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Publication numberUS20050084407 A1
Publication typeApplication
Application numberUS 10/909,061
Publication dateApr 21, 2005
Filing dateJul 30, 2004
Priority dateAug 7, 2003
Publication number10909061, 909061, US 2005/0084407 A1, US 2005/084407 A1, US 20050084407 A1, US 20050084407A1, US 2005084407 A1, US 2005084407A1, US-A1-20050084407, US-A1-2005084407, US2005/0084407A1, US2005/084407A1, US20050084407 A1, US20050084407A1, US2005084407 A1, US2005084407A1
InventorsJames Myrick
Original AssigneeMyrick James J.
Export CitationBiBTeX, EndNote, RefMan
External Links: USPTO, USPTO Assignment, Espacenet
Titanium group powder metallurgy
US 20050084407 A1
Abstract
Methods and compositions relating to powder metallurgy in which an amorphous-titanium-based metal glass alloy is compressed above its glass transition temperature Tg with a titanium alloy powder which is a solid at the compression temperature, to produce a compact with a relative density of at least 98%.
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Claims(1)
1. A powder manufacturing method for manufacturing an titanium group alloy product comprising the steps of blending from about 1 to about 25 volume percent of an amorphous metal powder having a glass transition temperature Tg, with from about 75 to about 99 volume percent of an titanium group metal powder, and compressing the blend at a temperature of at least Tg to produce a metal component having a density of at least 98% with about 2 percent or less porosity.
Description
CROSS REFERENCE TO RELATED APPLICATION

This application claims the benefit of the filing date of U.S. Provisional Application No. 60/493,176 filed Aug. 7, 2003.

FIELD OF THE INVENTION

The present invention is directed to powder metallurgy, and more particularly, to high performance powder metallurgy for titanium group metal alloys.

BACKGROUND OF THE INVENTION

Titanium, zirconium and hafnium (Titanium Group) alloys have high utility, but are relatively difficult to fabricate because of their susceptibility to oxidation and reaction with other materials at their high melting and forging temperatures.

Because of its relatively high temperature capability and high strength to weight ratio, titanium and its alloys are desirable for a variety of aerospace, industrial, marine, military and commercial applications where weight and/or high temperature performance are important, such as fan blades, compressor blades, discs, hubs and other components of turbine engines, automotive vehicle components such as engine valves, rocker arms, connecting rods and frames and sporting equipment. Because of its biocompatibility, corrosion resistance, titanium and its alloys are also used for chemical processing, desalination, power generation equipment, valve and pump parts, marine hardware, and prosthetic devices, implants, surgical devices and pacemaker cases. Because of its corrosion resistance, titanium is used for marine ball valves, fire pumps, heat exchangers, castings, hulks, water jet propulsion systems, shipboard cooling and piping systems. Titanium valve train components can substantially improve fuel efficiency because of their lightweight and high temperature operation.

Titanium alloys continue to be widely used in military and NASA space applications. In addition to manned space craft, titanium alloys are extensively employed in solid missile and rocket cases, and guidance control pressure vessels. Zirconium alloys are used in medical, chemical and nuclear reactor applications, while hafnium alloys find use in nuclear absorber components and rocket motors.

Titanium alloys which are reinforced with fibers or filaments such as high strength SiC filaments, are important materials. These fiber reinforced materials, such as titanium metal matrix composites, are conventionally produced by powder metallurgy and related technology in which the high strength, high modulus filaments are aligned and compacted at elevated temperatures with titanium alloy powders to form a dense metal matrix surrounding the SiC filaments. The preparation of titanium alloy base foils and sheets and of reinforced structures in which silicon carbide fibers are embedded in a titanium base alloy are described for example, in U.S. Pat. Nos. 4,775,547; 4,782,884; 4,786,566; 4,805,294; 4,805,833; 4,838,337; 5,939,313; 6,190,133 and 6,122,884.

Powder metallurgy is an important technology in which metal powders are formed and sintered to produce consolidated articles of manufacture. Powder metallurgy techniques include relatively simple procedures such as uniaxial powder compression in a mold followed by sintering at an elevated temperature somewhat below the melting point of the metal powder, as well as more complicated and expensive techniques such as hot isostatic pressing (HIPing) and metal injection molding (MIM). Titanium group metals are typically sintered under vacuum, or in inert or reducing gas atmosphere such as argon or hydrogen. Benefits of powder metallurgy fabrication can include near net shape manufacture, and control of crystal morphology. However, it is difficult to economically eliminate porosity and achieve full density using strong titanium group metals, because their relatively high temperature performance and high yield strength, which are desirable in the final products, make compaction difficult during manufacture. The relative density (D, in percent of full density) of a metal powder compacted during hot isostatic pressing by plastic yielding may be approximated by D=100[1−e(−3P/2Y)] where P is the consolidation pressure (in MPa) and Y is the yield strength (in MPa) of the alloy at the compression temperature [A. S. Helle, et al, Acta Metall. 33, 2163 (1985)]. For example, a titanium alloy powder having a nominal yield strength of just 300 MPa at the compacting temperature, and a practical maximum compression pressure of 750 MPa, the density of the compact is only about 98 percent, with about 2 percent remaining porosity in the finished product. Similarly, in metal injection molding (MIM) processes, powdered steel alloy powders may be mixed with a flowable thermoplastic organopolymer or other viscous binder to form a homogeneous, highly loaded mixture having, for example, approximately 60% volume metal powder and 40% volume of the flowable binder. The highly-solids-loaded mixture may be injection molded using conventional plastic injection molding systems, to produce molded “green” parts of considerable geometric complexity, which are highly filled with the titanium group metal alloy powder. The thermoplastic resin or other binder is subsequently thermally vaporized from the molded green parts in a debinding step, to leave a shaped metal part having high porosity. The porous, formed part may be subsequently sintered in an inert gas atmosphere, which densifies the part isotropically, while retaining the complex shape of the original molded part to relatively close tolerances. However, the finished, sintered parts retain considerable porosity, and carbide components from the thermoplastic binder.

Titanium forms a variety of refractory ceramic materials such as TiB2 and TiC and TiSi with small metalloid elements C, B and Si, and also forms relatively stable intermetallic compounds with metals including aluminum, nickel, cobalt and iron. Intermetallic compounds have a variety of attractive properties for high-temperature use, but have limited room-temperature ductility. Significant effort has been directed to the improvement of the ductility of titanium-aluminum intermetallic compounds. One approach which may provide increased ductility is reduction of grain size through compaction of controlled crystalline powder structures. Unfortunately, properties of consolidated powders may be affected by the conventional compaction methods such as hot-pressing because the prolonged heating can lead to coarsening of the microstructure. Efforts to avoid prolonged heating include explosive compaction in which a high-amplitude explosive stress wave compacts the titanium alloy powder. [E. Szewczak et al, “explosive consolidation of mechanically alloyed Ti—Al alloys”, materials science and engineering A226-228, pp 115-118 (1997)].

PM product density can be increased by use of high compression and sintering temperatures. However, properties of consolidated metal alloy powders may also be adversely affected because prolonged and/or high-temperature heating may lead to coarsening of the microstructure, and/or undesired reaction with reinforcing fibers or filaments.

DESCRIPTION OF THE DRAWINGS

FIG. 1 is a schematic cross sectional illustration of a low-temperature, low-pressure titanium alloy process in accordance with the present invention, utilizing a blend of a major amount of a crystalline titanium group metal powder, with a minor amount of a BMG alloy powder, prior to consolidation; and

FIG. 2 is schematic view of a process for forming a net-shape or net net-shape titanium alloy article of manufacture such as a gear, in accordance with an embodiment of the present invention.

SUMMARY OF THE INVENTION

The present invention is directed to methods for manufacturing titanium group powder metallurgy (PM) products, and to new titanium group PM products.

