# Patent

## Title
L12 aluminum alloys with bimodal and trimodal distribution

## Abstract

A two or three phase aluminum alloy having high strength, modulus, ductility and toughness, comprising a fine grain matrix phase nano L12 alloy having a particle size ranging from about 20 nm to 5 microns and a more ductile larger aluminum alloy coarse grain phase having a particle size ranging from about 25 to 250 microns. The fine grain matrix phase alloy comprises aluminum, at least one of scandium, erbium, thulium, ytterbium, and lutetium; and at least one of gadolinium, yttrium, zirconium, titanium, hafnium, and niobium. The alloy may also include ceramic reinforcements in addition to the fine grain matrix phase and the coarse grain phase.

## Background
CROSS-REFERENCE TO RELATED APPLICATIONS
This application is related to the following co-pending applications that are filed on even date herewith and are assigned to the same assignee: DISPERSION STRENGTHENED L12 ALUMINUM ALLOYS, Ser. No. 12/148,432; HEAT TREATABLE L12 ALUMINUM ALLOYS, Ser. No. 12/148,383; HIGH STRENGTH L12 ALUMINUM ALLOYS, Ser. No. 12/148,394; HIGH STRENGTH L12 ALUMINUM ALLOYS, Ser. No. 12/148,382; HEAT TREATABLE L12 ALUMINUM ALLOYS, Ser. No. 12/148,396; HIGH STRENGTH L12 ALUMINUM ALLOYS, Ser. No. 12/148,387; HIGH STRENGTH ALUMINUM ALLOYS WITH L12 PRECIPITATES, Ser. No. 12/148,426; HIGH STRENGTH L12ALUMINUM ALLOYS, Ser. No. 12/148,459; and L12 STRENGTHENED AMORPHOUS ALUMINUM ALLOYS, Ser. No. 12/148,458.
BACKGROUND
The present invention relates generally to aluminum alloys and more specifically to two and three phase L12 aluminum alloys with high strength and improved ductility.
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 Lu. The Al3X particles are coherent with the aluminum alloy matrix and are resistant to coarsening at elevated temperatures. The improved mechanical properties of the disclosed dispersion strengthened L12 aluminum alloys are stable up to 572° F. (300° C.). U.S. Patent Application Publication No. 2006/0269437 A1 discloses a high strength aluminum alloy that contains scandium and other elements that is strengthened by L12 dispersoids.
L12 strengthened aluminum alloys have high strength and improved fatigue properties compared to commercial aluminum alloys. Fine grain size results in improved mechanical properties of materials. Hall-Petch strengthening has been known for decades where strength increases as grain size decreases. An optimum grain size for optimum strength is in the nano range of about 30 to 100 nm. These alloys also have lower ductility.
Fine grain aluminum alloys exhibit high strength but the lower ductility leads to lower fracture toughness. It would be desirable to develop a high strength aluminum alloy with acceptable fracture toughness.
SUMMARY
The present invention is aluminum alloys with high strength and acceptable fracture toughness. In embodiments, these properties are achieved with a dual phase microstructure. The microstructure consists of a fine grain matrix strengthened by a dispersion of coherent L12 Al3X dispersoids where X is at least one first element selected from scandium, erbium, thulium, ytterbium, and lutetium, and at least one second element selected from gadolinium, yttrium, zirconium, titanium, hafnium, and niobium. The balance is substantially aluminum containing at least one alloying element selected from silicon, magnesium, lithium, copper, zinc, and nickel.
The addition of larger ductile aluminum alloy second phase particles (i.e., coarse grains) dispersed throughout the high strength fine grain matrix act to increase the fracture toughness by acting to blunt crack tips, thereby reducing the driving force for crack growth. The size of the coarse grains ranges from about 25 to about 250 microns, more preferably about 50 to about 200 microns, and even more preferably about 100 to about 150 microns. To provide a balanced combination of strength, ductility and toughness, the volume fraction of the coarse grains ranges from about 10 to about 50 volume percent, more preferably about 15 to about 40 volume percent, and even more preferably about 20 to about 30 volume percent. The resulting aluminum alloy is a bimodal system alloy.
It has also been discovered that ceramic reinforcements may be added to the bimodal system alloys to further increase their strength and modulus, thereby forming a trimodal system alloy. Trimodal system alloys have a higher load transfer mechanism that results in higher strength and modulus. Aluminum oxide, silicon carbide, boron carbide, aluminum nitride, titanium boride, titanium diboride and titanium carbide are suitable ceramic reinforcements.
In order to be effective, the reinforcing ceramic particles need to have fine size, moderate volume fraction and a good interface between the matrix and reinforcement. These ceramic reinforcements can have particle sizes ranging from about 0.5 to about 50 microns, more preferably about 1 to about 20 microns, and even more preferably 1 to about 10 microns. The volume fraction of ceramic reinforcements that may be added ranges from about 5 to about 40 volume percent, more preferably about 10 to about 30 volume percent, and even more preferably about 15 to about 25 volume percent. These fine ceramic reinforcement particles located at the grain boundary and within the grain boundary will restrict the dislocation from going around particles. The dislocations become attached with particles on the departure side, and thus require more energy to detach the dislocation.



BRIEF DESCRIPTION OF THE DRAWINGS
FIG. 1 is an aluminum scandium phase diagram.
FIG. 2 is an aluminum erbium phase diagram.
FIG. 3 is an aluminum thulium phase diagram.
FIG. 4 is an aluminum ytterbium phase diagram.
FIG. 5 is an aluminum lutetium phase diagram.



