Abstract:
A method for producing high strength aluminum alloy tanks and other vessels containing L1 2  dispersoids from an aluminum alloy powder containing the L1 2  dispersoids. The powder is consolidated into a billet having a density of about 100 percent. Tanks are formed by rolling consolidated billets into sheets, cutting preforms from said sheets, roll forming the performs into cylindrical shapes and friction stir welding the seams to form cylinders. L1 2  alloy domes are spin formed from the rolled sheet and friction stir welded to the cylinder. Circular bases are cut from the rolled sheet and friction stir welded to the domed cylinder to form bottoms of the tank.

Description:
CROSS-REFERENCE TO RELATED APPLICATION(S) 
     This application is related to the following co-pending applications that were filed on Dec. 9, 2008 herewith and are assigned to the same assignee: CONVERSION PROCESS FOR HEAT TREATABLE L1 2  ALUMINUM ALLOYS, Ser. No. 12/316,020; A METHOD FOR FORMING HIGH STRENGTH ALUMINUM ALLOYS CONTAINING L1 2  INTERMETALLIC DISPERSOIDS, Ser. No. 12/316,046; and A METHOD FOR PRODUCING HIGH STRENGTH ALUMINUM ALLOY POWDER CONTAINING L1 2  INTERMETALLIC DISPERSOIDS, Ser. No. 12/316,047. 
     This application is also related to the following co-pending applications that were filed on Apr. 18, 2008, and are assigned to the same assignee: L1 2  ALUMINUM ALLOYS WITH BIMODAL AND TRIMODAL DISTRIBUTION, U.S. Pat. No. 8,409,373, Ser. No. 12/148,395; DISPERSION STRENGTHENED L1 2  ALUMINUM ALLOYS, U.S. Pat. No. 8,017,072, Ser. No. 12/148,432; HEAT TREATABLE L1 2  ALUMINUM ALLOYS, ABANDONED Ser. No. 12/148,383; HIGH STRENGTH L1 2  ALUMINUM ALLOYS, U.S. Pat. No. 7,871,477, Ser. No. 12/148,394; HIGH STRENGTH L1 2  ALUMINUM ALLOYS, U.S. Pat. No. 811,395, Ser. No. 12/148,382; HEAT TREATABLE L1 2  ALUMINUM ALLOYS, U.S. Pat. No. 7,875,133, Ser. No. 12/148,396; HIGH STRENGTH L1 2  ALUMINUM ALLOYS, US20090263273A1, Ser. No. 12/148,387; HIGH STRENGTH ALUMINUM ALLOYS WITH L1 2  PRECIPITATES, U.S. Pat. No. 7,879,162, Ser. No. 12/148,426; HIGH STRENGTH L1 2  ALUMINUM ALLOYS, U.S. Pat. No. 8,002,912, Ser. No. 12/148,459; and L1 2  STRENGTHENED AMORPHOUS ALUMINUM ALLOYS, U.S. Pat. No. 7,875,131, Ser. No. 12/148,458. 
     BACKGROUND 
     The present invention relates generally to aluminum alloys and more specifically to a method for forming high strength aluminum alloy powder having L1 2  dispersoids therein into aluminum parts such as tanks, containers and other components of turbine engines as well as other products fabricated from aluminum alloys. 
     The combination of high strength, ductility, and fracture toughness, as well as low density, make aluminum alloys natural candidates for as variety of applications. Because of its low weight high strength, ductility and fracture toughness, aluminum alloys are of interest in the manufacture and use for many applications. 
     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 Al 3 X L1 2  intermetallic phases where X is selected from the group consisting of Sc, Er, Lu, Yb, Tm, and Lu. The Al 3 X particles are coherent with the aluminum alloy matrix and are resistant to coarsening at elevated temperatures. The improved mechanical properties of the disclosed dispersion strengthened L1 2  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 L1 2  dispersoids. 
     L1 2  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 higher ductility. 
     SUMMARY 
     The present invention is a method for consolidating aluminum alloy powders into useful components such as tanks and other vessels having improved strength and fracture toughness. In embodiments, powders include an aluminum alloy having coherent L1 2  Al 3 X dispersoids where X is at least one first element selected from scandium, erbium, thulium, ytterbium, and lutetium, and at least one second element selected from gadolinium, yttrium, zirconium, titanium, hafnium, and niobium. The balance is substantially aluminum containing at least one alloying element selected from silicon, magnesium, manganese, lithium, copper, zinc, and nickel. 
     In one embodiment the L1 2  aluminum alloy tanks and other vessels are formed by rolling consolidated billets into sheets, cutting preforms from said sheets, roll forming the preforms into a cylindrical shape and friction stir welding the seam to form a cylinder. A spin formed L1 2  alloy dome and a flat bottom are then friction stir welded to the cylinder to form a tank. 
    
