Patent Application
20100254850
The present invention relates generally to aluminum alloys and more specifically to a method for forming high strength aluminum alloy billets having L12 dispersoids therein.
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 owned by the assignee of the present invention 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 Al also owned commonly 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 commercially available aluminum alloys. Fine grain size results in improved mechanical properties of materials. Hall-Petch strengthening has been known for decades where strength increases as grain size decreases. An optimum grain size for optimum strength is in the nano range of about 30 to 100 nm. These alloys also have lower ductility.
The present invention is a method for forming aluminum alloy billets with high strength and acceptable fracture toughness. In embodiments, the alloys have 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 alloys are formed by encompassing a powder preform of an aluminum alloy body containing L12 dispersoid forming elements in a heated, flowable pressure transmitting medium, and rapidly compressing the powder to consolidate the perform to form the billet. Use of graphite or a ceramic as the pressure transfer medium causes a non-isostatic pressure field to form in the chamber. During consolidation, the powder preform undergoes an axial compression that exceeds radial expansion. The resulting shear strain cleans the surface oxide from the particles and increases metal-to-metal contact improving forging density.
This method is known in the industry as Ceracon forging. It has not been attempted using aluminum alloy containing L12 dispersoids. Examples of the Ceracon forging process are shown in U.S. Pat. No. 4,667,497, U.S. Patent Application Publication No. 2005/0147520 and U.S. Pat. No. 7,097,807 and are included in their entirety by reference.
The compression takes place in a closed container and results in billets of aluminum alloy containing L12 dispersoids with a density of essentially 100%.
The alloy products 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 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 can be complete solubility of lithium in the rapid solidified inventive alloys discussed herein.
The binary aluminum copper system is a simple eutectic at 32 weight percent copper and 1018° F. (548° C.). There can be complete solubility of copper in the rapidly solidified inventive alloys discussed herein.
The aluminum zinc binary system is a eutectic alloy system involving a monotectoid reaction and a miscibility gap in the solid state. There is a eutectic reaction at 94 weight percent zinc and 718° 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 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 nm respectively), indicating there is minimal driving force for causing growth of the Al3Yb dispersoids. This low interfacial energy makes the Al3Yb dispersoids thermally stable and resistant to coarsening up to temperatures as high as about 842° F. (450° C.). Additions of magnesium in 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 DO22 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.
L12 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 L12 dispersoids in the microstructure. The L12 dispersoid concentration following aging scales as the amount of L12 phase forming elements in solid solution in the aluminum alloy following quenching. Examples of L12 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 the bimodal system alloys of 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 (4-25) weight percent silicon, (1-8) weight percent magnesium, (0.5-3) weight percent lithium, (0.2-6.5) weight percent copper, (3-12) weight percent zinc, and (1-12) weight percent nickel.
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 6.5 weight percent, more preferably from about 0.5 to about 5.0 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 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
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
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
The amount of ytterbium present in the alloys of this invention, if any, may vary from about 0.1 to about 25 weight percent, more preferably from about 0.3 to about 20 weight percent, and even more preferably from about 0.4 to about 10 weight percent. The Al—Yb phase diagram shown in
The amount of lutetium present in the alloys of this invention, if any, may vary from about 0.1 to about 25 weight percent, more preferably from about 0.3 to about 20 weight percent, and even more preferably from about 0.4 to about 10 weight percent. The Al—Lu phase diagram shown in
The amount of gadolinium present in the alloys of this invention, if any, may vary from about 0.1 to about 20 weight percent, more preferably from about 0.3 to about 15 weight percent, and even more preferably from about 0.5 to about 10 weight percent.
The amount of yttrium present in the alloys of this invention, if any, may vary from about 0.1 to about 20 weight percent, more preferably from about 0.3 to about 15 weight percent, and even more preferably from about 0.5 to about 10 weight percent.
The amount of zirconium present in the alloys of this invention, if any, may vary from about 0.05 to about 4 weight percent, more preferably from about 0.1 to about 3 weight percent, and even more preferably from about 0.3 to about 2 weight percent.
The amount of titanium present in the alloys of this invention, if any, may vary from about 0.05 to about 10 weight percent, more preferably from about 0.2 to about 8 weight percent, and even more preferably from about 0.4 to about 4 weight percent.
The amount of hafnium present in the alloys of this invention, if any, may vary from about 0.05 to about 10 weight percent, more preferably from about 0.2 to about 8 weight percent, and even more preferably from about 0.4 to about 5 weight percent.
