Patent Grant
9611522
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.
The present invention relates generally to aluminum alloys and more specifically to a method for forming high strength aluminum alloy product 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 A1 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 higher ductility. Fine interparticle spacing provides higher yield strength through Orowan dislocation-particle interaction model. Fine interparticle spacing is achieved by controlling the precipitate particles to fine size for a given volume fraction.
The present invention is a method for forming aluminum alloys with high strength and 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 spray deposition in which a stream of molten aluminum alloy containing L12 dispersoid forming elements is contacted with high velocity inert gas stream to form liquid droplets that are directed toward a substrate. The droplets solidify upon impact and form a solid deposit with a low degree of porosity. The aluminum alloy product thus formed can be deformation processed and heat treated to develop improved strength and fracture toughness. The method is efficient because melting and consolidation are combined in a single step. In addition, the rapid cooling rate experienced during droplet flight and impact leads to high supersaturation of solute and an increased amount of metastable L12 dispersoids in the aged alloys.
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 rapidly 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 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 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 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 results in improved thermal and structural stability.
Hafnium forms metastable Al3Hf dispersoids in the aluminum matrix that have an L12 structure in the metastable condition and a D023 structure in the equilibrium condition. The Al3Hf dispersoids have a low diffusion coefficient, which makes them thermally stable and highly resistant to coarsening. Hafnium has a high solubility in the Al3X dispersoids allowing large amounts of hafnium to substitute for scandium, erbium, thulium, ytterbium, and lutetium in the above mentioned Al3X dispersoids, which results in stronger and more thermally stable dispersoids.
Niobium forms metastable Al3Nb dispersoids in the aluminum matrix that have an L12 structure in the metastable condition and a D022 structure in the equilibrium condition. Niobium has a lower solubility in the Al3X dispersoids than hafnium or yttrium, allowing relatively lower amounts of niobium than hafnium or yttrium to substitute for X in the Al3X dispersoids. Nonetheless, niobium can be very effective in slowing down the coarsening kinetics of the Al3X dispersoids because the Al3Nb dispersoids are thermally stable. The substitution of niobium for X in the above mentioned Al3X dispersoids results in stronger and more thermally stable dispersoids.
Al3X L12 precipitates improve elevated temperature mechanical properties in aluminum alloys for two reasons. First, the precipitates are ordered intermetallic compounds. As a result, when the particles are sheared by glide dislocations during deformation, the dislocations separate into two partial dislocations separated by an anti-phase boundary on the glide plane. The energy to create the anti-phase boundary is the origin of the strengthening. Second, the cubic L12 crystal structure and lattice parameter of the precipitates are closely matched to the aluminum solid solution matrix. This results in a lattice coherency at the precipitate/matrix boundary that resists coarsening. The lack of an interphase boundary results in a low driving force for particle growth and resulting elevated temperature stability. Alloying elements in solid solution in the dispersed strengthening particles and in the aluminum matrix that tend to decrease the lattice mismatch between the matrix and particles will tend to increase the strengthening and elevated temperature stability of the alloy.
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 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.
Spray deposition is similar to gas atomization formation of powder, in that metal droplets are formed by the interaction of a high pressure gas stream with a secondary stream of molten metal. The gas atomizes the metal into molten droplets and accelerates the droplets. In powder production the droplets are allowed to solidify and are collected in a collection chamber. In spray forming, the stream of molten droplets impacts a target in a semi-molten state (mushy state) before they solidify to produce a solid near net shape. The benefits of spray forming or spray deposition, as it is sometimes referred to, are first, the process results in near net shaped product directly from the melt rather than going through a series of process steps including powder canning, long degassing time and consolidation of powder used in a powder metallurgy process and thereby eliminating a number of intermediate processes usually involved in forming metal parts. Secondly, the rapid solidification rate causes high solute supersaturation that, in the case of L12 aluminum alloys, results in maximizing the amount of L12 strengthening dispersoid content in the alloys. Thirdly, the rapid solidification rate minimizes alloy segregation in the billet. An additional advantage of spray forming is that nonmetallic materials, e.g. ceramics can be injected into the molten metal spray that are incorporated into the final billet as an additional strengthening dispersion.