    
     
       BRIEF DESCRIPTION OF THE DRAWINGS 
         FIG. 1  is an aluminum scandium phase diagram. 
         FIG. 2  is an aluminum erbium phase diagram. 
         FIG. 3  is an aluminum thulium phase diagram. 
         FIG. 4  is an aluminum ytterbium phase diagram. 
         FIG. 5  is an aluminum lutetium phase diagram. 
         FIG. 6A  is a schematic diagram of a vertical gas atomizer. 
         FIG. 6B  is a close up view of nozzle  108  in  FIG. 6A . 
         FIGS. 7A and 7B  are SEM photos of the inventive aluminum alloy powder. 
         FIGS. 8A and 8B  are optical micrographs showing the microstructure of gas atomized L1 2  aluminum alloy powder. 
         FIG. 9  is a diagram showing the steps of the gas atomization process. 
         FIG. 10  is a diagram showing the processing steps to consolidate the L1 2  aluminum alloy powder. 
         FIG. 11  is a schematic diagram of blind die compaction. 
         FIG. 12  is a flow diagram of a method of forming an L1 2  aluminum alloy tank. 
         FIG. 13  is a photo of rolled L1 2  aluminum alloy sheet. 
         FIG. 14  is a schematic of a roll bending operation. 
         FIG. 15  is a perspective view of a friction stir welding operation. 
         FIG. 16  is a schematic of a spin forming operation. 
         FIG. 17  is a sketch of a cylindrical L1 2  aluminum alloy tank with a hemispherical top and a flat bottom. 
         FIG. 18  is a photograph of a cylindrical L1 2  aluminum alloy tank with a hemispherical top and a flat bottom. 
     
    
    