DETAILED DESCRIPTION
The alloys of this invention are based on aluminum based alloys with a bimodal or trimodal microstructure with high strength and fracture toughness for applications at temperatures from about −420° F. (−251° C.) up to about 650° F. (343° C.). A bimodal microstructure consists of relatively coarse ductile grains in a high strength ultra fine grain aluminum alloy matrix strengthened with coherent L12 Al3X dispersoids where X is at least one first element selected from scandium, erbium, thulium, ytterbium, and lutetium, and at least one second element selected from gadolinium, yttrium, zirconium, titanium, hafnium, and niobium. A trimodal microstructure is created when ceramic reinforcements are added to the bimodal system alloy.
The coarse grains can comprise any aluminum alloy with sufficient strength and ductility.
The high strength fine grain matrix of this invention comprises a solid solution of aluminum and at least one element selected from silicon, magnesium, lithium, copper, zinc, and nickel strengthened by L12 coherent precipitates where X is at least one first element selected from scandium, erbium, thulium, ytterbium, and lutetium, and at least one second element selected from gadolinium, yttrium, zirconium, titanium, hafnium, and niobium.
The aluminum silicon system is a simple eutectic alloy system with a eutectic reaction at 12.5 weight percent silicon and 1077° F. (577° C.). There is little solubility of silicon in aluminum at temperatures up to 930° F. (500° C.) and none of aluminum in silicon. However, the solubility can be extended significantly by utilizing rapid solidification techniques.
The binary aluminum magnesium system is a simple eutectic at 36 weight percent magnesium and 842° F. (450° C.). There is complete solubility of magnesium and aluminum in the rapidly solidified inventive alloys discussed herein.
The binary aluminum lithium system is a simple eutectic at 8 weight percent lithium and 1105° (596° C.). The equilibrium solubility of 4 weight percent lithium can be extended significantly by rapid solidification techniques. There is complete solubility of lithium in the rapid solidified inventive alloys discussed herein.
The binary aluminum copper system is a simple eutectic at 32 weight percent copper and 1018° F. (548° C.). There is complete solubility of copper in the rapidly solidified inventive alloys discussed herein.
The aluminum zinc binary system is a eutectic alloy system involving a monotectoid reaction and a miscibility gap in the solid state. There is a eutectic reaction at 94 weight percent zinc and 718° F. (381° C.). Zinc has maximum solid solubility of 83.1 weight percent in aluminum at 717.8° F. (381° C.) which can be extended by rapid solidification processes. Decomposition of the super saturated solid solution of zinc in aluminum gives rise to spherical and ellipsoidal GP zones which are coherent with the matrix and act to strengthen the alloy.
The aluminum nickel binary system is a simple eutectic at 5.7 weight percent nickel and 1183.8° F. (639.9° C.). There is little solubility of nickel in aluminum. However, the solubility can be extended significantly by utilizing rapid solidification processes. The equilibrium phase in the aluminum nickel eutectic system is L12 intermetallic Al3Ni.
In the 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, thulium, ytterbium, and lutetium, that have an L12 structure that is an ordered face centered cubic structure with the X atoms located at the corners and aluminum atoms located on the cube faces of the unit cell.
Scandium forms Al3Sc dispersoids that are fine and coherent with the aluminum matrix. Lattice parameters of aluminum and Al3Sc are very close (0.405 nm and 0.410 nm respectively), indicating that there is minimal or no driving force for causing growth of the Al3Sc dispersoids. This low interfacial energy makes the Al3Sc dispersoids thermally stable and resistant to coarsening up to temperatures as high as about 842° F. (450° C.). Additions of magnesium in aluminum increase the lattice parameter of the aluminum matrix, and decrease the lattice parameter mismatch further increasing the resistance of the Al3Sc to coarsening. Additions of zinc, copper, lithium, silicon, and nickel provide solid solution and precipitation strengthening in the aluminum alloys. In the alloys of this invention these Al3Sc dispersoids are made stronger and more resistant to coarsening at elevated temperatures by adding suitable alloying elements such as gadolinium, yttrium, zirconium, titanium, hafnium, niobium, or combinations thereof, that enter Al3Sc in solution.
Erbium forms Al3Er dispersoids in the aluminum matrix that are fine and coherent with the aluminum matrix. The lattice parameters of aluminum and Al3Er are close (0.405 nm and 0.417 nm respectively), indicating there is minimal driving force for causing growth of the Al3Er dispersoids. This low interfacial energy makes the Al3Er dispersoids thermally stable and resistant to coarsening up to temperatures as high as about 842° F. (450° C.). Additions of magnesium in aluminum increase the lattice parameter of the aluminum matrix, and decrease the lattice parameter mismatch further increasing the resistance of the Al3Er to coarsening. Additions of zinc, copper, lithium, silicon, and nickel provide solid solution and precipitation strengthening in the aluminum alloys. In the alloys of this invention, these Al3Er dispersoids are made stronger and more resistant to coarsening at elevated temperatures by adding suitable alloying elements such as gadolinium, yttrium, zirconium, titanium, hafnium, niobium, or combinations thereof that enter Al3Er in solution.
Thulium forms metastable Al3Tm dispersoids in the aluminum matrix that are fine and coherent with the aluminum matrix. The lattice parameters of aluminum and Al3Tm are close (0.405 nm and 0.420 nm respectively), indicating there is minimal driving force for causing growth of the Al3Tm dispersoids. This low interfacial energy makes the Al3Tm dispersoids thermally stable and resistant to coarsening up to temperatures as high as about 842° F. (450° C.). Additions of magnesium in aluminum increase the lattice parameter of the aluminum matrix, and decrease the lattice parameter mismatch further increasing the resistance of the Al3Tm to coarsening. Additions of zinc, copper, lithium, silicon, and nickel provide solid solution and precipitation strengthening in the aluminum alloys. In the alloys of this invention these Al3Tm dispersoids are made stronger and more resistant to coarsening at elevated temperatures by adding suitable alloying elements such as gadolinium, yttrium, zirconium, titanium, hafnium, niobium, or combinations thereof that enter Al3Tm in solution.