In accordance with such methods, a relatively small amount of an amorphous metal alloy powder having a glass transition temperature at a temperature Tg and supercooled liquid state at a temperature at or above its glass transition temperature, is blended with or coated on a major amount of a titanium group metal alloy powder. The powder blend is compressed at a temperature substantially below the melting point and within the supercooled liquid range of the amorphous alloy, to decrease the porosity of the blend and increase the total surface area of the amorphous powder. In this regard, the porosity of the mixture is preferably decreased to less than about 3 volume percent by compaction or other compression. When using an independent BMG amorphous metal powder (e.g., not a uniform amorphous metal glass coating on crystalline metal particles, the surface area of the amorphous metal alloy is preferably increased by at least 50 percent, and more preferably at least 100 percent during the compression step. The blend may also contain reinforcing powders or fibers such as B, C, Al2O3, SiC, SiCN, TiC powders, filaments or fibers, which will be incorporated in the finished PM products. Desirably, the titanium group powders and amorphous metal powders should have a particle size of less than about 300 microns, preferably with a range of particle sizes to facilitate packing. For some blends, it is desirable that the mean particle size of the amorphous metal component be less than half that of the titanium group metal particles, to facilitate more uniform dispersion in the interstices of the titanium group particles. The titanium group and amorphous powder components may, respectively, be produced in any suitable manner, such as by inert gas atomization, or centrifugal atomization of a molten metal stream, mechanical grinding or milling, chemical reaction, precipitation from the vapor phase, and/or electrolytic or electrodes deposition aqueous from or non-aqueous electrolytes.

Various preferred production methods of the present disclosure comprise blending from about 2 to about 30 volume percent, and more preferably from about 5 to about 25 volume percent, of amorphous metal alloy powder having a glass transition temperature below its crystallization temperature such that it has a supercooled liquid region, with from about 70 volume percent to about 98 volume percent, and more preferably from about 25 to about 95 volume percent, of a titanium group metal powder which does not have a supercooled liquid region (e.g., a conventional crystalline titanium or titanium group allow powder), based on the total volume (excluding voids) of the titanium group metal powder and the amorphous metal powder. The metal powder mixture is preferably blended to form a substantially homogenous mixture. The resulting blended metal powder mixture is formed and preferably compacted at a temperature in the supercooled liquid temperature range of amorphous metal powder, to produce a metal compact having a relative density of at least 98 percent, more desirably at least 99 percent, and most preferably substantially 100 percent (substantially fully dense). The powder mixture may be first compacted or otherwise compressed at a temperature below the glass transition temperature of the amorphous metal alloy component(s), which may also provide shear deformation at particle interfaces. This deformation produces new material surfaces, without surface oxidation or other passivation, which more readily bond to adjacent particles. Below Tg, the amorphous alloy may be typically stronger than the crystalline titanium group alloy powder, in which case shear deformation below Tg will largely occur at the crystalline titanium group powder surfaces, and even the reinforcement fiber surfaces.

The blend is subsequently compressed to high relative density at a temperature in the supercooled liquid temperature range at which the amorphous metal alloy powder is a viscous liquid, rather than a metal melt. This compression produces viscous flow of the amorphous metal glass, forcing it into the interstices of the titanium group metal particles, substantially increasing the total surface area of the amorphous metal component and desirably forming it into an at least partially interconnected matrix at least partially enclosing the titanium group metal particles.

This substantial increase in surface area, and decrease in the median thickness of the amorphous metal component facilitates reactive diffusion with the titanium group metal component at elevated temperature, if desired, to form new alloy(s) having higher melting points than the amorphous metal component, which has a relatively low melting point.

As indicated, the powder mixture comprises a major volume fraction V of titanium group metal powder and any filler or reinforcing agent powders, fibers or filaments, and a minor volume fraction v of the BMG amorphous metal powder or coating component. The compressive pressure applied to compact the mixed powder will desirably be at least that necessary to compact the titanium group metal powder and the filler or reinforcing component(s) (if any) to a relative density R of at least 97-100 v, and more preferably at least 98-100 v. This compression pressure Pc may be approximated by Pc=−2Y{ln(1−R)/3} where Pc is the minimum compression pressure at the supercooled liquid temperature, Y is the yield strength of the titanium group metal powder at the compression temperature, and R is the relative density of the titanium group metal powder (in the absence of the BMG amorphous alloy component).

For example, an titanium group metal powder having a yield strength of 400 MPa at 700° K. would itself (without the presence of a BMG amorphous metal component) be compressed to a relative density R, of 0.9 of full density, by a compression pressure Pc of approximately −2×400 (ln(1−0.9)/3 or approximately Pc=615 MPa. To compress the titanium group metal powder itself (without a BMG amorphous metal component) to a relative density of 0.98, leaving 2 percent porosity, would require over 1 Gpa (over 100,000 psi) which is difficult and expensive to provide in an economical production process under vacuum or inert atmosphere.

A mixture of 87 volume percent of this titanium group metal powder, and 13 volume percent of a BMG amorphous metal powder (or a BMG coating on a crystalline titanium alloy powder) with a glass transition at a temperature of 650° K. and a supercooled viscous glass liquid region of 100° (Tx=750° K.) could accordingly be compressed at 700° C using a compression pressure Pc of about 615 MPa, to produce a compact product having a relative density of 0.97 or higher.

The compacted PM product may be cooled to provide a dense metal compact in which the crystalline titanium group metal particles are embedded in an amorphous metal matrix which has substantially increased in total surface area to fill zones between the titanium group powders. The compacted PM product may also be heated to a temperature above the crystallization temperature of the amorphous metal continuous phase. In this manner, the iron metal group particles are embedded in a fully or partially crystalline continuous phase which is still substantially distinct from the titanium group metal particles. This can be advantageous for certain BMG alloys which form a strong, partially or fully crystalline alloy upon heating to an initial crystallization temperature Tx or above. The compact may then be heated to a temperature above the crystallization temperature of the amorphous metal, to nanocrystallize the amorphous component. There is very little shrinkage, if any, because the metal glass is already close to crystal density. The crystallization can also initiate from the crystalline metal particles, to get a good boundary.

Such PM products with discrete titanium group powder zones (even though compressed or distorted) and BMG alloy composition zones, may not have optimum heat resistance under stress, however, because the at least partially discrete continuous amorphous phase may remain as a eutectic or near-eutectic material with a relatively low melting point, Tm. In order to provide a finished PM product with increased heat resistance, the compacted product may also be further heated to a sintering temperature at which the titanium group alloy particles and the relatively continuous matrix phase diffuse together to form a new composition having a higher melt temperature than the original SLM amorphous alloy.

After diffusion of the original crystalline metal powder components with the original amorphous metal components, the composite does not “remelt” at the relatively low melting temperature Tm of the bulk metal glass component, because the different effective composition of the reacted compact has a higher melting point than the original BMG.

While some amorphous metals fully or partially crystallize over a limited period of time at temperatures coextensive with or only slightly above their effective glass transition temperature, a wide variety of amorphous metal compositions are relatively stable at temperatures at or slightly above their glass transition temperature, Tg, and do not initiate substantial crystallization unless raised to a crystallization temperature, Tx, which may be 10 to 100 or more degrees Celsius higher than Tg. By “bulk metal glass” (BMG) is meant an amorphous metal alloy composition having a glass transition temperature, Tg, at which it exhibits a supercooled liquid phase for at least one second, and preferably at least 30 seconds.

The glass transition temperature Tg (if any) and the crystallization temperature(s) Tx of an amorphous alloy are typically determined by differential scanning calorimetry, in which the temperature of a sample is slowly raised, and correlated, as a function of temperature, with the amount of energy necessary to raise the temperature. The glass transition phase change is typically an endothermic process, while crystallization is typically an exothermic process involving slight volume increase. Many, if not most, amorphous metal compositions do not have a glass transition temperature, but instead crystallize at one or more elevated temperatures without going through a distinct viscous glass transition phase. An amorphous metal composition may have a number of distinct crystallization temperatures Tx(1), Tx(2) . . . at which various components crystallize or recrystallize from components crystallized in a less stable or metastable crystalline phase at a lower crystallization temperature. As will be discussed, a variety of amorphous metal alloys have a distinct glass transition Tg, at which they undergo a slight volume expansion upon phase transition to a viscous glass state, and undergo partial crystallization, typically forming nanoscale crystallites in an amorphous matrix which remains in a viscous glassy state. These partially-nano-crystalline bulk metal glasses retain a viscous glassy matrix above Tg, and are useful in the present methods and are considered to have a supercooled liquid temperature region in which they form a viscous glass, albeit one with nanoscale crystallites at high temperatures, still below their metal temperature Tm, they will fully crystallize and lose their viscous, supercooled glass condition.