The amount of niobium present in the alloys of this invention, if any, may vary from about 0.05 to about 5 weight percent, more preferably from about 0.1 to about 3 weight percent, and even more preferably from about 0.2 to about 2 weight percent.
In order to have the best properties for the 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.
It is advantageous to form L12 strengthened aluminum alloy product from powder. The major reason is that the rapid cooling rate experienced during powder formation from the melt results in high supersaturation of intermetallic L12 phase forming elements in the powder. The high supersaturation leads to a maximum amount of the strengthening phase dispersed throughout the structure in the final consolidated part.
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 106° C./second. Cooling rates greater than 103° C./second are typically specified to ensure supersaturation of alloying elements in gas atomized L12 aluminum alloy powder in the inventive process described herein.
A schematic of typical vertical gas atomizer 100 is shown in
There are many effective nozzle designs known in the art to produce spherical metal powder. Nozzle 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 L12 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 rates favor production of finer powders. The particle size of gas atomized melts typically has a log normal distribution. An example of spherical L12 aluminum alloy powder is shown in the scanning electron micrograph (SEM) of
Oxygen and hydrogen in the powder can degrade the mechanical properties of the final part. It is preferred to limit the oxygen in the L12 alloy powder to about 1 ppm to 2000 ppm. Oxygen is intentionally introduced as a component of the helium gas during atomization. A thin oxide coating on the L12 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.
The process of consolidating the inventive alloy powders into useful forms is schematically illustrated in
The sieved, blended and (optionally) cryomilled powders are then put in a can (step 250) and vacuum degassed (step 260). Following vacuum degassing, the can is sealed (step 270) under vacuum and forged (step 280) to produce a densified preform. Finally, the preform is Ceracon forged (step 290) to produce a product with improved mechanical properties useful for subsequent service as a high temperature L12 strengthened aluminum alloy. Non-isostatic Ceracon forging will be described later.
Sieving (step 220) is a critical step in consolidation because the final mechanical properties relate directly to the particle size. Finer particle size results in finer L12 particle dispersion and finer grain size. Optimum mechanical properties have been observed with −450 mesh (30 micron) powder. Sieving (step 220) also limits the defect size in the powder. Before sieving, the powder is passivated with nitrogen gas in order to prevent agglomeration. Ultrasonic sieving is preferred for its efficiency.
Blending (step 230) is another critical step in the consolidation process because it results in improved uniformity of particle size distribution. Gas atomized L12 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 230) is also necessary when separate metal and/or ceramic powders are added to the L12 base powder to form bimodal and trimodal consolidated alloy microstructures.
Cryomilling (step 240) can be used to refine the grain size of gas atomized L12 aluminum alloy powder as well as the final consolidated alloy microstructure. Cryomilling is described in U.S. Pat. No. 6,902,699, Fritzemeier et al. and in U.S. Pat. No. 7,344,675, Van Daam et al. and are incorporated herein in their entirety by reference. Cryomilling involves high-energy ball milling under liquid nitrogen. The liquid nitrogen environment prevents oxidation and prevents frictional heating of the powder and the resulting grain coarsening. During the process, the powder particles are repeatedly sheared, fractured and cold welded which results in a severely deformed microstructure containing a high dislocation density that, with continued deformation, evolves into a cellular structure consisting of extremely small dislocation free grains separated by high angle grain boundaries with high dislocation density. The grain size of the cellular microstructure is typically less than 100 nm and the microstructure is considered a nanostructure.
In addition, the nitrogen environment results in the formation of nitride particles that reside at the grain boundaries and in the grain interiors and resist coarsening at higher temperatures. Stearic acid is preferably added to the powder charge to prevent excessive agglomeration and to promote fracturing and rewelding of the L12 aluminum alloy particles during milling.
Following sieving (step 220), blending (step 230) and (optionally) cryomilling (Step 240), the powders are transferred to a can (step 250) where the powder is vacuum degassed (step 260) for about 12 hours to over 8 days at elevated temperatures. A temperature range of about 500° F. (260° C.) to about 900° F. (482° C.) is preferred and about 750° F. (399° C.) is more preferred. Dynamic degassing large amounts of powder are preferred to static degassing to expose all of the powder to a uniform temperature. Degassing removes the stearic acid lubricant as well as oxygen and hydrogen from the charge.
Following vacuum degassing (step 260), the vacuum line is crimped and welded shut. The powder is then consolidated into a dense preform by closed die forging or by quasi-isostatic forging (step 280).
