Abstract:
Aluminum alloys having improved strength at 300° C. characterized by formation from an intermediate amorphous state to a final fcc matrix hardened by optimal 25 nm-diameter Ll 2  precipitates with an interphase misfit less than about 4% in all three dimensions and Al 23 Ni 6 M 4  precipitates where M is one or more elements selected from the group consisting of Y and Yb. An appropriate melt of aluminum with selected transition metals (Co, Cu, Fe, Ni, Ti, Y) and Ll 2  stabilizers (Sc, Yb) in amounts of about 2 to 12 and 2 to 15 atomic percent, respectively, is processed to achieve an intermediate amorphous state to dissolve Ll 2 -forming components. The amorphous alloys are then thermo-mechanically devitrified to a final crystalline microstructure. The alloys have good ductility and a short-term tensile strength exceeding about 275 MPa (40 ksi) at 300° C., and are useful for applications such as high-temperature turbine engine components or aircraft structural components.

Description:
CROSS REFERENCE TO RELATED APPLICATIONS 
       [0001]    This application is a continuation-in-part utility application of application Ser. No. 10/422,234 filed Apr. 24, 2003 entitled Nanophase Precipitation-Strengthened Al Alloys Processed Through the Amorphous State, which is based upon previously filed provisional applications: Ser. No. 60/375,940 filed Apr. 24, 2002 entitled “Amorphous metal alloy compositions” and Ser. No. 60/450,114 filed Feb. 25, 2003 entitled “Amorphous metal alloy compositions”, all of which are incorporated by reference and for which priority is claimed. 
     
    
     GOVERNMENT INTERESTS 
       [0002]    Activities relating to the development of the subject matter of this invention were funded at least in part by United States Government, U.S. Army Aviation &amp; Missile Command Contract No. DAAH01-02-C-R125, and thus may be subject to license rights and other rights in the United States. 
     
    
     BACKGROUND OF THE INVENTION 
       [0003]    In a principal aspect, the present invention relates to Al-based alloys processed through an amorphous state, preferably by means of a Rapid Solidification Process (RSP) from molten alloy, and then devitrified to a primarily crystalline microscale fine grain structure by thermo-mechanical processing. To promote glass-forming ability, the Al alloys comprise selected transition metal (TM) and lanthanide rare earth (RE) elements. The final crystalline microstructure has a combination of stable strength at or above about 300° C. and good ductility, characterized by optimal  25  nm-diameter Ll 2  precipitates in an fcc matrix with an interphase misfit typically less than about 4% in all three dimensions, and rod-shaped Al 23 Ni 6 M 4  precipitates. 
         [0004]    Improved strength at elevated temperatures has been a continuing goal in Al alloy development for more than three decades. Currently available commercial Al alloys, either manufactured with ingot or powder processing, are not capable of simultaneously achieving high strength and high-temperature stability near 300° C.; such characteristics being particularly important in applications such as fan components in turbine engines. Precipitation hardening introduced by aging is a known method to strengthen Al alloys. Conventional high-strength Al alloys in commercial applications employ Guinier-Preston zones and subsequent precipitation at or below 250° C. Examples of Al alloys processed with relatively high aging temperatures in commercial practice include alloy 2618 (200° C. for 20 hours), 4032 (170-175° C. for 8-12 hours), and 2218 (240° C. for 6 hours) [Metals Handbook-Properties and Selection: Nonferrous Alloys and Special-Purpose Materials, Volume 2, 10 th  Edition, ASM International]. At the noted aging temperatures, these alloys have an improved microstructure stability relative to other commercial Al alloys. These Al alloys, when precipitation-hardened, usually possess a room temperature yield strength of about 600 MPa. (85 ksi). However, at temperatures approaching 300° C., the precipitation hardening efficiency in these alloys quickly and significantly degrades as a result of precipitate coarsening and/or dissolution. Due to the unstable strengthening precipitate size distribution at such high temperatures, the yield strength of currently available aluminum alloys at 300° C. is often only 10% of the yield strength at room temperature, and thus renders such alloys unsuitable for high-temperature applications above 150° C. For high-temperature turbine engine components or aircraft structural components, a short-term tensile strength exceeding about 275 MPa (40 ksi) at 300° C. is desired. 