     DETAILED DESCRIPTION 
     1. L1 2  Aluminum Alloys 
     Alloy powders of this invention are formed from aluminum based alloys with high strength and fracture toughness for applications at temperatures from about −420° F. (−251° C.) up to about 650° F. (343° C.). The aluminum alloy comprises a solid solution of aluminum and at least one element selected from silicon, magnesium, lithium, copper, zinc, and nickel strengthened by L1 2  Al 3 X coherent precipitates where X is at least one first element selected from scandium, erbium, thulium, ytterbium, and lutetium, and at least one second element selected from gadolinium, yttrium, zirconium, titanium, hafnium, and niobium. 
     The alloys may also include at least one ceramic reinforcement. Aluminum oxide, silicon carbide, aluminum nitride, titanium diboride, titanium boride, boron carbide and titanium carbide are suitable ceramic reinforcements. 
     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 supersaturated 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 L1 2  intermetallic Al 3 Ni. 
     In the aluminum based alloys disclosed herein, scandium, erbium, thulium, ytterbium, and lutetium are potent strengtheners that have low diffusivity and low solubility in aluminum. All these elements form equilibrium Al 3 X intermetallic dispersoids where X is at least one of scandium, erbium, thulium, ytterbium, and 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. 
     Scandium forms Al 3 Sc dispersoids that are fine and coherent with the aluminum matrix. Lattice parameters of aluminum and Al 3 Sc are very close (0.405 nm and 0.410 nm respectively), indicating that there is minimal or no driving force for causing growth of the Al 3 Sc dispersoids. This low interfacial energy makes the Al 3 Sc dispersoids thermally stable and resistant to coarsening up to temperatures as high as about 842° 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 Al 3 Sc to coarsening. Additions of zinc, copper, lithium, silicon, manganese and nickel provide solid solution and precipitation strengthening in the aluminum alloys. 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. 
     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.). Additions of magnesium in aluminum increase the lattice parameter of the aluminum matrix, and decrease the lattice parameter mismatch further increasing the resistance of the Al 3 Er to coarsening. Additions of zinc, copper, lithium, silicon, manganese and nickel provide solid solution and precipitation strengthening in the aluminum alloys. These 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.). Additions of magnesium in aluminum increase the lattice parameter of the aluminum matrix, and decrease the lattice parameter mismatch further increasing the resistance of the Al 3 Tm to coarsening. Additions of zinc, copper, lithium, silicon, manganese and nickel provide solid solution and precipitation strengthening in the aluminum alloys. These 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.). Additions of magnesium in aluminum increase the lattice parameter of the aluminum matrix, and decrease the lattice parameter mismatch further increasing the resistance of the Al 3 Yb to coarsening. Additions of zinc, copper, lithium, silicon, manganese and nickel provide solid solution and precipitation strengthening in the aluminum alloys. These Al 3 Yb dispersoids are made stronger and more resistant to coarsening at elevated temperatures by adding suitable alloying elements such as gadolinium, yttrium, zirconium, titanium, hafnium, niobium, or combinations thereof that enter Al 3 Yb in solution. 
     Lutetium 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.). Additions of magnesium in aluminum increase the lattice parameter of the aluminum matrix, and decrease the lattice parameter mismatch further increasing the resistance of the Al 3 Lu to coarsening. Additions of zinc, copper, lithium, silicon, manganese and nickel provide solid solution and precipitation strengthening in the aluminum alloys. These 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 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. Despite its large atomic size, 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 despersoids have a low diffusion coefficient, which makes them thermally stable and highly resistant to coarsening. Titanium has a high solubility in the Al 3 X dispersoids allowing large amounts of titanium to substitute for X in the Al 3 X dispersoids, which result in improved thermal and structural stability. 
     Hafnium forms metastable Al 3 Hf dispersoids in the aluminum matrix that have an 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. 
     The aluminum oxide, silicon carbide, aluminum nitride, titanium di-boride, titanium boride, boron carbide and titanium carbide locate at the grain boundary and within the grain boundary to restrict dislocations from going around particles of the ceramic particles when the alloy is under stress. When dislocations form, they become attached with the ceramic particles on the departure side. Thus, more energy is required to detach the dislocation and the alloy has increased strength. To accomplish this, the particles of ceramic have to have a fine size, a moderate volume fraction in the alloy, and form a good interface between the matrix and the reinforcement. A working range of particle sizes is from about 0.5 to about 50 microns, more preferably about 1 to about 20 microns, and even more preferably about 1 to about 10 microns. The ceramic particles can break during blending and the average particle size will decrease as a result. 
     Al 3 X L1 2  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 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. 
     L1 2  phase strengthened aluminum alloys are important structural materials because of their excellent mechanical properties and the stability of these properties at elevated temperature due to the resistance of the coherent dispersoids in the microstructure to particle coarsening. The mechanical properties are optimized by maintaining a high volume fraction of L1 2  dispersoids in the microstructure. The L1 2  dispersoid concentration following aging scales as the amount of L1 2  phase forming elements in solid solution in the aluminum alloy following quenching. Examples of L1 2  phase forming elements include but are not limited to Sc, Er, Th, Yb, and Lu. The concentration of alloying elements in solid solution in alloys cooled from the melt is directly proportional to the cooling rate. 
     Exemplary aluminum alloys for this invention include, but are not limited to (in weight percent unless otherwise specified): 
     about Al-M-(0.1-4)Sc-(0.1-20)Gd; 
     about Al-M-(0.1-20)Er-(0.1-20)Gd; 
     about Al-M-(0.1-15)Tm-(0.1-20)Gd; 
     about Al-M-(0.1-25)Yb-(0.1-20)Gd; 
     about Al-M-(0.1-25)Lu-(0.1-20)Gd; 
     about Al-M-(0.1-4)Sc-(0.1-20)Y; 
     about Al-M-(0.1-20)Er-(0.1-20)Y; 
     about Al-M-(0.1-15)Tm-(0.1-20)Y; 
     about Al-M-(0.1-25)Yb-(0.1-20)Y; 
