 4874440 | Sawtell et al. | Oct 1989 | A |
| 4915605 | Chan et al. | Apr 1990 | A |
| 4927470 | Cho | May 1990 | A |
| 4933140 | Oslin | Jun 1990 | A |
| 4946517 | Cho | Aug 1990 | A |
| 4964927 | Shiflet et al. | Oct 1990 | A |
| 4988464 | Riley | Jan 1991 | A |
| 5032352 | Meeks et al. | Jul 1991 | A |
| 5053084 | Masumoto et al. | Oct 1991 | A |
| 5055257 | Chakrabarti et al. | Oct 1991 | A |
| 5059390 | Burleigh et al. | Oct 1991 | A |
| 5066342 | Rioja et al. | Nov 1991 | A |
| 5076340 | Bruski et al. | Dec 1991 | A |
| 5076865 | Hashimoto et al. | Dec 1991 | A |
| 5130209 | Das et al. | Jul 1992 | A |
| 5133931 | Cho | Jul 1992 | A |
| 5198045 | Cho et al. | Mar 1993 | A |
| 5211910 | Pickens et al. | May 1993 | A |
| 5226983 | Skinner et al. | Jul 1993 | A |
| 5256215 | Horimura | Oct 1993 | A |
| 5308410 | Horimura et al. | May 1994 | A |
| 5312494 | Horimura et al. | May 1994 | A |
| 5318641 | Masumoto et al. | Jun 1994 | A |
| 5397403 | Horimura et al. | Mar 1995 | A |
| 5458700 | Masumoto et al. | Oct 1995 | A |
| 5462712 | Langan et al. | Oct 1995 | A |
| 5480470 | Miller et al. | Jan 1996 | A |
| 5597529 | Tack | Jan 1997 | A |
| 5620652 | Tack et al. | Apr 1997 | A |
| 5624632 | Baumann et al. | Apr 1997 | A |
| 5882449 | Waldron et al. | Mar 1999 | A |
| 6139653 | Fernandes et al. | Oct 2000 | A |
| 6149737 | Hattori et al. | Nov 2000 | A |
| 6248453 | Watson | Jun 2001 | B1 |
| 6254704 | Laul et al. | Jul 2001 | B1 |
| 6258318 | Lenczowski et al. | Jul 2001 | B1 |
| 6309594 | Meeks, III et al. | Oct 2001 | B1 |
| 6312643 | Upadhya et al. | Nov 2001 | B1 |
| 6315948 | Lenczowski et al. | Nov 2001 | B1 |
| 6331218 | Inoue et al. | Dec 2001 | B1 |
| 6355209 | Dilmore et al. | Mar 2002 | B1 |
| 6368427 | Sigworth | Apr 2002 | B1 |
| 6506503 | Mergen et al. | Jan 2003 | B1 |
| 6517954 | Mergen et al. | Feb 2003 | B1 |
| 6524410 | Kramer et al. | Feb 2003 | B1 |
| 6531004 | Lenczowski et al. | Mar 2003 | B1 |
| 6562154 | Rioja et al. | May 2003 | B1 |
| 6630008 | Meeks, III et al. | Oct 2003 | B1 |
| 6702982 | Chin et al. | Mar 2004 | B1 |
| 6902699 | Fritzemeier et al. | Jun 2005 | B2 |
| 6918970 | Lee et al. | Jul 2005 | B2 |
| 6974510 | Watson | Dec 2005 | B2 |
| 7048815 | Senkov et al. | May 2006 | B2 |
| 7097807 | Meeks, III et al. | Aug 2006 | B1 |
| 7241328 | Keener | Jul 2007 | B2 |
| 7344675 | Van Daam et al. | Mar 2008 | B2 |
| 20010054247 | Stall et al. | Dec 2001 | A1 |
| 20030192627 | Lee et al. | Oct 2003 | A1 |
| 20040046402 | Winardi | Mar 2004 | A1 |
| 20040055671 | Olson et al. | Mar 2004 | A1 |
| 20040089382 | Senkov et al. | May 2004 | A1 |
| 20040170522 | Watson | Sep 2004 | A1 |
| 20040191111 | Nie et al. | Sep 2004 | A1 |
| 20050147520 | Canzona | Jul 2005 | A1 |
| 20060011272 | Lin et al. | Jan 2006 | A1 |
| 20060093512 | Pandey | May 2006 | A1 |
| 20060172073 | Groza et al. | Aug 2006 | A1 |
| 20060269437 | Pandey | Nov 2006 | A1 |
| 20070048167 | Yano | Mar 2007 | A1 |
| 20070062669 | Song et al. | Mar 2007 | A1 |
| 20080066833 | Lin et al. | Mar 2008 | A1 |
| Number | Date | Country |
|---|---|---|
| 1436870 | Aug 2003 | CN |
| 101205578 | Jun 2008 | CN |
| 0 208 631 | Jun 1986 | EP |
| 0 584 596 | Mar 1995 | EP |
| 1 111 079 | Dec 1999 | EP |
| 1 111 078 | Jun 2001 | EP |
| 1 249 303 | Oct 2002 | EP |
| 1 170 394 | Apr 2004 | EP |
| 1 439 239 | Jul 2004 | EP |
| 1 471 157 | Oct 2004 | EP |
| 1 728 881 | Dec 2006 | EP |
| 1 788 102 | May 2007 | EP |
| 2 656 629 | Dec 1990 | FR |
| 2843754 | Feb 2004 | FR |
| 04218638 | Dec 1990 | JP |
| 9104940 | Apr 1997 | JP |
| 9279284 | Oct 1997 | JP |
| 11156584 | Jun 1999 | JP |
| 2000119786 | Apr 2000 | JP |
| 2001038442 | Feb 2001 | JP |
| 2007188878 | Jul 2007 | JP |
| 2001144 | Oct 1993 | RU |
| 2001145 | Oct 1993 | RU |
| 90 02620 | Mar 1990 | WO |
| 91 10755 | Jan 1991 | WO |
| 9111540 | Aug 1991 | WO |
| WO 95 32074 | Nov 1995 | WO |
| WO 9610099 | Apr 1996 | WO |
| 9833947 | Aug 1998 | WO |
| 00 37696 | Jun 2000 | WO |
| 02 29139 | Apr 2002 | WO |
| 03 052154 | Jun 2003 | WO |
| WO 03 085145 | Oct 2003 | WO |
| WO 03 085146 | Oct 2003 | WO |
| 03 104505 | Dec 2003 | WO |
| 2004 005562 | Jan 2004 | WO |
| 2004046402 | Jun 2004 | WO |
| 2005045080 | May 2005 | WO |
| 2005047554 | May 2005 | WO |
| Number | Date | Country | |
|---|---|---|---|
| 20090263276 A1 | Oct 2009 | US |