         [0005]    In order to achieve a combination of high strength and usable high-temperature properties in Al alloys, researchers have investigated a variety of intermetallic precipitation dispersions. The Al-based Ll 2  phase is one of the best-known precipitates to achieve a good combination of high strength and high toughness of ambient temperatures. There are reportedly only seven elements stabilizing the Ll 2  phase: Er, Lu, Np, Sc, Tm, U, and Yb [Knipling, K. E. et al. Z.  Metallkd  97:246-265]. Since crystalline Al has very limited solubility for these Ll 2  stabilizers, it is difficult to produce a fine dispersion through crystalline solid-state heat treatments. Alternatively, with RSP from the liquid state, it is possible to either (1) directly produce a fine crystalline structure, or (2) produce partially amorphous Al alloys. Nonetheless, crystalline RSP Al alloys have not been able to meet the high-temperature strength requirements due to the difficulty of producing small, stable particles at adequate volume fraction. The focus on amorphous RSP Al alloys has been primarily on face centered cubic (fcc)-Al nanocrystals to enhance the ambient strength [Kim, Inoue, Masumoto, Mater Trans JIM 1990; 31: 747]. Upon devitrification, Al nanocrystals of up to 30% volume fraction can be dispersed within the amorphous matrix. However, this nanoscale ultra-fine grain stricture is undesirable because at high temperatures, typically ≧0.4-0.5 T m  where T m  is the material&#39;s absolute melting temperature, the contribution of grain boundary strengthening is minimal and the refined grain structures promote rapid diffusional creep. In addition, it has been reported that ultra-fine grain sizes may be undesirable when, considering formability and fracture toughness [Hornbogen, Starke, Acta metall. mater. 1993; 41: 1]. 
         [0006]    In sum, previous development on Al-based materials with high strength at elevated temperatures failed to meet the property objective of 275 MPa at 300° C. 
       SUMMARY OF THE INVENTION 
       [0007]    The present invention is directed to a new class of Al alloys characterized by formation from an intermediate amorphous state to a final fcc-Al matrix hardened by a combination of Ll 2  precipitates and Al 23 Ni 6 M 4  precipitates in order to establish the Al-based analogue of Ni-base superalloys and achieve high-temperature strength with usable ductility. 
         [0008]    An appropriate melt of Al with selected TM and RE is first processed to achieve an intermediate amorphous state to dissolve Ll 2 -forming components. The preferred method to achieve a primarily (above 70% in volume) amorphous state is RSP from the molten alloy by process techniques such as powder atomization, melt spinning, and spray casting. The RSP process should have a cooling rate of at least about 10 3 ° C./sec, preferably at least 10 4 ° C./sec. Other methods to achieve amorphous microstructure through a solid-state process, such as mechanical milling, may also be used. The intermediate amorphous alloys are then thermo-mechanically devitrified to a final primarily (above 70% in volume) fcc/Ll 2 /Al 23 Ni 6 M 4  crystalline microstructure with at least about 70% fcc phase in volume, at least about 0.10% Ll 2  phase, and at least about 10% Al 23 Ni 6 M 4  phase in volume. 
         [0009]    The selection of alloying elements is based on (1) good glass-forming ability with RSP, (2) long-term strength at or above 300° C., and (3) composition tolerance for a robust design. For glass-forming ability, elements with strong short-range ordering effects, and slow long-range diffusing kinetics in molten Al are employed. For long-term strength at or above 300° C., the alloy of the present invention employs 25 nm-diameter Ll 2  particles which are reported to provide optimal creep resistance [E. A. Marquis, Microstructural Evolution and Strengthening Mechanisms in Al—Sc and Al—Mg—Sc Alloys, Ph.D. thesis, Northwestern University, 2002.]. For a robust design, the present invention employs Al 23 Ni 6 M 4 , where M is one or more elements selected from the group consisting of Y and Yb. When there is deficiency of the Ll 2 -formers, the incoherent D0 11 -Al 3 Ni phase is expected to precipitate, leading to low ductility. Al 23 Ni 6 M 4  is more solute-rich that the Al 3 X phase and will consume less Al for a given amount of solute, giving rise to a higher amount of fcc matrix which in turn increases the ductility. 