Ytterbium forms Al3Yb dispersoids in the aluminum matrix that are fine and coherent with the aluminum matrix. The lattice parameters of Al and Al3Yb are close (0.405 nm and 0.420 mm respectively), indicating there is minimal driving force for causing growth of the Al3Yb dispersoids. This low interfacial energy makes the Al3Yb dispersoids thermally stable and resistant to coarsening up to temperatures as high as about 842° F. (450° C.). Additions of magnesium in aluminum increase the lattice parameter of the aluminum matrix, and decrease the lattice parameter mismatch further increasing the resistance of the Al3Yb to coarsening. Additions of zinc, copper, lithium, silicon, and nickel provide solid solution and precipitation strengthening in the aluminum alloys. In the alloys of this invention, these Al3Yb dispersoids are made stronger and more resistant to coarsening at elevated temperatures by adding suitable alloying elements such as gadolinium, yttrium, zirconium, titanium, hafnium, niobium, or combinations thereof that enter Al3Yb in solution.
Lutetium forms Al3Lu dispersoids in the aluminum matrix that are fine and coherent with the aluminum matrix. The lattice parameters of Al and Al3Lu are close (0.405 nm and 0.419 nm respectively), indicating there is minimal driving force for causing growth of the Al3Lu dispersoids. This low interfacial energy makes the Al3Lu dispersoids thermally stable and resistant to coarsening up to temperatures as high as about 842° F. (450° C.). Additions of magnesium in aluminum increase the lattice parameter of the aluminum matrix, and decrease the lattice parameter mismatch further increasing the resistance of the Al3Lu to coarsening. Additions of zinc, copper, lithium, silicon, and nickel provide solid solution and precipitation strengthening in the aluminum alloys. In the alloys of this invention, these Al3Lu dispersoids are made stronger and more resistant to coarsening at elevated temperatures by adding suitable alloying elements such as gadolinium, yttrium, zirconium, titanium, hafnium, niobium, or mixtures thereof that enter Al3Lu in solution.
Gadolinium forms metastable Al3Gd dispersoids in the aluminum matrix that are stable up to temperatures as high as about 842° F. (450° C.) due to their low diffusivity in aluminum. The Al3Gd dispersoids have a D019 structure in the equilibrium condition. Despite its large atomic size, gadolinium has fairly high solubility in the Al3X intermetallic dispersoids (where X is scandium, erbium, thulium, ytterbium or lutetium). Gadolinium can substitute for the X atoms in Al3X intermetallic, thereby forming an ordered L12 phase which results in improved thermal and structural stability.
Yttrium forms metastable Al3Y dispersoids in the aluminum matrix that have an L12 structure in the metastable condition and a D019 structure in the equilibrium condition. The metastable Al3Y dispersoids have a low diffusion coefficient which makes them thermally stable and highly resistant to coarsening. Yttrium has a high solubility in the Al3X intermetallic dispersoids allowing large amounts of yttrium to substitute for X in the Al3X L12 dispersoids which results in improved thermal and structural stability.
Zirconium forms Al3Zr dispersoids in the aluminum matrix that have an L12 structure in the metastable condition and D023 structure in the equilibrium condition. The metastable Al3Zr dispersoids have a low diffusion coefficient which makes them thermally stable and highly resistant to coarsening. Zirconium has a high solubility in the Al3X dispersoids allowing large amounts of zirconium to substitute for X in the Al3X dispersoids, which results in improved thermal and structural stability.
Titanium forms Al3Ti dispersoids in the aluminum matrix that have an L12 structure in the metastable condition and D022 structure in the equilibrium condition. The metastable Al3Ti despersoids have a low diffusion coefficient which makes them thermally stable and highly resistant to coarsening. Titanium has a high solubility in the Al3X dispersoids allowing large amounts of titanium to substitute for X in the Al3X dispersoids, which result in improved thermal and structural stability.
Hafnium forms metastable Al3Hf dispersoids in the aluminum matrix that have an L12 structure in the metastable condition and a D023 structure in the equilibrium condition. The Al3Hf dispersoids have a low diffusion coefficient, which makes them thermally stable and highly resistant to coarsening. Hafnium has a high solubility in the Al3X dispersoids allowing large amounts of hafnium to substitute for scandium, erbium, thulium, ytterbium, and lutetium in the above mentioned Al3X dispersoids, which results in stronger and more thermally stable dispersoids.
Niobium forms metastable Al3Nb dispersoids in the aluminum matrix that have an L12 structure in the metastable condition and a D022 structure in the equilibrium condition. Niobium has a lower solubility in the Al3X dispersoids than hafnium or yttrium, allowing relatively lower amounts of niobium than hafnium or yttrium to substitute for X in the Al3X dispersoids. Nonetheless, niobium can be very effective in slowing down the coarsening kinetics of the Al3X dispersoids because the Al3Nb dispersoids are thermally stable. The substitution of niobium for X in the above mentioned Al3X dispersoids results in stronger and more thermally stable dispersoids.