The determination of glass transition temperature and crystalline temperature(s) is typically a function of the rate at which the temperature of the metal glass foam is increased. For purposes of this disclosure, a rate of temperature increase of 0.25 degrees Celsius per second may be used to determine Tg, although other rates are used in determining reported Tg and Tx values herein.

Bulk metal glasses (BMGs) used herein preferably have a crystallization temperature, Tx, which is at least 20° C. and more preferably at least 40° C. higher than the glass transition temperature, Tg, of the bulk metal glass.

Amorphous metal alloys may be produced by rapidly cooling appropriate alloy compositions from their homogeneous melts. The temperature Tm at which an amorphous metal alloy fully melts, is generally significantly higher than the crystallization temperature(s) Tx, which in turn is higher than the glass transition temperature Tg (if any), of the alloy. Accordingly, a “typical” BMG may be heated from ambient temperature to its glass transition Tg (e.g., about 300° C.) at which is undergoes a phase change to a viscous supercooled liquid phase and undergoes a small volume expansion. Further heating produces a reduction of viscosity in the supercooled liquid until the crystallization temperature(s) Tx is (are) reached (e.g., 400° C.) at which crystalline phases nucleate and grow. There may be a number of crystallization temperatures Tx, at which different crystal phases and compositions are formed. If the BMG is heated above its Tg but below Tx, it may be cooled below its Tg and retain its glassy state. However, if such a BMG is raised to its crystallization temperature(s), Tx (Tx1, Tx2, etc.) it remains partially or fully crystallized when cooled, unless it is fully remelted at a much higher temperature (e.g., 1500° C.) and then rapidly cooled below its Tg. As indicated, some BMG alloys initially partially crystallize while retaining their viscous glass properties in their supercooled liquid region, and accordingly are useful as BMG powders and coatings in the present methods and products.

Amorphous metal alloys may have exceptionally high impact resistance and strength, which are important qualities for various metal products and components. For example, Bulk Metal Glasses (BMGs) based on Fe, Zr, Ti, Cu, Mg and/or Al metal systems can exhibit unique combinations of high hardness, strength, toughness and corrosion resistance, which are ideal for cellular foam materials. BMG alloys such as Fe—(Zr,Ti,Ni,Co,Mo)—(B,C,Si,P); Zr—Ni—Al—Cu; and Zr—Ti—Cu—Ni—(Si,Be) exhibit very good bulk glass-forming ability with high thermal stability in the supercooled glass state, and low critical cooling rates. [See, e.g., A. Inoue, et al., Mater. Trans. JIM, 31 (1991), p. 425; T. Zhang, et al., Mater. Trans. JIM, 32 (1991), p. 1005; A. Inoue et al., Mater. Trans. JIM, 32 (1991), p. 609; A. Peker, et al., Appl. Phys. Lett., 63 (1993), p. 2342.

The toughness of amorphous metals, including bulk metal glasses (BMGs) can increase with increasing impact or shear rates, to relatively high levels, which would be an important characteristic for metal foams designed for energy absorption. The more stable BMG alloys typically form dense, deep eutectic liquids with relatively small free volume, and relatively high melt viscosity, above their glass transition temperature, Tg. They typically comprise three, and preferably four or more components having negative heats of mixing and at least 12% difference in atomic size, in proportions which permit high packing density and short-range order. Being energetically close to the crystalline state in this manner, can provide slow crystallization kinetics, with high viscosity and high glass forming ability. R. Busch, “The Thermophysical Properties of Bulk Metallic Glass-Forming Liquids”, JOM, 52:7 (2000), pp. 39-42. However, the thickness of amorphous metal alloys which can be formed directly by casting from the melt is generally limited by the cooling rate and thermal conductivity. Reducing the thermal conductivity of a high temperature amorphous metal alloy melt by introducing gas bubbles prior to rapid cooling is possible, but does not facilitate the rapid cooling necessary for thick section casting.

A wide variety of high-strength glassy metal alloys may be used in the present methods. Amorphous Al94V4Fe2, with Al nanocrystals in an amorphous matrix, has relatively high tensile strength (1390 MPa) for a very light alloy. [Inoue, A., et al., “Deformation and fracture behavior of high-strength Al94(V, Ti)4Fe2 and Al93Ti5Fe2 Chemistry and Physics of Nanostructures and Related Non-Equilibrium Materials, TMS Annual Meeting, p 201-209 (1997); “Alloys consisting of nanogranular amorphous and Al phases”, TMS Annual Meeting, p 201-209 (1997)] Lightweight Ti-based bulk amorphous alloys such as Ti50Ni20Cu23Sn7 and Ti50Cu20Ni25Si4B2 can have very high tensile strength (2200 MPa) resulting in high strength-to-weight ratio, and a wide supercooled liquid region of viscous flow. [See, e.g., Louzguine, D. V., et al., “Nanocrystallization of Ti—Ni—Cu—Sn Amorphous Alloy”, Scripta mater., 43:371-376 (2000); Zhang, T., et al., Mater. Trans. JIM, 39:1001 (1998)]

Fe75CoNiSi8B 17 glassy alloys have a Young's modulus of 110 GPa, compressive fracture strength of 2800 MPa and fracture elongation of 1.9%. The glassy iron-based alloys exhibit a distinct glass transition, followed by a supercooled liquid region of over 50° C. before crystallization in a wide composition range of 7.5 to 45 at % Co and 7.5 to 60 at % Ni. Zhang, T.; et al., “Bulk glassy alloys in (Fe, Co, Ni)—Si—B system”, Materials Transactions, Vol. 42, pp. 1015-1018 (2001).

As indicated previously, these amorphous alloys can be used as a minor component of a powder blend with a relatively large amount of titanium group metal powder. The compacted composite may be formed with the titanium group metal powder component in a distinct matrix of the BMG amorphous metal alloy, or the compacted composite may be further heated to diffuse and react the components to form new alloys having a higher melt temperature than the original BMG. Desirably, the final composition after reaction above Tx of the BMG alloy, will have a melting temperature, Tm at least 100° C. higher than the BMG composition melting temperature Tm, and more preferably at least 250° higher.

Having described the BMG component, the titanium group alloy component will now be described. As indicated, the major component of the PM mixture is a titanium group alloy powder. By “titanium group alloy” is meant an alloy comprising at least 35 weight percent of the Group 4 periodic table elements titanium, zirconium and hafnium. Preferably, the titanium group alloy will comprise at least about 45 weight percent Ti, Zr and/or Hf, and more preferably at least about 60 weight percent of one or more of these elements, based on the weight percent total of the alloy powder. The titanium group alloy powder, as used herein, may also include substantially pure Ti, Zr and/or Hf powders, although titanium group alloys also containing other elements are preferred. Examples of other elements and their weight percent range in the alloy, are as follows:

Element Weight Percent
B 0-20
Be 0-5 
C  0-2.5
Co  0-100
Cr 0-25
Fe 0-60
Mn 0-10
Mo 0-15
Nb 0-15
Ni 0-60
Re 0-10
Si 0-10
Sn 0-5 
Ta 0-5 
W 0-20
Bi 0-5 
Ga 0-5 

Other elements may be present as well, and total Ti, Zr, Hf should best be at least weight percent, based on the total titanium group alloy weight.