     about Al-M-(0.1-25)Lu-(0.1-20)Y; 
     about Al-M-(0.1-4)Sc-(0.05-4)Zr; 
     about Al-M-(0.1-20)Er-(0.05-4)Zr; 
     about Al-M-(0.1-15)Tm-(0.05-4)Zr; 
     about Al-M-(0.1-25)Yb-(0.05-4)Zr; 
     about Al-M-(0.1-25)Lu-(0.05-4)Zr; 
     about Al-M-(0.1-4)Sc-(0.05-10)Ti; 
     about Al-M-(0.1-20)Er-(0.05-10)Ti; 
     about Al-M-(0.1-15)Tm-(0.05-10)Ti; 
     about Al-M-(0.1-25)Yb-(0.05-10)Ti; 
     about Al-M-(0.1-25)Lu-(0.05-10)Ti; 
     about Al-M-(0.1-4)Sc-(0.05-10)Hf; 
     about Al-M-(0.1-20)Er-(0.05-10)Hf; 
     about Al-M-(0.1-15)Tm-(0.05-10)Hf; 
     about Al-M-(0.1-25)Yb-(0.05-10)Hf; 
     about Al-M-(0.1-25)Lu-(0.05-10)Hf; 
     about Al-M-(0.1-4)Sc-(0.05-5)Nb; 
     about Al-M-(0.1-20)Er-(0.05-5)Nb; 
     about Al-M-(0.1-15)Tm-(0.05-5)Nb; 
     about Al-M-(0.1-25)Yb-(0.05-5)Nb; and 
     about Al-M-(0.1-25)Lu-(0.05-5)Nb. 
     M is at least one of about (1-8) weight percent magnesium, about (4-25) weight percent silicon, about (0.1-3) weight percent manganese, about (0.5-3) weight percent lithium, about (0.2-6) weight percent copper, about (3-12) weight percent zinc, and about (1-12) weight percent nickel. 
     The amount of magnesium present in the fine grain matrix, 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 binary aluminum silicon system is a simple eutectic at 12.6 weight percent silicon and 1070.6° F. (577° C.). There is complete solubility of silicon and aluminum in the rapidly solidified inventive alloys discussed herein 
     The binary aluminum manganese system is a simple eutectic at about 2 weight percent manganese and 1216.4° F. (658° C.). There is complete solubility of manganese and aluminum in the rapidly solidified inventive alloys discussed herein 
     The amount of lithium present in the fine grain matrix, 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, if any, may vary from about 0.2 to about 6 weight percent, more preferably from about 0.5 to about 5 weight percent, and even more preferably from about 2 to about 4.5 weight percent. 
     The amount of zinc present in the fine grain matrix, 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, 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, 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 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. 
     The amount of erbium present in the fine grain matrix, 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 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 103° C./second. 
     The amount of thulium present in the alloys, if any, may vary from about 0.1 to about 15 weight percent, more preferably from about 0.2 to about 10 weight percent, and even more preferably from about 0.4 to about 6 weight percent. 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 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 103° C./second. 
     The amount of ytterbium present in the alloys, 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 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 103° C./second. 
     The amount of lutetium present in the alloys, 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 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 103° C./second. 
     The amount of gadolinium present in the alloys, 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, 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, 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, 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, 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, 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, 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 vanadium, and 0.1 weight percent cobalt. The total quantity of additional elements should not exceed about 1% by weight, including the above listed impurities and other elements. 
     2. L1 2  Alloy Powder Formation and Consolidation 
     The highest cooling rates observed in commercially viable processes are achieved by gas atomization of molten metals to produce powder. Gas atomization is a two fluid process wherein a stream of molten metal is disintegrated by a high velocity gas stream. The end result is that the particles of molten metal eventually become spherical due to surface tension and finely solidify in powder form. Heat from the liquid droplets is transferred to the atomization gas by convection. The solidification rates, depending on the gas and the surrounding environment, can be very high and can exceed 10 6 ° C./second. Cooling rates greater than 10 3 ° C./second are typically specified to ensure supersaturation of alloying elements in gas atomized L1 2  aluminum alloy powder in the inventive process described herein. 
     A schematic of typical vertical gas atomizer  100  is shown in  FIG. 6A .  FIG. 6A  is taken from R. Germain, Powder Metallurgy Science Second Edition MPIF (1994) (chapter 3, p. 101) and is included herein for reference. Vacuum or inert gas induction melter  102  is positioned at the top of free flight chamber  104 . Vacuum induction melter  102  contains melt  106  which flows by gravity or gas overpressure through nozzle  108 . A close up view of nozzle  108  is shown in  FIG. 6B . Melt  106  enters nozzle  108  and flows downward till it meets the high pressure gas stream from gas source  110  where it is transformed into a spray of droplets. The droplets eventually become spherical due to surface tension and rapidly solidify into spherical powder  112  which collects in collection chamber  114 . The gas recirculates through cyclone collector  116  which collects fine powder  118  before returning to the input gas stream. As can be seen from  FIG. 6A , the surroundings to which the melt and eventual powder are exposed are completely controlled. 
     There are many effective nozzle designs known in the art to produce spherical metal powder. Designs with short gas-to-melt separation distances produce finer powders. Confined nozzle designs where gas meets the molten stream at a short distance just after it leaves the atomization nozzle are preferred for the production of the inventive L1 2  aluminum alloy powders disclosed herein. Higher superheat temperatures cause lower melt viscosity and longer cooling times. Both result in smaller spherical particles. 
     A large number of processing parameters are associated with gas atomization that affect the final product. Examples include melt superheat, gas pressure, metal flow rate, gas type, and gas purity. In gas atomization, the particle size is related to the energy input to the metal. Higher gas pressures, higher superheat temperatures and lower metal flow rates result in smaller particle sizes. Higher gas pressures provide higher gas velocities for a given atomization nozzle design. 
     To maintain purity, inert gases are used, such as helium, argon, and nitrogen. Helium is preferred for rapid solidification because the high heat transfer coefficient of the gas leads to high quenching rates and high supersaturation of alloying elements. 