Al3X L12 precipitates improve elevated temperature mechanical properties in aluminum alloys for two reasons. First, the precipitates are ordered intermetallic compounds. As a result, when the particles are sheared by glide dislocations during deformation, the dislocations separate into two partial dislocations separated by an anti-phase boundary on the glide plane. The energy to create the anti-phase boundary is the origin of the strengthening. Second, the cubic L12 crystal structure and lattice parameter of the precipitates are closely matched to the aluminum solid solution matrix. This results in a lattice coherency at the precipitate/matrix boundary that resists coarsening. The lack of an interphase boundary results in a low driving force for particle growth and resulting elevated temperature stability. Alloying elements in solid solution in the dispersed strengthening particles and in the aluminum matrix that tend to decrease the lattice mismatch between the matrix and particles will tend to increase the strengthening and elevated temperature stability of the alloy.
Exemplary aluminum alloys for the bimodal system alloys of this invention include, but are not limited to (in weight percent unless otherwise specified):
about Al-M-(10-50 vol. %)CG-(0.1-4)Sc-(0.1-20)Gd;
about Al-M-(10-50 vol. %)CG-(0.1-20)Er-(0.1-20)Gd;
about Al-M-(10-50 vol. %)CG-(0.1-15)Tm-(0.1-20)Gd;
about Al-M-(10-50 vol. %)CG-(0.1-25)Yb-(0.1-20)Gd;
about Al-M-(10-50 vol. %)CG-(0.1-25)Lu-(0.1-20)Gd;
about Al-M-(10-50 vol. %)CG-(0.1-4)Sc-(0.1-20)Y;
about Al-M-(10-50 vol. %)CG-(0.1-20)Er-(0.1-20)Y;
about Al-M-(10-50 vol. %)CG-(0.1-15)Tm-(0.1-20)Y;
about Al-M-(10-50 vol. %)CG-(0.1-25)Yb-(0.1-20)Y;
about Al-M-(10-50 vol. %)CG-(0.1-25)Lu-(0.1-20)Y;
about Al-M-(10-50 vol. %)CG-(0.1-4)Sc-(0.05-4)Zr;
about Al-M-(10-50 vol. %)CG-(0.1-20)Er-(0.05-4)Zr;
about Al-M-(10-50 vol. %)CG-(0.1-15)Tm-(0.05-4)Zr;
about Al-M-(10-50 vol. %)CG-(0.1-25)Yb-(0.05-4)Zr;
about Al-M-(10-50 vol. %)CG-(0.1-25)Lu-(0.05-4)Zr;
about Al-M-(10-50 vol. %)CG-(0.1-4)Sc-(0.05-10)Ti;
about Al-M-(10-50 vol. %)CG-(0.1-20)Er-(0.05-10)Ti;
about Al-M-(10-50 vol. %)CG-(0.1-15)Tm-(0.05-10)Ti;
about Al-M-(10-50 vol. %)CG-(0.1-25)Yb-(0.05-10)Ti;
about Al-M-(10-50 vol. %)CG-(0.1-25)Lu-(0.05-10)Ti;
about Al-M-(10-50 vol. %)CG-(0.1-4)Sc-(0.05-10)Hf;
about Al-M-(10-50 vol. %)CG-(0.1-20)Er-(0.05-10)Hf;
about Al-M-(10-50 vol. %)CG-(0.1-15)Tm-(0.05-10)Hf;
about Al-M-(10-50 vol. %)CG-(0.1-25)Yb-(0.05-10)Hf;
about Al-M-(10-50 vol. %)CG-(0.1-25)Lu-(0.05-10)Hf;
about Al-M-(10-50 vol. %)CG-(0.1-4)Sc-(0.05-5)Nb;
about Al-M-(10-50 vol. %)CG-(0.1-20)Er-(0.05-5)Nb;
about Al-M-(10-50 vol. %)CG-(0.1-15)Tm-(0.05-5)Nb;
about Al-M-(10-50 vol. %)CG-(0.1-25)Yb-(0.05-5)Nb; and
about Al-M-(10-50 vol. %)CG-(0.1-25)Lu-(0.05-5)Nb.
M is at least one of about (4-25) weight percent silicon, (1-8) weight percent magnesium, (0.5-3) weight percent lithium, (0.2-3) weight percent copper, (3-12) weight percent zinc, and (1-12) weight percent nickel. CG is a coarse grain ductile aluminum alloy having a particle size of about 25 to about 250 microns.
The amount of silicon present in the fine grain matrix of this invention if any may vary from about 4 to about 25 weight percent, more preferably from about 4 to about 18 weight percent, and even more preferably from about 5 to about 11 weight percent.
The amount of magnesium present in the fine grain matrix of this invention if any may vary from about 1 to about 8 weight percent, more preferably from about 3 to about 7.5 weight percent, and even more preferably from about 4 to about 6.5 weight percent.
The amount of lithium present in the fine grain matrix of this invention if any may vary from about 0.5 to about 3 weight percent, more preferably from about 1 to about 2.5 weight percent, and even more preferably from about 1 to about 2 weight percent.
The amount of copper present in the fine grain matrix of this invention if any may vary from about 0.2 to about 3 weight percent, more preferably from about 0.5 to about 2.5 weight percent, and even more preferably from about 1 to about 2.5 weight percent.
The amount of zinc present in the fine grain matrix of this invention if any may vary from about 3 to about 12 weight percent, more preferably from about 4 to about 10 weight percent, and even more preferably from about 5 to about 9 weight percent.
The amount of nickel present in the fine grain matrix of this invention if any may vary from about 1 to about 12 weight percent, more preferably from about 2 to about 10 weight percent, and even more preferably from about 4 to about 10 weight percent.
The amount of scandium present in the fine grain matrix of this invention if any may vary from 0.1 to about 4 weight percent, more preferably from about 0.1 to about 3 weight percent, and even more preferably from about 0.2 to about 2.5 weight percent. The Al—Sc phase diagram shown in FIG. 1 indicates a eutectic reaction at about 0.5 weight percent scandium at about 1219° F. (659° C.) resulting in a solid solution of scandium and aluminum and Al3Sc dispersoids. Aluminum alloys with less than 0.5 weight percent scandium can be quenched from the melt to retain scandium in solid solution that may precipitate as dispersed L12 intermetallic Al3Sc following an aging treatment. Alloys with scandium in excess of the eutectic composition (hypereutectic alloys) can only retain scandium in solid solution by rapid solidification processing (RSP) where cooling rates are in excess of about 103° C./second.
The amount of erbium present in the fine grain matrix of this invention, if any, may vary from about 0.1 to about 20 weight percent, more preferably from about 0.3 to about 15 weight percent, and even more preferably from about 0.5 to about 10 weight percent. The Al—Er phase diagram shown in FIG. 2 indicates a eutectic reaction at about 6 weight percent erbium at about 1211° F. (655° C.). Aluminum alloys with less than about 6 weight percent erbium can be quenched from the melt to retain erbium in solid solutions that may precipitate as dispersed L12 intermetallic Al3Er following an aging treatment. Alloys with erbium in excess of the eutectic composition can only retain erbium in solid solution by rapid solidification processing (RSP) where cooling rates are in excess of about 103° C./second.