Examples of preferred titanium group metal powders include titanium alloys, such as alpha, near alpha, alpha-beta and beta titanium alloys:

TITANIUM ALLOYS (weight percent)
Alpha Near Alpha Alpha Beta Beta
Ti—5Al—2.5Sn Ti—2.25—Al—11Sn—5Zr—1Mo—0.2Si Ti—10V—2Fe—3Al Ti—11.5Mo—6Zr—4.5Sn
Ti—5Al—5Sn—2Zr—2Mo—0.25Si Ti—3Al—2.5V Ti—13V—11Cr—3Al
Ti—6Al—2Nb—1Ta—1Mo Ti—6Al—2Sn—2Zr—2Mo—2Cr—0.25Si Ti—3Al—8V—6Cr—4Mo—4Zr
Ti—6Al—2Sn—4Zr—2Mo Ti—6Al—2Sn—4Zr—6Mo Ti—5Al—2Sn—2Zr—4Mo—4Cr
Ti—6Al—2Sn—1.5Zr—1Mo—0.35Bi—0.1Si Ti—6Al—4V Ti—8Mo—8V—2Fe—3Al
Ti—6Al—6V—2Sn—0.75Cu
Ti—7Al—4Mo
Ti—8Mn

Examples of commercial alloys with practical total weight percent ranges are exemplified as follows:

    • Ti-10V-2Fe-3Al Ti=83-86.8%
    • Ti-15V-3Cr-3Sn-3Al Ti=72.6-78.5%
    • Ti-2.5Cu═Ti=96.1-98% (high strength)
    • Ti-3Al-2.5V-0.05Pd Ti=92.5-95.5%
    • Ti-6Al-7Nb Ti=84.5-88% (high strength)
    • Ti-5Al-1Sn-1Zr-1V-0.8Mo Ti=88.5-93%
    • Ti-4Al-4Mo-2Sn Ti=85.9-92.2%
    • Ti-4Al-4Mo-4Sn-0.5Si Ti=83.2-90.7%
    • Ti-6Al-2Sn-4Zr-2Mo-0.08Si Ti=83.6-87.2%
    • Ti-6Al-2Sn-4Zr-6Mo Ti=79.4-83.7% (high temperature)
    • Ti-6Al-2Fe-0.1Si Ti=90.7-93%
    • Ti-11Sn-5Zr-2.25Al-1Mo-0.2Si Ti=77.9-82.6% (high temperature)
    • Ti-6Al-5Zr-0.5Mo-0.25Si Ti=85.8-89.5%
    • Ti-5.5Al-3.5Sn-3Zr-1Nb-0.25Mo-0.3Si Ti=84.2-88.1% (high temperature)
    • Ti-5.8Al-4Sn-3.5Zr-0.7Nb-0.5Mo-0.35Si-0.06C Ti=81-87.4% (high temperature)

The alpha phase is typically fostered and strengthened by aluminum and/or tin. Zr may also be present in alpha alloys.

Beta phase stabilizers (e.g., Mo, V, Fe, Cr, Ni, Nb, Hf, Ta, Mn, Co, Cu, W) promote the relatively ductile beta phase and thus enhance heat treatability and fabricability. Beta eutectoid stabilizers (e.g., Fe, Ni and Cr) may be used in combination with beta isomorphous stabilizers (e.g., Mo and V) provide definite advantages over additions of beta isomorphous stabilizers alone.

Alpha-beta titanium alloys may contain alpha-forming strengthening elements (aluminum, tin and zirconium) together with beta phase solid solution elements such as molybdenum, chromium, vanadium, iron and nickel to promote strength and ductility.

Examples of titanium group based alloy compositions suitable as crystalline major component powders, or as final composition alloys after diffusion-reaction of the BMG amorphous metal matrix with an titanium group alloy in the compacted part, include:

Titanium alpha alloys typically contain aluminum and tin, as well as molybdenum, zirconium, nitrogen, vanadium, columbium, tantalum, and silicon. Alpha alloys are not generally designed for heat treatment, but are weldable and are commonly used for cryogenic applications, aircraft parts, and chemical processing equipment.

Alpha-beta alloys can be strengthened by heat treatment and aging, and therefore can undergo processing and fabrication while the material is still ductile, then undergo heat treatment to strengthen the material for use in aircraft and aircraft turbine parts, chemical processing equipment, marine hardware, and prosthetic devices.

Beta titanium alloys have good hardenability, and cold formability when they are solution-treated, and high strength when they are aged. Beta alloys tend to be denser than other titanium alloys but have high yield strengths.

Zirconium alloys such as 99Zr, 1Nb and Zircaloy-2 98.5Zr0.1Cr0.1Fe and 0.05 Ni-1.4Sn have high temperature performance and are used for chemical and nuclear applications. Hf alloys have high reaction absorption for nuclear applications.

As indicated, titanium and its alloys are used as the matrix component in high performance metal-matrix composites (MMCs) for gas turbine engine components, including fan blades, fan frames, actuators, rotors, vanes, cases, ducting, shafts, and liners. Titanium-based metal matrix composites are important structural materials which are strong, temperature resistant, and relatively stiff for their relatively light weight.

Substantial research effort has been directed to developing Ti MMC materials, but despite this effort, current Ti MMC manufacturing capabilities are still relatively limited and material costs are relatively high. It would be desirable to reduce the cost of Ti MMC parts to affordable levels, to facilitate use of such parts in gas turbine engine and other aerospace applications.

DETAILED DESCRIPTION OF THE INVENTION

Bulk Metallic Glasses (BMGs) having a broad supercooled liquid range (preferably at least about 25° C., and desirably at least 50 K or more) can be fabricated from their low-temperature glassy melts in a bulk form at slow cooling rates of the order of 1-100K/s. The metallic glasses can have very high strength at ambient temperature. They can be heated to form a relatively viscous, “soft” glass without crystallizing, typically at temperatures below 500° C., e.g., 200-450° C. Bulk metallic glasses with the same or similar tensile strength as the cast bulk glassy alloy and melt-spun metallic glass ribbon can also be shaped by extrusion of gas-atomized glass powders. Kawamura, Y. et al., “Development and industrialization of powder-consolidation and forging technologies in metallic glasses” Funtai Oyobi Fummatsu Yakin/Journal of the Japan Society of Powder and Powder Metallurgy, Vol. 48, p 845-853 (2001).

BMGs such as Zr—Al Ni—Cu glasses and La—Al—Ni glasses can be easily forged into complex shapes above their glass transition temperatures, Tg. See, A. Inoue at al, “Novel Superplasticity of Supercooled Liquid for Bulk Amorphous Alloys”, Materials Science Forum, Vol. 243-245, pp. 197-206 (1997). The stress-strain rate curves for such supercooled ZrAl—Ni—Cu glass liquid shows the decreasingly viscous nature of the supercooled metal glass liquid as the temperature is increased above Tg, but below Tx. For example, the bulk metal glass Zr65Al10Cu15Ni10 has a glass transition temperature of about 652° K. (˜379° C.). At temperatures slightly above Tg (379 to 440° C.), this metal forms a viscous, easily shapable glass which is easily deformed or extruded. Such viscous metal glasses can be readily shaped, formed and/or extruded as desired. In accordance with the present disclosure, they may be used to manufacture cellular metal foams with a wide variety of high performance characteristics.

Mg—Cu—Y based bulk amorphous alloys such as Mg60Cu30Y10 have a high strength to weight ratio, and good flow behavior in the supercooled liquid region, which provides excellent formability at relatively low temperatures. Mg60Cu30Y10 has a Tg of about 400° K. (125° C.) and a Tx of about 450° K. (175° C.).

A variety of amorphous metals with their Tg, Tx and supercooled liquid region are listed in the following Table (with compositions given at atomic %):