The spray forming process is described in detail in U.S. Pat. No. 4,938,275 Leatham et al. and is included herein in entirety. The patent is assigned to Osprey Metals Ltd. and is commonly called the Osprey process by those skilled in the art.
A schematic of typical vertical spray forming process 10 is shown in
As gas builds up in spray chamber 55 during deposition, pressure in spray chamber 55 is controlled by gas exiting exhaust port 100 as indicated by arrow 110.
An added benefit of spray forming is that particulates such as ceramic powder can be added to spray 60 to provide additional strengthening. This process is indicated by particle reservoir 120 holding particles 130 and particle stream 140 entering spray 60 as spray forming proceeds.
Control of the following process parameters are critical to successful billet formation: melt superheat, melt flow rate, gas pressure, spray motion, spray height (distance between gas nozzles 50 and substrate 80) and substrate 90 motion (rotation rate and withdrawal rate). In spray forming, any material that can withstand the thermal spray environment without causing any reaction with aluminum to form undesirable intermetallic particles can be used as a substrate material. Stainless steel is preferred over other materials for a substrate due to its availability, high strength and inability to form intermatallic particles upon contact with aluminum.
Processing parameters are critical for forming solid billets with low porosity and high structural integrity. Important processing parameters include, among others, metal flow rate, superheat temperature, gas pressure, spray height, metal poring temperature, metal stream diameter, substrate preheat, atomizing gas, substrate rotation, substrate size, deposit thickness and deposit length.
Metal flow rates of about 5 lb/min (2.5 Kg/min) to about 50 lb/min (25 Kg/min) are preferred at superheat temperatures of about 150° F. (66° C.) to about 200° F. (93° C.). Lower flow rate results in finer powder and higher flow rate gives coarser powder for a given amount of gas and metal superheat. Finer powder is beneficial; however, in order to produce dense deposits due to good bonding with other powder particles, the metal powder needs to be in a semimolten stage instead of being completely solid. Higher metal flow rate results in relatively coarser powder which takes longer to solidify. Very high flow rate is undesirable because it provides droplets that remain completely liquid until they impact the substrate. The metal flow rate range given above provides good bonding characteristics resulting in dense deposits.
Gas pressures of about 80 psi (0.55 MPa) to about 500 psi (3.45 MPa) are preferred for the alloys disclosed herein. Lower gas pressure gives larger powder size and higher gas pressure results in finer powder size for a given metal flow rate. Lower gas pressure is still more preferred for spray deposition because it produces slightly coarser powder which remains in semi molten stage which is desirable for good bonding. However, in order to produce good properties the microstructure needs to be fine which is derived from finer powder size. Therefore, it is required to produce powder with sizes which contain an adequate fraction of liquid in order to produce balanced properties.
Fractions of liquid in atomized droplets of about 10 percent to about 50 percent are preferred for the alloys disclosed herein. It is extremely important to retain liquid in droplets before the powder is impacted to the substrate to allow good bonding with other powder particles. If the liquid fraction is too high, it will stick better with other powder. However, the properties will be inferior equivalent to cast products. If the liquid fraction is too low, the powder will not stick together. It is most preferred to have an adequate amount of liquid in droplets in order to produce dense deposits with good mechanical properties.
Spray heights of about 20 in (508 mm) to about 27 in (686 mm) are preferred for the alloys discussed herein. Spray height is the distance between the nozzle and where the metal stream impacts the substrate. Droplets solidify by liberating heat to the atomizing gas by convection. As the spray height is increased, powder particles have more time for solidification before they impact the substrate which reduces bonding of the powder to the substrate. Lower spray height allows less time for the powder to travel before it can solidify and therefore the droplet has a higher liquid fraction. The range given above for spray heights provides good bonding characteristics of the powder resulting in dense deposits.