The amount of thulium present in the alloys of this invention, if any, may vary from about 0.1 to about 15 weight percent, more preferably from about 0.2 to about 10 weight percent, and even more preferably from about 0.4 to about 6 weight percent. The Al—Tm phase diagram shown in FIG. 3 indicates a eutectic reaction at about 10 weight percent thulium at about 1193° F. (645° C.). Thulium forms metastable Al3Tm dispersoids in the aluminum matrix that have an L12 structure in the equilibrium condition. The Al3Tm dispersoids have a low diffusion coefficient which makes them thermally stable and highly resistant to coarsening. Aluminum alloys with less than 10 weight percent thulium can be quenched from the melt to retain thulium in solid solution that may precipitate as dispersed metastable L12 intermetallic Al3Tm following an aging treatment. Alloys with thulium in excess of the eutectic composition can only retain Tm in solid solution by rapid solidification processing (RSP) where cooling rates are in excess of about 103° C./second.
The amount of ytterbium present in the alloys of this invention, if any, may vary from about 0.1 to about 25 weight percent, more preferably from about 0.3 to about 20 weight percent, and even more preferably from about 0.4 to about 10 weight percent. The Al—Yb phase diagram shown in FIG. 4 indicates a eutectic reaction at about 21 weight percent ytterbium at about 1157° F. (625° C.). Aluminum alloys with less than about 21 weight percent ytterbium can be quenched from the melt to retain ytterbium in solid solution that may precipitate as dispersed L12 intermetallic Al3Yb following an aging treatment. Alloys with ytterbium in excess of the eutectic composition can only retain ytterbium in solid solution by rapid solidification processing (RSP) where cooling rates are in excess of about 103° C. per second.
The amount of lutetium present in the alloys of this invention, if any, may vary from about 0.1 to about 25 weight percent, more preferably from about 0.3 to about 20 weight percent, and even more preferably from about 0.4 to about 10 weight percent. The Al—Lu phase diagram shown in FIG. 5 indicates a eutectic reaction at about 11.7 weight percent Lu at about 1202° F. (650° C.). Aluminum alloys with less than about 11.7 weight percent lutetium can be quenched from the melt to retain Lu in solid solution that may precipitate as dispersed L12 intermetallic Al3Lu following an aging treatment. Alloys with Lu in excess of the eutectic composition can only retain Lu in solid solution by rapid solidification processing (RSP) where cooling rates are in excess of about 103° C./second. The amount of gadolinium present in the alloys of this invention, if any, may vary from about 0.1 to about 20 weight percent, more preferably from about 0.3 to about 15 weight percent, and even more preferably from about 0.5 to about 10 weight percent.
The amount of gadolinium present in the alloys of this invention, if any, may vary from about 0.1 to about 20 weight percent, more preferably from about 0.3 to about 15 weight percent, and even more preferably from about 0.5 to about 10 weight percent.
The amount of yttrium present in the alloys of this invention, if any, may vary from about 0.1 to about 20 weight percent, more preferably from about 0.3 to about 15 weight percent, and even more preferably from about 0.5 to about 10 weight percent.
The amount of zirconium present in the alloys of this invention, if any, may vary from about 0.05 to about 4 weight percent, more preferably from about 0.1 to about 3 weight percent, and even more preferably from about 0.3 to about 2 weight percent.
The amount of titanium present in the alloys of this invention, if any, may vary from about 0.05 to about 10 weight percent, more preferably from about 0.2 to about 8 weight percent, and even more preferably from about 0.4 to about 4 weight percent.
The amount of hafnium present in the alloys of this invention, if any, may vary from about 0.05 to about 10 weight percent, more preferably from about 0.2 to about 8 weight percent, and even more preferably from about 0.4 to about 5 weight percent.
The amount of niobium present in the alloys of this invention, if any, may vary from about 0.05 to about 5 weight percent, more preferably from about 0.1 to about 3 weight percent, and even more preferably from about 0.2 to about 2 weight percent.
In order to have the best properties for the fine grain matrix of this invention, it is desirable to limit the amount of other elements. Specific elements that should be reduced or eliminated include no more than about 0.1 weight percent iron, 0.1 weight percent chromium, 0.1 weight percent manganese, 0.1 weight percent vanadium, and 0.1 weight percent cobalt. The total quantity of additional elements should not exceed about 1% by weight, including the above listed impurities and other elements.
More preferred exemplary aluminum alloys of the bimodal system alloys of this invention include, but are not limited to (in weight percent unless otherwise specified):
about Al-M-(15-40 vol. %)CG-(0.1-3)Sc-(0.3-15)Gd;
about Al-M-(15-40 vol. %)CG-(0.3-15)Er-(0.3-15)Gd;
about Al-M-(15-40 vol. %)CG-(0.2-10)Tm-(0.3-15)Gd;
about Al-M-(15-40 vol. %)CG-(0.3-20)Yb-(0.3-15)Gd;
about Al-M-(15-40 vol. %)CG-(0.3-20)Lu-(0.3-15)Gd;
about Al-M-(15-40 vol. %)CG-(0.1-3)Sc-(0.3-15)Y;
about Al-M-(15-40 vol. %)CG-(0.3-15)Er-(0.3-15)Y;
about Al-M-(15-40 vol. %)CG-(0.2-10)Tm-(0.3-15)Y;
about Al-M-(15-40 vol. %)CG-(0.3-20)Yb-(0.3-15)Y;
about Al-M-(15-40 vol. %)CG-(0.3-20)Lu-(0.3-15)Y;
about Al-M-(15-40 vol. %)CG-(0.1-3)Sc-(0.1-3)Zr;
about Al-M-(15-40 vol. %)CG-(0.3-15)Er-(0.1-3)Zr;
about Al-M-(15-40 vol. %)CG-(0.2-10)Tm-(0.1-3)Zr;
about Al-M-(15-40 vol. %)CG-(0.3-20)Yb-(0.1-3)Zr;
about Al-M-(15-40 vol. %)CG-(0.3-20)Lu-(0.1-3)Zr;
about Al-M-(15-40 vol. %)CG-(0.1-3)Sc-(0.2-8)Ti;
about Al-M-(15-40 vol. %)CG-(0.3-15)Er-(0.2-8)Ti;
about Al-M-(15-40 vol. %)CG-(0.2-10)Tm-(0.2-8)Ti;
about Al-M-(15-40 vol. %)CG-(0.3-20)Yb-(0.2-8)Ti;
about Al-M-(15-40 vol. %)CG-(0.3-20)Lu-(0.2-8)Ti;
about Al-M-(15-40 vol. %)CG-(0.1-3)Sc-(0.2-8)Hf;
about Al-M-(15-40 vol. %)CG-(0.3-15)Er-(0.2-8)Hf;
about Al-M-(15-40 vol. %)CG-(0.2-10)Tm-(0.2-8)Hf;
about Al-M-(15-40 vol. %)CG-(0.3-20)Yb-(0.2-8)Hf;
about Al-M-(15-40 vol. %)CG-(0.3-20)Lu-(0.2-8)Hf;
about Al-M-(15-40 vol. %)CG-(0.1-3)Sc-(0.1-3)Nb;
about Al-M-(15-40 vol. %)CG-(0.3-15)Er-(0.1-3)Nb;
about Al-M-(15-40 vol. %)CG-(0.2-10)Tm-(0.1-3)Nb;
about Al-M-(15-40 vol. %)CG-(0.3-20)Yb-(0.1-3)Nb; and
about Al-M-(15-40 vol. %)CG-(0.3-20)Lu-(0.1-3)Nb.