     Lower metal flow rates and higher gas flow ratios favor production of finer powders. The particle size of gas atomized melts typically has a log normal distribution. In the turbulent conditions existing at the gas/metal interface during atomization, ultra fine particles can form that may reenter the gas expansion zone. These solidified fine particles can be carried into the flight path of molten larger droplets resulting in agglomeration of small satellite particles on the surfaces of larger particles. An example of small satellite particles attached to inventive spherical L1 2  aluminum alloy powder is shown in the scanning electron microscopy (SEM) micrographs of  FIGS. 7A and 7B  at two magnifications. The spherical shape of gas atomized aluminum powder is evident. The spherical shape of the powder is suggestive of clean powder without excessive oxidation. Higher oxygen in the powder results in irregular powder shape. Spherical powder helps in improving the flowability of powder which results in higher apparent density and tap density of the powder. The satellite particles can be minimized by adjusting processing parameters to reduce or even eliminate turbulence in the gas atomization process. The microstructure of gas atomized aluminum alloy powder is predominantly cellular as shown in the optical micrographs of cross-sections of the inventive alloy in  FIGS. 8A and 8B  at two magnifications. The rapid cooling rate suppresses dendritic solidification common at slower cooling rates resulting in a finer microstructure with minimum alloy segregation. 
     Oxygen and hydrogen in the powder can degrade the mechanical properties of the final part. It is preferred to limit the oxygen in the L1 2  alloy powder to about 1 ppm to 2000 ppm. Oxygen is intentionally introduced as a component of the helium gas during atomization. An oxide coating on the L1 2  aluminum powder is beneficial for two reasons. First, the coating prevents agglomeration by contact sintering and secondly, the coating inhibits the chance of explosion of the powder. A controlled amount of oxygen is important in order to provide good ductility and fracture toughness in the final consolidated material. Hydrogen content in the powder is controlled by ensuring the dew point of the helium gas is low. A dew point of about minus 50° F. (minus 45.5° C.) to minus 100° F. (minus 73.3° C.) is preferred. 
     In preparation for final processing, the powder is classified according to size by sieving. To prepare the powder for sieving, if the powder has zero percent oxygen content, the powder may be exposed to nitrogen gas which passivates the powder surface and prevents agglomeration. Finer powder sizes result in improved mechanical properties of the end product. While minus 325 mesh (about 45 microns) powder can be used, minus 450 mesh (about 30 microns) powder is a preferred size in order to provide good mechanical properties in the end product. During the atomization process, powder is collected in collection chambers in order to prevent oxidation of the powder. Collection chambers are used at the bottom of atomization chamber  104  as well as at the bottom of cyclone collector  116 . The powder is transported and stored in the collection chambers also. Collection chambers are maintained under positive pressure with nitrogen gas which prevents oxidation of the powder. 
     A schematic of the L1 2  aluminum powder manufacturing process is shown in  FIG. 9 . In the process aluminum  200  and L1 2  forming (and other alloying) elements  210  are melted in furnace  220  to a predetermined superheat temperature under vacuum or inert atmosphere. Preferred charge for furnace  220  is prealloyed aluminum  200  and L1 2  and other alloying elements before charging furnace  220 . Melt  230  is then passed through nozzle  240  where it is impacted by pressurized gas stream  250 . Gas stream  250  is an inert gas such as nitrogen, argon or helium, preferably helium. Melt  230  can flow through nozzle  240  under gravity or under pressure. Gravity flow is preferred for the inventive process disclosed herein. Preferred pressures for pressurized gas stream  250  are about 50 psi (10.35 MPa) to about 750 psi (5.17 MPa) depending on the alloy. 
     The atomization process creates molten droplets  260  which rapidly solidify as they travel through agglomeration chamber  270  forming spherical powder particles  280 . The molten droplets transfer heat to the atomizing gas by convention. The role of the atomizing gas is two fold: one is to disintegrate the molten metal stream into fine droplets by transferring kinetic energy from the gas to the melt stream and the other is to extract heat from the molten droplets to rapidly solidify them into spherical powder. The solidification time and cooling rate vary with droplet size. Larger droplets take longer to solidify and their resulting cooling rate is lower. On the other hand, the atomizing gas will extract heat efficiently from smaller droplets resulting in a higher cooling rate. Finer powder size is therefore preferred as higher cooling rates provide finer microstructures and higher mechanical properties in the end product. Higher cooling rates lead to finer cellular microstructures which are preferred for higher mechanical properties. Finer cellular microstructures result in finer grain sizes in consolidated product. Finer grain size provides higher yield strength of the material through the Hall-Petch strengthening model. 
     Key process variables for gas atomization include superheat temperature, nozzle diameter, helium content and dew point of the gas, and metal flow rate. Superheat temperatures of from about 150° F. (66° C.) to 200° F. (93° C.) are preferred. Nozzle diameters of about 0.07 in. (1.8 mm) to 0.12 in. (3.0 mm) are preferred depending on the alloy. The gas stream used herein was a helium nitrogen mixture containing 74 to 87 vol. % helium. The metal flow rate ranged from about 0.8 lb/min (0.36 kg/min) to 4.0 lb/min (1.81 kg/min). The oxygen content of the L1 2  aluminum alloy powders was observed to consistently decrease as a run progressed. This is suggested to be the result of the oxygen gettering capability of the aluminum powder in a closed system. The dew point of the gas was controlled to minimize hydrogen content of the powder. Dew points in the gases used in the examples ranged from −10° F. (−23° C.) to −110° F. (−79° C.). 
     The powder is then classified by sieving process  290  to create classified powder  300 . Sieving of powder is performed under an inert environment to minimize oxygen and hydrogen pickup from the environment. While the yield of minus 450 mesh powder is extremely high (95%), there are always larger particle sizes, flakes and ligaments that are removed by the sieving. Sieving also ensures a narrow size distribution and provides a more uniform powder size. Sieving also ensures that flaw sizes cannot be greater than minus 450 mesh which will be required for nondestructive inspection of the final product. 
     Processing parameters of exemplary gas atomization runs are listed in Table 1. 
     