Metal pouring temperatures of about 1400° F. (760° C.) to about 2200° F. (1204° C.) are preferred for the alloys discussed herein. Higher metal pouring temperature provides finer powder particle sizes due to more efficient disintegration of the metallic stream. Lower metal pouring temperature provides larger powder size. The metal pouring temperature range given above is wide because two different alloys have considerable differences in melting characteristics based on their compositions. The above metal pouring temperature range provides powders with good bonding characteristics resulting in dense deposits.
Metal stream diameters of about 4 mm to about 12 mm are preferred for the alloys discussed herein. Metal stream diameter controls the molten metal flow rate. Small metal stream diameters provide finer powder particles for a given gas pressure due to higher energy available for more efficient disintegration of metal stream. Too small metal stream diameters can create problems due to plugging of the nozzle. Large metal stream diameters provide large powder size due to inefficient break up of the metallic stream by the same amount of gas which was used for small metal stream diameter. The above range for metal stream diameters provides powder with good bonding characteristics resulting in dense deposits.
Substrate preheats of about 500° F. (260° C.) to about 800° F. (477° C.) are preferred for the alloys discussed herein. Substrate preheating improves the bonding of powder particles and produces dense deposits by reducing porosity. In addition, substrate preheat improves closer contact between the deposited metal and the substrate which makes it difficult for oxygen to penetrate. While it is hard to keep substrate temperatures too high because the atomizing gas cools off the substrate, it is desirable to heat the substrate. If the substrate is not hot, the first layer of metal that gets deposited will have a very fine microstructure because it takes away the heat more efficiently. However, subsequent layers which are deposited have coarser microstructures because the metal powder does not contact with the substrate. Instead, it contacts the hot deposit which can not extract heat quickly. Substrate preheat temperatures given above provide dense deposits.
The atomizing gas is preferred to be nitrogen, argon or helium. Helium has a higher transfer coefficient than those of nitrogen and argon resulting in finer microstructures due to higher cooling rate experienced. Argon provides a finer powder due to more efficient disintegration and is cheaper than helium. Nitrogen provides good powder sizes that are desired for good bonding. Nitrogen is cheaper than argon and helium. Based on the cost and good powder bonding characteristics obtained, nitrogen is even more preferred for spray deposition.
Substrate rotations of about 150 rpm to about 300 rpm are preferred for the alloys discussed herein. Substrate rotation provides uniform deposition of metal powder for making cylindrical products. In order to make different shapes, a substrate needs to be moved in different ways.
Substrate sizes of about 3″ diameter to about 6″ diameter and about 12″ long to about 60″ long are preferred for the alloys discussed herein. While it is preferred to use the substrate sizes and lengths described here, other sizes and lengths can be used also.
Deposit thicknesses of about 1″ to about 5″ and deposit lengths of about 5″ to about 50″ are preferred for the alloys discussed herein. While the deposit thicknesses and lengths mentioned here are preferred, other sizes can be used also.
Ceramic particles comprised of, but not limited to, SiC, B4C, TiB2, TiC, TiB and Al2O3 can be introduced to the alloy powder by coinjecting them in the atomization spray to improve the properties. The particle sizes of these reinforcements can range from about 1 micron to about 20 microns and volume fractions can range from about 5 percent to 25 percent. Ceramic reinforcements provide higher modulus, higher strength and higher wear resistance of the L12 aluminum alloys. However, ductility and fracture toughness often decrease due to lower ductility and fracture toughness of these ceramic reinforcements.
The sprayed deposit can be extruded, forged or rolled to further densify and to produce different shapes. Porosity is observed in the deposit if process parameters are not controlled properly. In that case, hot working including extrusion, forging and rolling is needed to further densify the deposit. In addition, hot working also breaks up the oxide in the deposit, distributes it more uniformly and provides improved mechanical properties. Extrusion, forging and rolling can also be used to produce different shaped components. It is preferred to perform these hot working operations in the temperature range of about 400° F. (204° C.) to about 800° F. (477° C.).
Although a vertical billet is described in
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.
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