         [0010]    Thus, it is an object of the invention to provide a new class of high-temperature high-strength Ll 2 -phase strengthened Al alloys processed through the amorphous state, preferably with RSP, and then subsequently devitrified with thermo-mechanical, processes. 
         [0011]    A further object of the invention is to combine selected TM and RE to provide good glass forming ability during RSP such as powder atomization or melt spinning to form an amorphous Al alloy, dissolving the Ll 2 -stabilizers before the devitrification process. 
         [0012]    Another object of the invention is to provide aluminum alloys with usable strength at or above about 300° C. by selecting Ll 2 -stabilizers which reduce the interphase lattice misfit in all three dimensions to promote a finer dispersion. 
         [0013]    Another object of the invention is to employ Al 23 Ni 6 M 4  precipitates to provide composition tolerance and maintain reasonable alloy ductility. 
         [0014]    These and a other objects, advantages and features will be set forth in the detailed description which follows. 
     
    
     
       BRIEF DESCRIPTION OF THE DRAWING 
         [0015]    In the detailed description which follows, reference will be made to the drawing comprised of the following figures: 
           [0016]      FIG. 1  is an X-ray diffractogram of the alloy of Example 1 as melt-spun with positions of fcc pure aluminum reflections indicating a fully amorphous state; 
           [0017]      FIG. 2  is an X-ray diffractogram of the alloy of Example 1 after devitrification at 550° C. for 24 hours, with positions of reflections of pure fcc Al, Al 3 Yb, and Al 23 Ni 6 Yb 4  phases, indicating the desired phases: fcc+Ll 2 ; 
           [0018]      FIG. 3  is a Scanning Electron Microscope (SEM) secondary electron image of devitrified alloy of Example 1 indicating phase constituents fcc+Ll 2 Al 23 Ni 6 Yb 4 ; and 
           [0019]      FIG. 4  is an SEM secondary electron image of devitrified alloy of Example 2 indicating phase constituents fcc+Ll 2 +Al 23 Ni 6 Yb 4 . 
       
    
    
     DESCRIPTION OF THE PREFERRED EMBODIMENTS 
       [0020]    General Summary 
         [0021]    In general, the subject matter of the invention comprises an Al alloy in crystalline form having higher or greater strength particularly at elevated temperatures, i.e. greater than about 300° C. The Al alloy is; made by compounding a mixture of Al with selected TM and RE in amorphous state followed by devitrification to a mixed crystalline state comprising fcc, Ll 2 , and Al 23 Ni 6 M 4  phases wherein the ratios of the crystalline states are within certain preferred ranges. Preferably the resultant alloy has at least about 70% by volume fcc phase, at least about 10% by volume Ll 2  phase, and at least about 10% by volume Al 23 Ni 6 M 4  phase where M is selected from the group consisting of Y, Yb and a combination of Y and Yb with limited residual amorphous or quasi-crystalline phase material. 
         [0022]    The choice of starting materials may vary, as may the compounding processes, the glass formation processes and the devitrification processes. In the amorphous state, there may be some crystalline material contained therein, but preferably no more than about 30% by volume. The particle size of alloys passing through a fully or almost fully glassy state is much finer than that of alloys without passing through the glassy state or only passing through a partially glassy state with Ll 2  already present in the as-spun condition. Thus forming the mixture in the amorphous intermediate state constitutes a very important aspect of the invention. 
         [0023]    The alloy materials, in addition to Al, include one or more TM taken or selected from the group of Cu, Ni, Co, Ti, Fe, Y, and Sc, and one or more RE selected or taken from the group of Er, Tm, Yb, and Lu. TM metals are utilized in the range of about 2 to 12 at %, and RE materials are utilized in the range of about 2 to 15 at %. 
         [0024]    The processes for mixing or forming the starting materials in the amorphous state are not necessarily limiting. Thus, it is contemplated that solid state processing, liquid or melt processing as well as gas phase processing may be utilized, though liquid phase processing is preferred. The completeness of the amorphous state is at least about 70% by volume and preferably greater. 