M is at least one of about (4-18) weight percent silicon, (3-7.5) weight percent magnesium, (1-2.5) weight percent lithium, (0.5-2.5) weight percent copper, (4-10) weight percent zinc, and (2-10) weight percent nickel. CG is a coarse grain ductile aluminum alloy having a particle size of about 50 to about 200 microns.
Even more preferred exemplary aluminum alloys of the bimodal system alloys of this invention include, but are not limited to (in weight percent unless otherwise specified):
about Al-M-(20-30 vol. %)CG-(0.2-2.5)Sc-(0.5-10)Gd;
about Al-M-(20-30 vol. %)CG-(0.5-10)Er-(0.5-10)Gd;
about Al-M-(20-30 vol. %)CG-(0.4-6)Tm-(0.5-10)Gd;
about Al-M-(20-30 vol. %)CG-(0.4-10)Yb-(0.5-10)Gd;
about Al-M-(20-30 vol. %)CG-(0.4-10)Lu-(0.5-10)Gd;
about Al-M-(20-30 vol. %)CG-(0.2-2.5)Sc-(0.5-10)Y;
about Al-M-(20-30 vol. %)CG-(0.5-10)Er-(0.5-10)Y;
about Al-M-(20-30 vol. %)CG-(0.4-6)Tm-(0.5-10)Y;
about Al-M-(20-30 vol. %)CG-(0.4-10)Yb-(0.5-10)Y;
about Al-M-(20-30 vol. %)CG-(0.4-10)Lu-(0.5-10)Y;
about Al-M-(20-30 vol. %)CG-(0.2-2.5)Sc-(0.3-2)Zr;
about Al-M-(20-30 vol. %)CG-(0.5-10)Er-(0.3-2)Zr;
about Al-M-(20-30 vol. %)CG-(0.4-6)Tm-(0.3-2)Zr;
about Al-M-(20-30 vol. %)CG-(0.4-10)Yb-(0.3-2)Zr;
about Al-M-(20-30 vol. %)CG-(0.4-10)Lu-(0.3-2)Zr;
about Al-M-(20-30 vol. %)CG-(0.2-25)Sc-(0.4-5)Ti;
about Al-M-(20-30 vol. %)CG-(0.5-10)Er-(0.4-5)Ti;
about Al-M-(20-30 vol. %)CG-(0.4-6)Tm-(0.4-5)Ti;
about Al-M-(20-30 vol. %)CG-(0.4-10)Yb-(0.4-5)Ti;
about Al-M-(20-30 vol. %)CG-(0.4-10)Lu-(0.4-5)Ti;
about Al-M-(20-30 vol. %)CG-(0.2-2.5)Sc-(0.4-5)Hf;
about Al-M-(20-30 vol. %)CG-(0.5-10)Er-(0.4-5)Hf;
about Al-M-(20-30 vol. %)CG-(0.4-6)Tm-(0.4-5)Hf;
about Al-M-(20-30 vol. %)CG-(0.4-10)Yb-(0.4-5)Hf;
about Al-M-(20-30 vol. %)CG-(0.4-10)Lu-(0.4-5)Hf;
about Al-M-(20-30 vol. %)CG-(0.2-2.5)Sc-(0.2-2)Nb;
about Al-M-(20-30 vol. %)CG-(0.5-10)Er-(0.2-2)Nb;
about Al-M-(20-30 vol. %)CG-(0.4-6)Tm-(0.2-2)Nb;
about Al-M-(20-30 vol. %)CG-(0.4-10)Yb-(0.2-2)Nb; and
about Al-M-(20-30 vol. %)CG-(0.4-10)Lu-(0.2-2)Nb.
M is at least one of about (5-11) weight percent silicon, (4-6.5) weight percent magnesium, (1-2) weight percent lithium, (1-2.5) weight percent copper, (5-9) weight percent zinc, and (4-10) weight percent nickel. CG is a coarse grain ductile aluminum alloy having a particle size of about 100 to 150 microns.
Exemplary aluminum alloys of the trimodal system alloys of this invention include, but are not limited to (in weight percent unless otherwise specified):
about Al-M-(10-50 vol. %)CG-(0.1-4)Sc-(0.1-20)Gd-(5-40 vol. %)Al2O3;
about Al-M-(10-50 vol. %)CG-(0.1-20)Er-(0.1-20)Gd-(5-40 vol. %)Al2O3;
about Al-M-(10-50 vol. %)CG-(0.1-15)Tm-(0.1-20)Gd-(5-40 vol. %)Al2O3;
about Al-M-(10-50 vol. %)CG-(0.1-25)Yb-(0.1-20)Gd-(5-40 vol. %)Al2O3;
about Al-M-(10-50 vol. %)CG-(0.1-25)Lu-(0.1-20)Gd-(5-40 vol. %) Al2O3;
about Al-M-(10-50 vol. %)CG-(0.1-4)Sc-(0.1-20)Y-(5-40 vol. %)B4C;
about Al-M-(10-50 vol. %)CG-(0.1-20)Er-(0.1-20)Y-(5-40 vol. %)B4C;
about Al-M-(10-50 vol. %)CG-(0.1-15)Tm-(0.1-20)Y-(5-40 vol. %)B4C;
about Al-M-(10-50 vol. %)CG-(0.1-25)Yb-(0.1-20)Y-(5-40 vol. %)B4C
about Al-M-(10-50 vol. %)CG-(0.1-25)Lu-(0.1-20)Y-(5-40 vol. %)B4C;
about Al-M-(10-50 vol. %)CG-(0.1-4)Sc-(0.05-3.0)Zr-(5-40 vol. %)SiC;
about Al-M-(10-50 vol. %)CG-(0.1-20)Er-(0.05-4.0)Zr-(5-40 vol. %)SiC;
about Al-M-(10-50 vol. %)CG-(0.1-15)Tm-(0.05-4.0)Zr-(5-40 vol. %)SiC;
about Al-M-(10-50 vol. %)CG-(0.1-25)Yb-(0.05-4.0)Zr-(5-40 vol. %)SiC;
about Al-M-(10-50 vol. %)CG-(0.1-25)Lu-(0.05-4.0)Zr-(5-40 vol. %)SiC;
about Al-M-(10-50 vol. %)CG-(0.1-4)Sc-(0.05-10)Ti-(5-40 vol. %)TiB2;
about Al-M-(10-50 vol. %)CG-(0.1-20)Er-(0.05-10)Ti-(5-40 vol. %)TiB2;
about Al-M-(10-50 vol. %)CG-(0.1-15)Tm-(0.05-10)Ti-(5-40 vol. %)TiB2;
about Al-M-(10-50 vol. %)CG-(0.1-25)Yb-(0.05-10)Ti-(5-40 vol. %)TiB2;
about Al-M-(10-50 vol. %)CG-(0.1-25)Lu-(0.05-10)Ti-(5-40 vol. %)TiB2;
about Al-M-(10-50 vol. %)CG-(0.1-4)Sc-(0.05-10)Hf-(5-40 vol. %)TiB;
about Al-M-(10-50 vol. %)CG-(0.1-20)Er-(0.05-10)Hf-(5-40 vol. %)TiB;
about Al-M-(10-50 vol. %)CG-(0.1-15)Tm-(0.05-10)Hf-(5-40 vol. %)TiB;
about Al-M-(10-50 vol. %)CG-(0.1-25)Yb-(0.05-10)Hf-(5-40 vol. %)TiB;
about Al-M-(10-50 vol. %)CG-(0.1-25)Lu-(0.05-10)Hf-(5-40 vol. %)TiB;
about Al-M-(10-50 vol. %)CG-(0.1-4)Sc-(0.05-5)Nb-(5-40 vol. %)TiC;
about Al-M-(10-50 vol. %)CG-(0.1-20)Er-(0.05-5)Nb-(5-40 vol. %)TiC;
about Al-M-(10-50 vol. %)CG-(0.1-15)Tm-(0.05-5)Nb-(5-40 vol. %)TiC;
about Al-M-(10-50 vol. %)CG-(0.1-25)Yb-(0.05-5)Nb-(5-40 vol. %)TiC; and
about Al-M-(10-50 vol. %)CG-(0.1-25)Lu-(0.05-5)Nb-(5-40 vol. %)TiC.
M is at least one of about (4-25) silicon, (1-8) magnesium, (0.5-3) lithium, (0.2-3) copper, (3-12) zinc, and (1-12) nickel. CG is a coarse grain ductile aluminum alloy having a particle size of about 25 to about 250 microns.
The ceramic reinforcements added to create these trimodal system alloys provide additional strength and modulus enhancements which depend on particle size and volume fraction of the ceramic reinforcements.
These L12 bimodal and trimodal alloys may be made using standard powder metallurgy processing wherein the fine rapidly solidified matrix powder is mixed with the coarse ductile phase powder in an inert environment to prevent oxidation. The powder mix is then degassed and compacted in any suitable manner such as, for example, by vacuum hot pressing or blind die compaction (where compaction occurs in both by shear deformation) or by vacuum hot pressing (where compaction occurs by diffusion or creep). The fine grain matrix powder can be made by any rapid solidification technique that can provide elemental supersaturation such as, but not limited to, melt spinning, splat quenching, spray deposition, vacuum plasma spraying, cold spraying, laser melting, mechanical alloying, ball milling (i.e. at room temperature), cryomilling (i.e. in a liquid nitrogen environment), laser deposition, or atomization. Any processing technique utilizing cooling rates equivalent to or higher than about 103° C./second is considered to be a rapid solidification technique for these alloys. Therefore, the minimum desired cooling rate for the processing of these alloys is about 103° C./second, although higher cooling rates may be necessary for the fine grain matrix alloys having larger amounts of alloying additions.
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