Major
element Alloy Composition Tg (K) Tx (K) Tx − Tg Ref.
Mg— Mg80Ni10Nd10 454 471 17 k
Mg75Ni15Nd10 450 470 20 k
Mg60Cu30Y10 419 466 47 c
(Mg99Al1)60Cu30Y10 419 459 40 c
(Mg98Al2)60Cu30Y10 421 454 33 c
(Mg96Al4)60Cu30Y10 411 455 44 c
(Mg95Al5)60Cu30Y10 415 453 38 c
(Mg93Al7)60Cu30Y10 411 445 34 c
Mg70Ni15Nd15 467 489 22 k
Mg65Ni20Nd15 459 501 42 k
Mg65Cu25Y10 425 479 54 k
Mg60Cu30Y10 400 450 50
Zr— Zr66Al8Ni26 672 708 36 k
Zr66Al8Cu7Ni19 662 721 59 k
Zr66Al8Cu12Ni14 655 733 78 k
Zr66Al9Cu16Ni9 657 736 79 k
Zr65Al7.5Cu17.5Ni10 657 736 79 k
Zr57Ti5Al10Cu20Ni8 677 720 43 k
Zr41.2Ti13.8Cu12.5Ni10Be22.5 623 672 49 k
Zr38.5Ti16.5Ni9.75Cu15.25Be20 630 678 48 k
Zr39.88Ti15.12Ni9.98Cu13.77Be21.25 629 686 57
Zr42.63Ti12.37Cu11.25Ni10 Be23.75 623 712 89 k
Zr44Ti11Cu10Ni10Be25 625 739 114 k
Zr55Al10N15Cu30 683 748 65 d
Zr45.38Ti9.62Cu8.75Ni10Be26.25 623 740 117
Zr65Al10Ni10Cu15 652 757 105 e
Zr65Al7.5Cu17.5Ni10 633 749 116 i
(Zr65Al7.5Cu17.5Ni10)95Fe5 650 725 75 i
(Zr65Al7.5Cu17.5Ni10)90Fe10 670 730 60 i
(Zr65Al7.5Cu17.5Ni10)85Fe15 675 735 60 i
(Zr65Al7.5Cu17.5Ni10)80Fe20 680 740 60 i
Zr52.5Cu17.9Ni14.6Al10Ti5 686 725 39 j
(Zr67Hf33)52.5Cu117.9Ni14.6Al10Ti5 708 753 45 j
(Zr50Hf50)52.5Cu17.9Ni14.6Al10Ti5 722 767 45 j
(Zr33Hf67)52.5Cu17.9Ni14.6Al10Ti5 737 786 49 j
Zr52.5Cu17.9Ni14.6Al10Ti5 767 820 53 j
Zr52.2Ti16.7Cu17.7Ni8.7B4.7 564 668 104 l
Zr50.2Ti16.7Cu17.7Ni8.7B6.7 646 719 73 l
Zr48.2Ti16.7Cu17.7Ni8.7B8.7 682 720 38 l
Zr54.9Ti16.7Cu17.7Ni8.7P2.0 578 686 108 l
Zr53.9Ti16.7Cu17.7Ni8.7P3.0 636 722 86 l
Zr52.9Ti16.7Cu17.7Ni8.7P4.0 698 734 36 l
Zr54.9Ti16.7Cu17.7Ni8.7Si2.0 562 681 119 l
Zr53.9Ti16.7Cu17.7Ni8.7Si3.0 563 681 118 l
Zr52.9Ti16.7Cu17.7Ni8.7Si4.0 639 742 103 l
Zr41.2Ti13.8Cu12.5Ni10Be22.5 633 741 108 l
Zr70Fe20Ni10 646 673 27 o
Zr60Al10Cu30 680 750 70 p
La— La55Al25Ni10Cu10 467 547 80 k
La55Al25Ni5Cu15 459 520 61 k
La55Al25Cu20 456 495 39 k
La55Al25Ni5Cu10Co5 465 542 77 k
La66Al14Cu20 395 449 54 k
La60Al20Ni10Co5Cu5 451 523 72 g
Pd— Pd40Cu30Ni10P20 577 656 79 k
Pd81.5Cu2Si16.5 633 670 37 k
Pd79.5Cu4Si16.5 635 675 40 k
Pd77.5Cu6Si16.5 637 678 41 k
Pd77Cu6Si17 642 686 44 k
Pd73.5Cu10Si16.5 645 685 40 k
Pd71.5Cu12Si16.5 652 680 28 k
Pd40Ni40P20 590 671 80 k
Nd— Nd60Al15Ni10Cu10Fe5 430 475 45 k
Nd61Al11Ni8Co5Cu15 445 469 24 k
Cu— Cu60Zr30Ti10 713 763 50 k
Cu54Zr27Ti9Be10 720 762 42 k
Cu48Ti34Zr10Ni8
Ti— Ti34Zr11Cu47Ni8 698 727 29 k
Ti50Ni24Cu20B1Si2Sn3 726 800 74 k
Ti50Ni24Cu20B1Si2Sn3 726 800 74 h
Ti45Ni20Cu25Sn5Zr5
Ti50Cu25Ni25 713 753 40 m
Ti50Ni22Cu25Sn3 715 765 50 m
Ti50Ni20Cu25Sn5 710 770 60 m
Ti50 Ni20Cu23Sn7 710 759 49 m
Ti50Ni24Cu25Sb1 707 740 33 m
Ti50Ni22Cu25Sb3 763 718 45 m
Ti50Cu35Ni12Sn3
Ti50Cu40Ni4Si4B2 740 786 46 m
Ti50Cu20 Ni24Si4B2 745 810 65 m
Fe— Fe63Ni7Zr10B20 553 579 26 b
and/or Fe56Ni14Zr10B20 579 601 22 b
Co— Fe49Ni21Zr10B20 589 611 22 b
Fe42Ni28Zr10B20 602 619 18 b
Fe42Co7Ni21Zr10B20 580 611 30 b
Fe72Hf8Nb2B18 856 932 76 f
(Fe, Co)85 Zr7B6(Nb, Nd)2
Fe 74.5Si13.5B9Nb3
Fe58Co7Ni7Zr8B20 821 899 78 n
Fe52Co10Nb8B30 907 994 87 n
Fe62Co9.5RE3.5B25 (RE = Pr, Nd, >50
Sm, Gd, Tb, Dy, Er) (22.5-30 at % B, 0-30
at % Co and 2.5-6 at % RE)
Co63Fe7Zr6Ta4B20 858 895 37 n
Co40Fe22Nb8B30 895 976 81 n
Co43Fe20Ta5.5B31.5 910 980
Fe75-x-yCoxNiySi8B17 (7.5 to 45 Up to 54 s
at % Co and 7.5 to 60 at % Ni)
Fe90-xNb10Bx t
Fe85.5Zr2Nb4B8.5
Fe70B20Zr8Nb2 91 u
Al— Al85Ni5Y8Co2 538 570 32 a
Al87(La,Nd,Pr)8Ni5 500 553 53 r
Al92(La,Nd,Pr)4Ni4 525 608 83 r
Al—Ti—M (M = V, Fe, Co and/or Ni) v
alloys, such as Al94V4Fe2,
Al93Ti4Fe3, Al93Ti4V3
Al94V2Ti2Fe2
Al93Ti5Fe2

a. Kawamura, Y., et al., “Nanocrystalline Aluminum Bulk Alloys with a High Strength of 1420 MPa Produced by the Consolidation of Amorphous Powders”, Scripta mater., 44; 1599-1604 (2001)

b. Liu, Y. J., et al., “The correlation of microstructural development and thermal stability of mechanically alloyed multicomponent Fe—Co—Ni—Zr—B alloys”, Acta Materialia, 50, 2747-2760 (2002)

c. Linderoth, S., et al., “On the stability and crystallization of bulk amorphous Mg—Cu—Y—Al Alloys”, Materials Science and Engineering A304-306, 656-659 (2001)

d. deOliveira, M. F., et al., “Effect of oxide particles on the crystallization behaviour of Zr55Al10Ni5Cu30 Alloy”, Materials Science & Engineering A304-306, 665-6659 (2001)

e. Kawamura, Y., et al., “Newtonian and non-Newtonian viscosity of supercooled liquid in metallic glasses”, Materials Science & Engineering, A304-306, 674-678 (2001)

f. Kawamura, Y., et al., “Superplasticity in Fe-based metallic glass with wide supercooled liquid region”, Materials Science & Engineering, A304-306, 674-678 (2001)

g. Saotome, Y., et al., “Superplastic micro/nano formability of La60Al20Ni10Co5Cu5 amorphous alloy in supercooled liquid state”, Materials Science & Engineering, A304-306, 716-720 (2001)

h. Zhang, T., et al., “Ti-based amorphous alloys with a large supercooled liquid region”, Materials Science & Engineering, A304-306, 771-774 (2001)

i. Mattern, N., et al., “Influence of iron additions on structure and properties of amorphous Zr65Al7.5Cu17.5Ni10”, Materials Science and Engineering A304-306, 311-314 (2001)

j. Glass-forming ability and crystallization of bulk metallic glass (HfxZr1-x)52.5Cu17.9Ni14.6Al10Ti5”, Journal of Non-Crystalline Solids”, 311 77-82 (2002)

k. Lu, Z. P., et al., “A new glass-forming ability criterion for bulk metallic glasses”, Acta Materialia, 50, 3501-3512 (2002)

l. Choi, et al., “Effect of Additive Elements on the Glass Forming Ability and Crystallization of Zr—Ti—Cu—Ni Metallic Glasses”, Journal of Metastable and Nanocrystalline Materials, Vols. 343-346, pp. 109-115 (2000)