         [0025]    Development Technique 
         [0026]    Precipitation-hardened Al alloys are difficult to develop for high strength due to limited solubility of alloying elements. Al alloys with high fractions of precipitate that cannot be completely solution-treated have very coarse particles that tend to limit strength, corrosion resistance and toughness. In contrast, the Al alloys of the present invention exhibit high strength, good ductility, and high-temperature stability at or above 300° C. 
         [0027]    By carefully selecting an appropriate Al alloy composition, processing techniques can achieve a fully amorphous state after rapid cooling. Furthermore, this glass can then be thermo-mechanically processed such that the glass devitrifies into a crystalline fcc matrix with nanophase precipitates. By passing through the glass state, the equilibrium solidification that would produce coarse precipitates is avoided. Certain TM such as Fe, Co, Ni and Cu promote short-range ordering in liquid Al, which leads to low partial molar volume, low thermal expansion, and high viscosity that are beneficial to glass-forming ability. RE elements such as Ce, Gd, Yb, and Er with large atomic size exhibit low diffusivity in Al and thus retard crystal nucleation. Therefore, Al-TM-RE comprise a class of glass-forming system for Al alloys of the present invention. 
         [0028]    The elements Er, Lu, Tm and Yb are reported as the only RE Ll 2 -stabilizers. Among these four RE elements, Yb has the smallest lattice parameter and relatively low-cost. Er has the lowest cost. To evaluate the effect on glass-forming ability of these alloying additions, a reduced glass transition temperature (T rg ) model was developed. In the Al-TM-RE system, this model predicts that Er has no beneficial effect to T rg . As a consequence, alloys of the invention utilize Yb as the preferred Ll 2 -stabilizer rather than Er, Tm, and Lu. 
         [0029]    Sc is the oily TM element that can form a stable Ll 2  with Al. Compared to RE Ll 2  formers, Sc can form Ll 2  with a smaller lattice parameter, reducing the misfit between Ll 2  and Al matrix. However, Sc is by far the most expensive of the Ll 2 -stabilizers and therefore embodiments of the invention seek to limit. Sc as much as possible. Efforts have been made to search for other TM to substitute for Sc. A preliminary requirement for such substitution is solubility. Ti has a substantial solubility in Al 3 Sc. In addition, Ti has the lowest diffusion coefficient in solid Al among TMs. Adding Ti to Al 3 Sc thus reduces the coarsening rate of Ll 2  precipitates. Moreover, addition of Ti decreases the lattice parameter of Al 3 (Sc,Ti) and hence minimizes the lattice misfit with Al. Thus, alloys of the invention incorporate Yb and Sc as base Ll 2  formers but are not limited to these elements. TM such as Ti, V, Zr, etc., which will result in low misfit and thus retard coarsening are considered useful. 
         [0030]    For a robust design, the present invention employs Al 23 Ni 6 M 4 , where M is one or more elements selected from the group consisting of Y and Yb. To introduce both Al 23 Ni 6 M 4  and Ll 2  in the design, thermodynamic equilibrium calculations were performed using the thermodynamic database and calculation package Thermo-Calc® [Sundman, B. B. Jansson, and J. O. Andersson. 1985 . Calphad  9: 153-190]. Thermodynamic calculations predict that Y has certain solubility in Ll 2 , which expands the Ll 2  lattice spacing, increasing the misfit. Therefore, a design criterion should be set to limit the partitioning of Y in Ll 2 . In addition, other phases such as Al 3 Ni, Al 3 Y and Al 9 CO 2  should be avoided. 
         [0031]    Al-base alloys will have good ductility when the amount of fcc is equal to or over about 70%. Thus, the total amount of Al 23 Ni 6 M 4  and Ll 2  is fixed to less than about 30%. At the desired phase constitution, Co content is set by [x Ni +x Co ]/x Y =6/4 because Co has a small solubility in Al 23 Ni 6 M 4  by substituting for Ni. After examining the effect of Co addition based on thermodynamic calculations, an optimum was found around 0.6 at % Co, at which partitioning of Y in Ll 2  is almost zero. If Co addition is significantly more than 0.6 at %, Al 9 CO 2  and Al 3 Y may precipitate. 