1. An aluminum alloy with at least a bimodal microstructure comprising large grains in a fine grained matrix wherein:
the large grains comprise any aluminum alloy having a grain size ranging from about 25 to about 250 microns; and
the fine grain matrix having a grain size ranging from about 20 nm to about 5 microns comprises an aluminum alloy strengthened with L12 Al3X dispersoids wherein X comprises 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.

2. The aluminum alloy of claim 1, wherein the fine grain matrix comprises a solid solution of aluminum and at least one element selected from silicon, magnesium, lithium, copper, zinc, and nickel.

3. The aluminum alloy of claim 2, wherein the fine grain matrix comprises at least one of about 4 to about 25 weight percent silicon, about 1 to about 8 weight percent magnesium, about 0.5 to about 3 weight percent lithium, about 0.2 to about 3 weight percent copper, about 3 to about 12 weight percent zinc, about 1 to about 12 weight percent nickel.

4. The aluminum alloy of claim 1 with at least a trimodal microstructure, further comprising at least one ceramic reinforcement selected from aluminum oxide, silicon carbide, boron carbide, aluminum nitride, titanium boride, titanium diboride and titanium carbide.

5. The aluminum alloy of claim 4, wherein the ceramic reinforcement particle size ranges from about 0.5 to about 50 microns.