m. Inoue, A., “Synthesis and Properties of Ti-Based Bulk Amorphous Alloys with a Large Supercooled Liquid Region”, Journal of Metastable and Nanocrystalline Materials, Vols. 2-6 (1999), pp. 307-314

n. Inoue, et al., “Ferromagnetic Bulk Glassy Alloys with Useful Engineering Properties”, Journal of Metastable and Nanocrystalline Materials, Vols. 343-346, pp. 81-90 (2000)

o. Saida, et al., “Nano-Icosahedral Quasicrystalline Phase Formation from a Supercooled Liquid State in Zr—Fe Ternary Metallic Glass”, Applied Physics Letters, Vol. 76, No. 21, pp. 3037-3039 (May 22, 2000)

p. Inoue, et al., “Synthesis of High Strength Bulk Nanocrystalline Alloys Containing Remaining Amorphous Phase”, Journal of Metastable and Nanocrystalline Materials, Vol. 1, pp. 1-8 (1999)

q. Eckert, J., “Mechanical Alloying of Highly Processable Glassy Alloys”, Materials Science and Engineering, A226-228, pp. 364-373 (1997)

r. Tong, et al., “Microstructure and Thermal Analysis of Amorphous Al87RE8Ni5 and AL92RE4Ni4 Alloys”, Materials Letters, Vol. 28, pp. 133-136 (1996)

s. Zhang, et al, “Bulk glassy alloys in (Fe, Co, Ni)—Si—B system”, Materials Transactions, v 42, (2001)

t. Imafuku, et al, “Structural variation of Fe—Nb—B metallic glasses during crystallization process”, Scripta Materialia, v 44 (2001)

u. Ma, et al, “Fe-based metallic glass with significant supercooled liquid region of over 90 K”, Journal of Materials Science Letters, v 17 (1998)

v. Kimura, et al, “Formation of nanogranular amorphous phase in rapidly solidified Al—Ti—M (M = V, Fe, Co or Ni) alloys and their mechanical strength”, Nanostructured Materials, v 8, p 833-844 (1997)

S. J. Poon, et al. in “Glass formability of ferrous- and aluminum-based structural metallic alloys.” Journal of Non-Crystalline Solids v317, pp1-9(2003) describe and refer to a variety of ferrous- and aluminum-based amorphous metals which are also incorporated herein by reference.

However, while BMGs are superb structural materials due to a unique combination of properties such as a large elastic strain limit, high strength and good fracture toughness, they also may be characterized by high concentration of shear deformation under intense loading conditions.

Bulk metallic glasses (BMGs) may have a reasonably large elastic strain limit (e.g., about 2%), high strength (e.g., up to 3.8 GPa or more) and a good fracture toughness (up to 55 MPa m ½). Shear deformation concentration in an amorphous metal matrix may be reduced through the use of nonamorphous reinforcements, such as in situ precipitates or composite reinforcements. [see, Clausen, B., et al., “Deformation of In-Situ-Reinforced Bulk Metallic Glass Matrix Composites”, Mater. Sci. Forum (2002)], and through the use of BMG composites, as described hereinafter.

Particulate reinforcements may be used, or nanoscale precipitates may be formed in the glass directly from the amorphous metal composition by partial crystallization which retains a supercooled liquid matrix. Bulk nanocrystalline amorphous metal alloys such as aluminum-based amorphous alloys of Al—Ni—Y—Co, Al—Si—Ni—Ce and Al—Fe—Ti-TM (TM=Cr,Mo,V,Zr) systems, such as A189 (Ni0.33Y0.54Co0.13) 11 may be produced by arc-melting the pure metals in the appropriate composition ratios, in an argon atmosphere. Rapidly solidified BMG alloy powders may be produced by high pressure He gas atomization at a dynamic pressure of about 10 MPa with powder below 25 micron particle size being rapidly cooled to retain the amorphous condition.

As shown in FIG. 1, upon heating and compressing (e.g., under vacuum), the mixture of BMG powder and a crystalline Ti or Ti alloy powder to a temperature above the glass transition Tg of the BMG powder but below its Tx, a substantially fully dense composite is formed in which the titanium particles are embedded in a BMG alloy matrix. Upon heating, the matrix substantially above the crystallization temperature(s) Tx of the BMG matrix, the matrix crystallizes. In addition, at higher temperatures still below the full melt temperature Tm, the crystallized matrix and the reinforcing powders will react to form new crystalline phases and crystal structures which would not form from the BMG composition itself. In this way, stronger crystalline alloys can be fabricated in desired shape in a supercooled liquid matrix state, followed by conversion to new, higher performance alloys by heating to a temperature below the finished alloy melt temperature, which is higher than the BMG melting temperature Tm. Preferably, the final composite alloy composition will have a melting temperature Tm at least 100° C., and more preferably at least 250° C. higher than the melt temperature Tm of the original amorphous metal (BMG) component of the powder blend.

Even upon fully melting the newly formed composition, it will not form the same eutectic BMG alloy on cooling when the BMG and the particulate reinforcement are of different compositions. For example, a high aluminum BMG mixed with a titanium-group powder can form high-melting titanium-aluminide compositions.

It should also be noted that amorphous metal composites of a plurality of two or more different amorphous metal compositions may be provided which have advantageous properties, and which can provide an increased range of final alloy compositions when blended with the major titanium-group alloy component of the powder mixture.

As indicated, embodiments of the present disclosure is directed to processing of relatively small amounts of amorphous metal with relatively large amounts of crystalline (including quasicrystalline) metal and/or ceramic powders. Relatively small amounts of amorphous metals may be used in powder metallurgy processes and formulations to produce high performance parts and components. In conventional powder metallizing processes, metal powders and blends of metal powders with reinforcing or hardening components, are formed, compressed and sintered (in various orders and combinations of these steps by MIM, HIPing, compressive sintering, etc.), to produce finished or semi-finished products. One disadvantage of powder conventional metal sintering processes is that it is difficult to achieve full density by compression and sintering, even under vacuum hot isostatic pressing or expensive compression wave techniques. While amorphous metal powders with a sufficiently wide supercooled region can be compressed to full density and formed into complex shapes, they generally have low crystallization temperatures, and low melting points, which limit their use. In this regard, a relatively small amount of from about 1 to about 25 volume percent of an appropriate amorphous metal powder (or BMG coating) may be blended with from about 60 to about 99 volume percent of a conventional crystalline (including quasicrystalline) or non glass-transitioning amorphous titanium group metal powder, based on the total volume of the blend (excluding voids). From about 0 to about 15 volume percent of additional reinforcing powders, fibers, filaments, lubricating, molding and/or blending agents may also be incorporated in the blend. The amorphous metal powder should desirably exhibit a supercooled liquid temperature range above its Tg of at least about 20° C., and more preferably at least about 35° C. for at least 5 minutes before fully crystallizing to a point that it loses its gear-like viscosity and under shear. In addition, the amorphous metal powder and the conventional metal powder component(s) should form an alloy which both has a higher melting temperature than the liquid melting temperature of the amorphous metal itself, and which has a tensile strength which is larger than the tensile strength of the fully crystallized amorphous alloy composition. For example, 5 volume percent of a Zr-based, Al-based or Ti-based amorphous metal powder having a particle size smaller than 10 microns may be blended with titanium power, a titanium alloy powder, such as TiAl6V4, or a titanium intermetallic powder such as Ti3Al or TiAl. The powders are mixed to produce a homogeneous blend of ingredients. The powder mixture may be pressure or gravity-fed into a die, and compacted at a temperature in the superfluid temperature range of the Fe-based amorphous metal powder, between Tg and Tx. Compacting pressures may be increased, for example, from an initial 100 KPa to a final pressure in the 30-50 tons per square inch range in conventional PM presses. The compacted parts may subsequently be heated to at least partially crystallize the metallic glass composition or matrix. During sintering the pressed parts may be conveyed through a controlled-atmosphere furnace. The pressed powder particles may crystallize and/or react together, to form a dense steel product when using a strong iron or steel crystalline powder with a small amount of an amorphous BMG alloy such as a Ti, Zr and/or Fe-based BMG in this manner.