       
         
               
             
               
               
               
               
               
               
               
               
               
             
               
               
               
               
               
               
               
               
               
             
           
               
                 TABLE 1 
               
             
             
               
                   
               
               
                 Gas atomization parameters used for producing powder 
               
             
          
           
               
                   
                   
                   
                   
                   
                   
                 Average 
                   
                   
               
               
                   
                   
                   
                   
                   
                   
                 Metal 
                 Oxygen 
                 Oxygen 
               
               
                   
                 Nozzle 
                 He 
                 Gas 
                 Dew 
                 Charge 
                 Flow 
                 Content 
                 Content 
               
               
                   
                 Diameter 
                 Content 
                 Pressure 
                 Point 
                 Temperature 
                 Rate 
                 (ppm) 
                 (ppm) 
               
               
                 Run 
                 (in) 
                 (vol %) 
                 (psi) 
                 (° F.) 
                 (° F.) 
                 (lbs/min) 
                 Start 
                 End 
               
               
                   
               
             
          
           
               
                 1 
                 0.10 
                 79 
                 190 
                 &lt;−58 
                 2200 
                 2.8 
                 340 
                 35 
               
               
                 2 
                 0.10 
                 83 
                 192 
                 −35 
                 1635 
                 0.8 
                 772 
                 27 
               
               
                 3 
                 0.09 
                 78 
                 190 
                 −10 
                 2230 
                 1.4 
                 297 
                 &lt;0.01 
               
               
                 4 
                 0.09 
                 85 
                 160 
                 −38 
                 1845 
                 2.2 
                 22 
                 4.1 
               
               
                 5 
                 0.10 
                 86 
                 207 
                 −88 
                 1885 
                 3.3 
                 286 
                 208 
               
               
                 6 
                 0.09 
                 86 
                 207 
                 −92 
                 1915 
                 2.6 
                 145 
                 88 
               
               
                   
               
             
          
         
       
     