An exemplary embodiment of this invention is to consolidate a canned L12 aluminum alloy powder preform into a substantially 100% dense billet by a quasi-isostatic Ceracon-type forging process. The forging process consists of uniaxially pressing the canned powder preform or solid part perform in a cylindrical press, wherein the preform is surrounded by pressure transmitting medium during the forging. A schematic of quasi-isostatic forging equipment 300 is shown in
The steps to Ceracon forge an L12 aluminum alloy powder preform are schematically illustrated in
One advantage of this process is that, if the preheating steps are followed, the run time is short. As a result, deleterious microstructural changes, such as grain and particle coarsening are minimized. Run times can be as short as a few minutes if automated loading and charging equipment is used. In addition, ductility and fracture toughness of L12 based aluminum alloys can be improved significantly due to dynamic bimodal pressure generated during quasi-isostatic forging that breaks apart continuous powder surface oxide that prevent metal-metal bonding and randomly distributes them in the material.
During quasi-isostatic Ceracon type forging, the billet axially deforms about 30% and radially deforms about 10%. Strain rates from about 0.1 min−1 to about 6 min−1 at forging temperatures from about 400° F. (204° C.) to about 900° F. (482° C.) are preferred. Forging pressures from 50 ksi (345 MPa) to 150 ksi (1034 MPa) are preferred.
The pressure transmitting medium can consist of graphite or other carbon containing powders or ceramic powders or both. By proper fixturing and preheating the pressure transmitting medium and the work piece, densification during quasi-isostatic forging can be extremely rapid on the order of minutes.
The unequal actual and radial deformation of the powder is schematically illustrated by the dotted line outline of deformed can 352 in
Quasi-isostatic forging of a 60 to 80 percent dense aluminum-7.5 weight percent magnesium powder preform by this technique resulted in about a 30 percent axial compression and about a 10 percent radial expansion as taught by Meeks III et al. U.S. Pat. No. 7,097,807 and included herein by reference. The quasi-isostatic forging deformation of L12 aluminum alloy powder of this embodiment has beneficial effects on the microstructure and resulting properties of the consolidated billet. The nonuniform stress and resulting strain field in the powder during forging results in extensive shear deformation. The shear deformation deforms the powder, strips off the surface oxide from the L12 alloy particles and redistributes it throughout the consolidated L12 aluminum alloy powder forging as finely divided dispersoids. As a result, there is increased metal to metal contact during forging and resulting increased mechanical integrity. The redistributed surface oxide particles act as additional strengthening agents by resisting dislocation motion by Orowan strengthening.
In other embodiments, Ceracon type forging can be followed by hot isostatic pressing (HIP), forging, rolling and other deformation processing techniques.
To summarize, Ceracon type quasi-isostatic forging can be low cost and efficient. The pressure transmitting medium, as well as the vacuum sealed preform can be preheated prior to introduction of the perform into the forging chamber to minimize runtime. In addition, forging can be accomplished at high strain rates resulting in forging runs lasting less than a few minutes. The short run time results in an economical process that also further inhibits grain and L12 particle growth in the billet during forging due to the limited time at temperature. The pressure transmitting medium can be reused and loading and unloading the preforms and resulting forgings can be an automated procedure.
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.
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 L12 ALUMINUM ALLOYS, Ser. No. 12/316,020; A METHOD FOR FORMING HIGH STRENGTH ALUMINUM ALLOYS CONTAINING L12 INTERMETALLIC DISPERSOIDS, Ser. No. 12/316,046; and A METHOD FOR PRODUCING HIGH STRENGTH ALUMINUM ALLOY POWDER CONTAINING L12 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: L12 ALUMINUM ALLOYS WITH BIMODAL AND TRIMODAL DISTRIBUTION, Ser. No. 12/148,395; DISPERSION STRENGTHENED L12 ALUMINUM ALLOYS, Ser. No. 12/148,432; HEAT TREATABLE L12 ALUMINUM ALLOYS, Ser. No. 12/148,383; HIGH STRENGTH L12 ALUMINUM ALLOYS, Ser. No. 12/148,394; HIGH STRENGTH L12 ALUMINUM ALLOYS, Ser. No. 12/148,382; HEAT TREATABLE L12 ALUMINUM ALLOYS, Ser. No. 12/148,396; HIGH STRENGTH L12 ALUMINUM ALLOYS, Ser. No. 12/148,387; HIGH STRENGTH ALUMINUM ALLOYS WITH L12 PRECIPITATES, Ser. No. 12/148,426; HIGH STRENGTH L12 ALUMINUM ALLOYS, Ser. No. 12/148,459; and L12 STRENGTHENED AMORPHOUS ALUMINUM ALLOYS, Ser. No. 12/148,458.