         [0032]    Experimental Results 
         [0033]    The present invention alloys, through, computational design of multi-component Al-TM-RE systems incorporate, desired processing properties-glass forming ability and the desired microstructure—a fine dispersion of Ll 2  after devitrification in the Al matrix. 
       Example 1 
       [0034]    Prototypes of preferred embodiments can be made by arc-melting, melt spinning or wedge casting. Through melt spinning, ribbons of Al-3.46Ni-2.78Y-0.72Co-0.42Yb-0.63Sc-0.42Zr-0.21Ti (at %) were made. Melt-spun ribbons are approximately 3-4 mm wide and 30-40μ in thickness. The ribbons were characterized using micro-hardness) x-ray diffraction, and SEM analysis. The x-ray diffraction pattern ( FIG. 1 ) of the as-spun ribbon indicates a partial amorphous microstructure without intermetallic precipitates. After devitrification at 550° C. for 24 hours, x-ray diffraction ( FIG. 2 ) shows precipitation of Al 23 Ni 6 Yb 4  and peaks of Ll 2 . It is noted that the peaks of Ll 2  are shifted compared to Ll 2 -Al 3 Yb, indicating; decrease of lattice parameters due to dissolution of Sc, Ti, and Zr in Al 3 Yb. Such decrease of the Ll 2  lattice parameter will reduce the misfit.  FIG. 3  shows an SEM image of the devitrified specimens confirming the phase constituents fcc+Ll 2 +Al 23 Ni 6 Yb 4 . The matrix is fcc-Al, the large sized grey phase material is Al 23 Ni 6 Yb 4 , and the small white particles are Ll 2  phase particles. The Ll 2  particles remain smaller than ˜50 nm in diameter, while the rod-shaped Al 23 Ni 6 Yb 4  phase material is less than 1μ in length. The small Ll 2  particles will provide optimal creep resistance at or above 300° C. and the Al 23 Ni 6 Yb 4  material is present to avoid detrimental compounds and improve the ductility at the high temperature. 
       Example 2 
       [0035]    Ribbons of Al-3Ni-2.42Y-0.62Co-0.6Yb-0.6Sc-0.6Zr-0.6Ti (at %) were made using the protocol of Example 1. The ribbons were characterized using micro-hardness, x-ray diffraction, and SEM analysis.  FIG. 4  shows an SEM image of the devitrified specimens confirming the phase constituents fcc+Ll 2 +Al 23 Ni 6 Yb 4 . The matrix is fcc-Al, the large sized grey phase material is Al 23 Ni 6 Yb 4 , and the small white particles are Ll 2  phase material. The Ll 2  particles remain smaller than ˜50 nm in diameter, while the rod-shaped Al 23 Ni 6 Yb 4  phase material is less than 1μ in length. The small Ll 2  particles will provide optimal creep resistance at or above 300° C. and the Al 23 Ni 6 Yb 4  material is present to avoid detrimental compounds and improve the ductility. 
       Example 3 
       [0036]    Scale-up processing of the alloy in Example 1 was engineered. Amorphous powder produced by high-pressure He atomization can be used as a raw material to produce an amorphous bulk by consolidation at high temperatures. The amorphous alloy powder is produced by gas atomization, followed by sieving, precompaction, canning and sealing into a Cu tube, carried out in a well-controlled atmosphere with an oxide or moisture concentration below 1 ppm. Powder of the alloy in Example 1 was successfully atomized and extruded. The extrusion is a thermo-mechanical process where the glass devitrifies into a crystalline fcc matrix with nanophase precipitates. 
       Example 4 
       [0037]    Powder of the alloy in Example 2 was successfully atomized and extruded using the protocol of Example 3. The extrusion is a thermo-mechanical process where the glass devitrifies into a crystalline fcc matrix with nanophase precipitates. 
         [0038]    Variations of the described aluminum alloy as well as the process for manufacture thereof and the product created by the process arc available to provide the expected functionality of high short-term and long-term strength at temperatures above about 300° C. Thus the invention is to be limited only by the following claims and equivalents thereof.