     The role of powder quality is extremely important to produce material with higher strength and ductility. Powder quality is determined by powder size, shape, size distribution, oxygen content, hydrogen content, and alloy chemistry. Over fifty gas atomization runs were performed to produce the inventive powder with finer powder size, finer size distribution, spherical shape, and lower oxygen and hydrogen contents. Processing parameters of some exemplary gas atomization runs are listed in Table 1. It is suggested that the observed decrease in oxygen content is attributed to oxygen gettering by the powder as the runs progressed. 
     Inventive L1 2  aluminum alloy powder was produced with over 95% yield of minus 450 mesh (30 microns) which includes powder from about 1 micron to about 30 microns. The average powder size was about 10 microns to about 15 microns. As noted above, finer powder size is preferred for higher mechanical properties. Finer powders have finer cellular microstructures. As a result, finer cell sizes lead to finer grain size by fragmentation and coalescence of cells during powder consolidation. Finer grain sizes produce higher yield strength through the Hall-Petch strengthening model where yield strength varies inversely as the square root of the grain size. It is preferred to use powder with an average particle size of 10-15 microns. Powders with a powder size less than 10-15 microns can be more challenging to handle due to the larger surface area of the powder. Powders with sizes larger than 10-15 microns will result in larger cell sizes in the consolidated product which, in turn, will lead to larger grain sizes and lower yield strengths. 
     Powders with narrow size distributions are preferred. Narrower powder size distributings produce product microstructures with more uniform grain size. Spherical powder was produced to provide higher apparent and tap densities which help in achieving 100% density in the consolidated product. Spherical shape is also an indication of cleaner and lower oxygen content powder. Lower oxygen and lower hydrogen contents are important in producing material with high ductility and fracture toughness. Although it is beneficial to maintain low oxygen and hydrogen content in powder to achieve good mechanical properties, lower oxygen may interfere with sieving due to self sintering. An oxygen content of about 25 ppm to about 500 ppm is preferred to provide good ductility and fracture toughness without any sieving issue. Lower hydrogen is also preferred for improving ductility and fracture toughness. It is preferred to have about 25-200 ppm of hydrogen in atomized powder by controlling the dew point in the atomization chamber. Hydrogen in the powder is further reduced by heating the powder in vacuum. Lower hydrogen in final product is preferred to achieve good ductility and fracture toughness. 
     A schematic of the L1 2  aluminum powder consolidation process is shown in  FIG. 10 . The starting material is sieved and classified L1 2  aluminum alloy powders (step  310 ). Blending (step  320 ) is a preferred step in the consolidation process because it results in improved uniformity of particle size distribution. Gas atomized L1 2  aluminum alloy powder generally exhibits a bimodal particle size distribution and cross blending of separate powder batches tends to homogenize the particle size distribution. Blending (step  320 ) is also preferred when separate metal and/or ceramic powders are added to the L1 2  base powder to form bimodal or trimodal consolidated alloy microstructures. 
     Following blending (step  320 ), the powders are transferred to a can (step  330 ) where the powder is vacuum degassed (step  340 ) at elevated temperatures. The can (step  330 ) is an aluminum container having a cylindrical, rectangular or other configuration with a central axis. Cylindrical configurations are preferred with hydraulic extrusion presses. Vacuum degassing times can range from about 0.5 hours to about 8 days. A temperature range of about 300° F. (149° C.) to about 900° F. (482° C.) is preferred. Dynamic degassing of large amounts of powder is preferred to static degassing. In dynamic degassing, the can is preferably rotated during degassing to expose all of the powder to a uniform temperature. Degassing removes oxygen and hydrogen from the powder. 
     Following vacuum degassing (step  340 ), the vacuum line is crimped and welded shut (step  350 ). The powder is then fully densified by blind die compaction or closed die forging as the process is sometimes called (step  360 ). At this point the can may be removed by machining (step  380 ) to form a useful billet (step  390 ). 
     A schematic showing blind die compaction (process  400 ) is shown in  FIGS. 11A and 11B . The equipment comprises base  410 , die  420 , ram  430 , and means to apply pressure to ram  430  indicated by arrow  450 . Prior to compaction, billet  440  does not fill die cavity  460 . After compaction, billet  445  completely fills the die cavity and has taken the shape of die cavity  460 . The die cavities can have any shape provided they have a central symmetrical axis parallel to arrow  450 . Rectangular shapes adopt well for rolling preforms. Canned L1 2  aluminum alloy powder preforms are easily densified due to the large capacity of modern hydraulic presses. 
     3. L1 2  Aluminum Alloy Container Fabrication 
     L1 2  aluminum alloys are ideal lightweight candidates for application as containers in an aerospace environment where temperatures approach 600° F. (316° C.). The high specific strength of these alloys qualifies them to be used as lightweight high pressure containers such as fuel or hydraulic tanks. Fabrication of one exemplary embodiment of the invention, a cylinder with a hemispherical top, will be described here as an example. It is to be understood that the inventive L1 2  alloys can be used for an unlimited number of container applications, particularly in the aerospace environment. 