Conventional P/M presses operate with a vertical stroke to compact metal powders in a single cavity die to form the desired parts or components, such as gears such as pinion gears, bevel gears (straight and spiral), face gears and sprockets, cams, counterweights, armatures, pole pieces, bearings, and bushings. A wide variety of parts for automotive, computer, medical, dental, musical, electronic, tool, and aerospace uses may be produced in this manner.

For example, as indicated above, suitable metal powders include titanium and titanium alloy powders (in weight percent) such as:

    • Ti-3Al-8V-6Cr-4Zr-4Mo (Beta C)
    • Ti-15Mo-3Nb-3Al-0.2Si

High strength

    • Ti-6Al-4V
    • Ti-5Al-2.5Sn
    • Ti-4Al-4Mo-2Sn
    • Ti-6Al-6V-2Sn
    • Ti-10V-2Fe-3Al
    • Ti-15V-3Cr-3Sn-3Al
    • Ti-5.5Al-3Sn-3Zr-0.5Nb
    • Ti-5Al-2Sn-4Mo-2Zr-4Cr
    • Ti-8Al-1Mo-1V

High Temperature

    • Alpha—Two Aluminide (Ti-24Al-11Nb, intermetallic compound based on Ti3Al)
    • Alpha—Two Aluminide (25/10/3/1)
    • Ti-6Al-2Sn-4Zr-2Mo
    • Ti-6Al-7Nb
    • Ti-6Al-2Sn-4Zr-6Mo
    • Ti-5.5Al-3.5Sn-3Zr-1Nb
    • Ti-5.8Al-4Sn-3.5Zr-0.7Nb
    • Ti-6Al-2Sn-4Zr-6Mo

Powder formation for both the titanium alloy compositions, and the powdered amorphous metal component (which is desirably cooled) at a rate, for example, of at least 1×104 degrees K. per second for some less-stable alloys, can be accomplished by a number of different processes, including, for example, melt-spinning, gas atomization, centrifugal atomization, and splat quenching. The powder can be consolidated by, for example, hipping, hot pressing, hot extrusion, powder rolling, powder forging and dynamic compaction. In centrifugal atomization, for example, the melt stream is discharged from a rapidly spinning centrifugal cup, and is contacted by high pressure cold helium gas to facilitate fast cooling (e.g., greater than 1×105 K/s). The helium gas can be collected, purified and reused. The speed of the rotating centrifugal cup may, for example, be about 20,000-45,000 (e.g., 40,000) RPM, the speed and other processing conditions can be adjusted to produce a fine powder with about a 25 micrometer mean particle size.

As indicated, it may be desirable to provide conventional (crystalline/nanocrystalline) metal surfaces, fibers, wires, sheets and powders with a thin BMG surface which facilitates metallurgical joining of these materials with the BMG. The surface layers may be provided in a variety of ways, including PVD. One way to deposit an amorphous metal layer on a substrate is to fully vaporize an amorphous metal composition, such as by E-beam or plasma jet, and subsequently rapidly condense and cool it on the powder, fiber or sheet surface of the crystalline metal or ceramic to be coated, and cool the condensed deposit within the amorphous metal cooling time required to prevent crystallization. For example, when utilizing a plasma gun, as a suitable BMG powder with a particle size of less than 5 microns, and preferably less than 3 microns, is fully vaporized in a plasma jet in an argon atmosphere. The plasma jet will typically reach a temperature of at least 8,000-15,000 K, and small particle size of the BMG facilitates its vaporization. The vaporized metal and the plasma jet stream may be expanded to cool it and substantially simultaneously mix it with a cooled (e.g, room temperature) crystallized titanium group metal powder, at a mass ratio of 1-100 times the mass of the vaporized metal. Vaporized BMG rapidly condenses on the crystalline powder, and can cool rapidly to form an amorphous, preferably metallurgically-bonded thin amorphous metal surface layer. An amorphous metal coating may also be electroplated or electrolessly deposited on the titanium group powder, or fibers.

The use of filament reinforcements is also an important option. The amorphous metal powder blend may also include other components such as reinforcing and/or alloying fibers or powders. Such fibers or powders may be densely consolidated within the metal glass matrix. The additional particulate component may also provide alloying constituents to modify the composition of the alloy(s) crystallized from the amorphous alloy upon heating above the crystallization temperature(s) Tx.

As illustrated in FIG. 1 a blend of a BMG amorphous alloy powder with a crystalline metal which, may be, for example, a Ti, Zr or Al-based powder or an Fe-based powder having a mean particle size from 0.5 to 25 microns. For Al, Zr and/or Ti-based BMG amorphous metal matrices, the crystalline metal powder may be, for example, a titanium powder, a titanium aluminide powder such as Ti3Al or TiAl, or a crystallized titanium group alloy powder, having a particle size of from about 0.5 to about 50 microns. Bulk Metallic Glasses having a broad supercooled liquid range (e.g., preferably 50 K or more) can be shaped and fabricated from their viscous, low-temperature glassy melts. When heated above Tg, they can form a “soft” viscous glass without crystallizing, typically at relatively low temperatures of 200-400° C. Bulk metallic glasses can be shaped by die forging and extrusion of amorphous metal powders [e.g., Kawamura, Y. et al., “Development and industrialization of powder-consolidation and forging technologies in metallic glasses” Funtai Oyobi Fummatsu Yakin/Journal of the Japan Society of Powder and Powder Metallurgy, v 48, p 845-853 (2001)].

Net-shape, low-temperature, low pressure powder metallurgy fabrication processes for rapidly producing full density high-performance Titanium alloys and lighter-weight Titanium alloy metal matrix composites with lightweight, high-stiffness filaments or fibers such as SiC or SiBCN are a significant advance for rapid prototyping and rapid manufacture of a very wide variety of high-performance Titanium-based components, which conventionally require much more expensive tooling, forging and shaping. As shown in FIG. 2, virtually any Titanium product shape can be produced using relatively inexpensive, low-pressure, low-temperature molds or dies, which themselves can be made quickly and inexpensively because of the relatively low temperature and strength requirements involved.

The disclosed process technology permits inexpensive fabrication of advanced Titanium-based materials, rapid prototyping of new components and structures, and rapid manufacturing of high-performance Titanium and Titanium composite components for new and current components and subcomponents.

The powder blend can be compacted at relatively low temperature, which is above the Tg of the Bulk Metal Glass (e.g., 300° C.) but well below the sintering temperature of titanium powders (e.g., 1200° C.). Compacting pressures are also relatively low, because it is primarily the viscous Bulk Metal Glass which is consolidated, not the hard titanium powders. The compacted parts may be used with a BMG matrix, or can subsequently be heated to at least partially crystallize the metallic glass composition or matrix. The particles may crystallize and/or react together, to form a dense product. The powder blend is compressed at a relatively low temperature (well below the melting point) within the supercooled liquid range of the amorphous alloy, to eliminate the porosity of the blend, and increase the total surface area of the amorphous metal powder. The blend may also contain reinforcing filaments or fibers such as strong, lightweight B, SiC, or SiCN fibers, which will be incorporated in the finished products. Desirably, the titanium group powders and amorphous metal powders should have a particle size of less than about 200 microns, preferably with a range of particle sizes to facilitate packing. The mean particle size of the BMG amorphous metal component should probably be less than half that of the titanium group metal particles, to facilitate more uniform dispersion in the interstices of the titanium group particles.

As shown in FIG. 2, the blend is subsequently compacted above the Tg of the BMG component, where it is a viscous liquid. This compression produces viscous flow of the amorphous metal glass, forcing it into the interstices of the titanium group metal particles, substantially increasing the total surface area of the amorphous metal component and desirably forming it into an at least partially interconnected matrix enclosing the titanium group metal particles. This decreases the median thickness of the amorphous metal component. This facilitates reactive diffusion with the titanium group metal component at elevated temperatures to form new high performance alloy(s) having much higher melt temperatures than the amorphous metal component.