6. The aluminum alloy of claim 4, wherein the ceramic reinforcements comprise about 5 to about 40 volume percent of the alloy.

7. The aluminum alloy of claim 1, wherein the alloy is fabricated using powder metallurgical techniques.

8. The aluminum alloy of claim 7, wherein the fine grain matrix is produced by a rapid solidification technique utilizing a cooling rate of at least about 103° C./second, and a casting process.

9. The aluminum alloy of claim 7, wherein the rapid solidification technique comprises at least one of: melt spinning, splat quenching, atomization, spray deposition, vacuum plasma spraying, cold spraying, laser melting, mechanical alloying, cryomilling, spin forming and ball milling.

10. The aluminum alloy of claim 7, wherein the casting process comprises at least one of: squeeze casting, die casting, sand casting, and permanent mold casting.

11. The aluminum alloy of claim 1, wherein the large grains comprise about 10 to about 50 volume percent of the alloy.

12. The aluminum alloy of claim 1, wherein the Al3X dispersoids comprise:
at least one first element selected from the group consisting of about 0.1 to about 4 weight percent scandium, about 0.1 to about 20 weight percent erbium, about 0.1 to about 15 weight percent thulium, about 0.1 to about 25 weight percent ytterbium, and about 0.1 to about 25 weight percent lutetium; and
at least one second element selected from the group consisting of about 0.1 to about 20 weight percent gadolinium, about 0.1 to about 20 weight percent yttrium, about 0.05 to about 4 weight percent zirconium, about 0.05 to about 10 weight percent titanium, about 0.05 to about 10 weight percent hafnium, and about 0.05 to about 5 weight percent niobium.

13. The aluminum alloy of claim 1, wherein the alloy is capable of being used at temperatures from about −420° F. (−251° C.) up to about 650° F. (343° C.).