     A flow diagram listing the steps to fabricate an L1 2  aluminum alloy cylinder with a hemispherical top is given in  FIG. 12 . The process starts with a consolidated L1 2  alloyed billet with the can removed (Step  510 ). The billet is then rolled into a thin sheet (Step  520 ). Rolling at ambient temperature with intermediate stress relief anneals between passes is preferred. A photograph of a rolled L1 2  aluminum alloy sheet is shown in  FIG. 13 . The sheet has an excellent surface finish. Rectangular blanks are then cut from the sheet that will form the cylindrical wall of the container (Step  520 ). The rectangular blanks are then roll formed into a cylindrical shape. There are many ways to roll form an alloy sheet into cylindrical shapes known in the art. One method is schematically illustrated in  FIG. 14 . Roll forming process  600  comprises rolled L1 2  alloy sheet  610  passing between removable mandrel  630  and urethane roll  640  mounted on drive shaft  650 . Top steel roll  620  is programmed to apply downward pressure to force mandrel  630  and alloy sheet  610  into urethane roll  640  as it is driven in the direction of arrow  650 . This allows alloy sheet  610  to assume the curvature of mandrel  630  thereby forming a cylindrical shape. 
     Roll formed cylindrical alloy sheet  610  is then friction stir welded to form a cylinder (Step  540 ).  FIG. 15  is a perspective sketch of friction stir welding operation  700 . Friction stir welding (FSW) operation  700  comprises welded structure  710  and FSW system  712 . Welded structure  710  comprises L1 2  alloy sheets  714  and  716  that abut each other at intersection  718 . In this invention, alloy sheets  714  and  716  are each end of roll formed cylinder  610 . As discussed below, alloy sheet  714  and  716  are welded together at intersection  718  with FSW system  712  to form welded joint  720  where joint  720  substantially retains the pre-weld strengths of alloy sheets  714  and  716 . 
     FSW system  712  includes controller  722 , tool  724 , and pin  726  (pin  726  shown with hidden lines). Pin  726  extends from the bottom surface of tool  724  and is pressed into metal parts  714  and  716  during a FSW operation. Controller  722  directs tool  724  and pin  726  to rotate in the direction of arrow  728  (or in an opposite rotational direction from arrow  728 ), and to press down into metal parts  714  and  716  in the direction of arrow  730 . This causes pin  726  to dig into alloy sheets  714  and  716  at intersection  718  until tool  724  reaches alloy sheets  714  and  716 . The depth of pin  726  determines the depth of the weld at intersection  718 . 
     While tool  724  and pin  726  are rotating, controller  722  directs tool  724  and pin  726  to move along intersection  718  in the direction of arrow  732 . As tool  724  and pin  726  move along intersection  718 , the rotation of tool  724  and pin  726  frictionally heat alloy sheets  714  and  716  at intersection  718 . The heated alloys enter a plastic-like state, and are stirred by the rotational motion of tool  724  and pin  726 , thereby creating welded joint  720  at intersection  718 . Alloy sheets  714  and  716  are desirably braced together to prevent alloy sheets  714  and  716  from moving apart during the FSW operation. Process variables include tool material, downward force, tool rotational speed, and traverse rate. The FSW operation is a solid-state welding process, in which the heated alloys do not melt. As such, the refined microstructures of the L1 2  aluminum alloys are substantially retained while forming welded joint  720 . This is in contrast to other welding techniques, such as fusion welding, in which the welded alloys are melted to form a welded joint. Melting L1 2  aluminum alloys destroys the refined L1 2  microstructure of the alloys, thereby lowering the strength and deformation resistance of the resulting welded structure. 
     The next step is to cut a circular preform from the L1 2  alloy rolled sheet as a starting piece to form a dome (Step  550 ). The dome is then formed by a number of sheet metal forming processes known in the art. The example discussed herein is spin forming (Step  560 ).  FIG. 16  is a schematic illustration of spin forming operation  800 . Spin forming operation  800  comprises L1 2  alloy sheet  810  on hemispherical dome pattern  820  on rotatable table  830  on rotatable shaft  840  rotating in direction of arrows  850  around axis  860 . 
     Spin forming tool  865  comprises tool holder  875  and tool  870 . Tool  870  comprises a rotatable wheel rotating in the direction of arrow  880 . During spin forming, tool  865  is moved in the direction of arrow  890  such that tool  870  forces alloy sheet  810  to conform to the external shape of hemispherical dome pattern  820  thereby spin forming a dome. 
     The next step in fabricating an L1 2  alloy container comprising a cylindrical body with a hemispherical dome top is to friction stir weld the spin formed dome to the roll formed and friction stir welded cylinder (Step  570 ). Finally, a closed container can be formed by cutting a circular bottom for the cylinder from L1 2  alloy rolled sheet (Step  580 ) and friction stir welding it to the domed L1 2  alloy cylinder. 
     A perspective sketch of the inventive L1 2  aluminum alloy container is shown in  FIG. 17 . Roll formed L1 2  alloy cylindrical body  910  is topped with spin formed hemispherical dome  930 . Friction stir weld  920  forms cylindrical body  910  and friction stir weld  940  attaches L1 2  alloy dome  930  to cylindrical body  910 . Bottom  950  is attached to cylindrical body  910  by friction stir weld  960 . A photograph of a friction stir welded L1 2  roll formed cylinder with a spin formed dome is shown in  FIG. 18  with the components identified by the same numbers as in  FIG. 17 . 
     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.