As indicated, the powder mixture comprises a major volume fraction V of titanium group metal powder and any reinforcing filaments, and a minor volume fraction v of the BMG amorphous metal powder or coating component. The compressive pressure Pc applied to compact the mixed powder is much less than that necessary to compact the titanium group metal powder. This compression pressure Pc may be approximated by:
Pc=−2Y{ln(1−R)/3}
where Pc is the minimum compression pressure at the supercooled liquid temperature, Y is the yield strength of the titanium group metal powder at the compression temperature, and R is the relative density of the titanium group metal powder (in the absence of the BMG amorphous alloy component).

For example, a titanium group metal powder having a yield strength of 300 MPa at 700° K. would itself (without the presence of a BMG amorphous metal component) be compressed to a relative density R of 0.85 (85% of full density), by a compression pressure Pc of approximately −2×300 (ln(1−0.85)/3 or approximately Pc=−380 MPa. To compress the titanium group metal powder itself (without a BMG amorphous metal component) to a relative full density of 0.999, could require over 1.3 Gpa, which is difficult and expensive to provide in an economical production process under vacuum or inert atmosphere.

The compacted PM product may be cooled to provide a dense metal compact in which the crystalline titanium group metal particles are embedded in an ultrastrong amorphous metal matrix.

Importantly, the fully compacted titanium alloy product may also be heated to a temperature above the crystallization temperature of the amorphous metal continuous phase. In this manner, the titanium metal group particles are embedded in a fully or partially crystalline continuous phase which is still substantially distinct from the titanium group metal particles. This can be advantageous for certain BMG alloys which form a strong, partially or fully crystalline alloy upon heating to an initial crystallization temperature Tx or above.

Such titanium alloy products with discrete titanium group powder zones (even though compressed or distorted) and BMG alloy composition zones, will not have optimum heat resistance under stress for high temperature operation, because the partially discrete continuous amorphous phase may remain as a eutectic material with a relatively low melting point, Tm. In order to provide a finished product with increased heat resistance, the compact product may be further heated to a sintering temperature at which the titanium group alloy particles and the relatively continuous matrix phase diffuse together to form a new composition of high performance titanium. The Bulk Metal Glass, and the titanium alloy components may be selected to provide high-performance titanium alloys upon diffusion-reaction. This reaction may form selected Ti-alloys, with formation of nano- or micron-scale ceramic reinforcing particles such as TiB2, TiNi, etc., which improve the high-temperature strength of the alloy.

After diffusion of the original crystalline metal powder components with the original amorphous metal components, the composite does not “remelt” at the relatively low melting temperature of the bulk metal glass component, because of the different composition of the reacted compact. The new, high performance titanium alloy may include highly refractory ceramic (e.g., TiB2) and intermetallic (e.g., TiCu, NiAl3, TiAl) reinforcements produced in situ, as previously indicated.

EXAMPLE 1

Al85Ni5Y8Co2 fully amorphous alloy has a glass transition temperature Tg of 538° K., a relatively high strength of 1250 MPa at room temperature and a relatively wide and a 38° K. supercooled liquid region [Y. Kawamura, et al, “Nanocrystalline Aluminum Bulk Alloys With A High Strength Of 1420 Mpa Produced By The Consolidation Of Amorphous Powders”, Scripta mater. 44 (2001) 1599-1604; A. Inoue, et al., Mater. Trans. JIM. 31, 493 (1990).]. Its initial crystallization occurs through the precipitation of fcc-Al particles while retaining viscous flow in the supercooled liquid region [Y. Kawamura, et al, Int. J. Powder Metall. 33, 50 (1997); Y. Kawamura, et al, J. Jpn. Soc. Powder Powder Metall. 38, 948 (1991); Y. Kawamura, et al, Mater. Trans. JIM. 40, 749 (1999)], and accordingly is used as a minor BMG component with a large volume of crystalline titanium powder.

A one Kg ingot of an Al85Ni5Y8Co2 (at %) alloy is prepared by arc melting a mixture of the pure elements. The ingot is re-melted at a temperature of 1573 K and then gas atomized through a melt stream with a diameter of 2.0 mm, using helium gas with a dynamic pressure of 10 MPa for atomization. The powder is sieved to below 25 microns. 200 grams of this amorphous alloy powder is uniformly mixed with 800 grams of CP titanium powder having a particle size of less than about 75 microns. The CP titanium has a tensile strength of about 220 MPa at ambient temperature. The entire blend (100 g) is flushed with argon, and heated in a steel die under vacuum to a temperature of 560° K., under a compaction pressure of 240 MPa, which upon cooling to 25° C. produces a fully dense composite cylinder of CP titanium powder embedded in a somewhat discontinuous matrix of the amorphous metal composition.

The composite cylinder is subsequently heated to a temperature of 1100° C., in an inert atmosphere. As the temperature increases over Tx of the amorphous material, it crystallizes. Upon additional heating, the CP titanium powder and the amorphous matrix reactively diffuses to “consume” the amorphous matrix to form a high-aluminum, unitary titanium alloy having a melting temperature much higher than the amorphous matrix. Similar runs using Fe, Zr, Al and Cu-based amorphous alloys as described above (compressed in their supercooled liquid region) produce similar results, with different final alloy compositions.

EXAMPLE 2

200 grams of SiBNC fibers having a tensile strength of 2-4 GPa, a density of 1.8/cm3, a diameter of 8-14 microns [See e.g., H. P. Balddus, et al., “Properties of Amorphous SiBNC-Ceramic Fibers”, Key Engineering Materials, Vol. 127-131, pp. 177-184 (1977)] are aligned in a blend of 900 grams of CP titanium powder and 100 grams of a BMG alloy powder, by weight, based on the total powder weight. The BMG alloy powder is Ti50Cu20Ni24Si4B2 (atomic percent), having a Tg of about 745° K. and a supercooled liquid region of about 65° K. The aligned fibers and powder blend are compressed between sheet platens under vacuum at a temperature of 8000 K, at a compression pressure of 465 MPa to produce a fully dense sheet “tape” with a monolayer of aligned fibers. The sheet is removed from the platen press and subsequently heated in an inert temperature to a temperature of about 1200° C. to form a fiber-reinforced, high-temperature composite, in which the fibers are embedded in a matrix having an overall alloy composition of Ti94.18, Cu2.63, Ni2.92, Si0.23 and B0.04 percent by weight.

EXAMPLE 3

89 grams of a titanium alloy powder which is 96.4 atomic percent titanium and 3.6 atomic percent vanadium is blended with 11 grams of a bulk metal glass powder having a composition of A194V4Fe2 atomic percent. The 100 grams of the powder blend is compressed in a gear-shaped mold under vacuum at a compression pressure of 25 kSI, at a temperature 25° above the glass transition temperature Tg of the amorphous alloy to form a substantially fully dense gear component. The gear is removed from the mold, and heated in a vacuum furnace to a temperature of 1000° C. to first crystallize the BMG and then reactively diffuse it with the Ti alloy powder, to form an alloy with a nominal composition in weight percent of 89.6Ti, 6.09Al, 4.04V and 0.27Fe.

EXAMPLE 4

Silicon carbide fibers (SCS-6 fibers of Textron Specialty Materials Company) are sputter coated with a layer of Nb about one micron in thickness. Alternatively, an amorphous metal alloy composition having predominantly beta-stabilizing elements, such as Zr70Fe20Ni10 may be sputter-coated on the filaments, while coating the filaments below 500° K. to maintain the coating in an amorphous condition.

900 grams of titanium alloy powder (particle size 50-100 microns) having a composition of 95 weight percent Ti, a5 weight percent V, 2 weight percent Mo, is blended with 10 weight percent of an amorphous Zr alloy powder (particle size 20-50 microns) having a composition of Zr66Al18Cu7Ni19 (atomic %). 200 grams of the coated SiC fibers are aligned and embedded in the powder blend, and compressed under vacuum at a temperature of 715° K. and a pressure of 45 ksi to form a substantially dense foil with SiC monofilaments embedded therein. The reinforced foil is subsequently cut and stacked with other like foils to form finished component shaped, and compressed and heated to 1100° C. in an inert atmosphere to reaction-diffuse the amorphous alloy with the titanium alloy powder, and to unify the sheets to form a high-strength, fiber reinforced composite article suitable for high temperature use.

Referenced by
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Classifications
U.S. Classification419/66
International ClassificationB22F1/00
Cooperative ClassificationB22F2999/00, B22F1/0003
European ClassificationB22F1/00A